light metal systems. part 3: selected systems from al-fe-v to al-ni-zr

XI

Landolt-BörnsteinNew Series IV/11A2

MSIT®

Introduction

Introduction

Data Covered

The series focuses on light metal ternary systems and includes phase equilibria of importance for alloydevelopment, processing or application, reporting on selected ternary systems of importance to industriallight alloy development and systems which gained otherwise scientific interest in the recent years.

General

The series provides consistent phase diagram descriptions for individual ternary systems. Therepresentation of the equilibria of ternary systems as a function of temperature results in spacial diagramswhose sections and projections are generally published in the literature. Phase equilibria are described interms of liquidus, solidus and solvus projections, isothermal and pseudobinary sections; data on invariantequilibria are generally given in the form of tables.

The world literature is thoroughly and systematically searched back to the year 1900. Then, thepublished data are critically evaluated by experts in materials science and reviewed. Conflicting informationis commented upon and errors and inconsistencies removed wherever possible. It considers those, and onlythose data, which are firmly established, comments on questionable findings and justifies re-interpretationsmade by the authors of the evaluation reports.

In general, the approach used to discuss the phase relationships is to consider changes in state and phasereactions which occur with decreasing temperature. This has influenced the terminology employed and isreflected in the tables and the reaction schemes presented.

The system reports present concise descriptions and hence do not repeat in the text facts which canclearly be read from the diagrams. For most purposes the use of the compendium is expected to be self-sufficient. However, a detailed bibliography of all cited references is given to enable original sources ofinformation to be studied if required.

Structure of a System Report

The constitutional description of an alloy system consists of text and a table/diagram section which areseparated by the bibliography referring to the original literature (see Fig. 1). The tables and diagrams carrythe essential constitutional information and are commented on in the text if necessary.

Where published data allow, the following sections are provided in each report:

Literature Data

The opening text reviews briefly the status of knowledge published on the system and outlines theexperimental methods that have been applied. Furthermore, attention may be drawn to questions which arestill open or to cases where conclusions from the evaluation work modified the published phase diagram.

Binary Systems

Where binary systems are accepted from standard compilations reference is made to these compilations.In other cases the accepted binary phase diagrams are reproduced for the convenience of the reader. Theselection of the binary systems used as a basis for the evaluation of the ternary system was at the discretionof the assessor.

XII


MSIT®

Introduction

Solid Phases

The tabular listing of solid phases incorporates knowledge of the phases which is necessary or helpfulfor understanding the text and diagrams. Throughout a system report a unique phase name and abbreviationis allocated to each phase.

Phases with the same formulae but different space lattices (e.g. allotropic transformation) aredistinguished by:

– small letters (h), high temperature modification (h2 > h1)(r), room temperature modification(1), low temperature modification (l1 > l2)

– Greek letters, e.g., , '– Roman numerals, e.g., (I) and (II) for different pressure modifications.In the table “Solid Phases” ternary phases are denoted by * and different phases are separated by

horizontal lines.

Heading

Literature Data

Binary Systems

Solid Phases

Pseudobinary Systems

Invariant Equilibria

Liquidus, Solidus, Solvus Surfaces

Isothermal Sections

Miscellaneous

Miscellaneous

Isothermal Sections




Solid Phases

Binary Systems

Text

References

Tables anddiagrams

Temperature-Composition Sections

Temperature-Composition Sections

Thermodynamics

Notes on Materials Properties and Applications

Thermodynamics


Fig. 1: Structure of a system report

XIII


MSIT®

Introduction


Pseudobinary (quasibinary) sections describe equilibria and can be read in the same way as binary diagrams.The notation used in pseudobinary systems is the same as that of vertical sections, which are reported under“Temperature – Composition Sections”.


The invariant equilibria of a system are listed in the table “Invariant Equilibria” and, where possible, aredescribed by a constitutional “Reaction Scheme” (Fig. 2).

The sequential numbering of invariant equilibria increases with decreasing temperature, one numberingfor all binaries together and one for the ternary system.

Equilibria notations are used to indicate the reactions by which phases will be– decomposed (e- and E-type reactions)– formed (p- and P-type reactions)– transformed (U-type reactions)For transition reactions the letter U (Übergangsreaktion) is used in order to reserve the letter T to denote

temperature. The letters d and D indicate degenerate equilibria which do not allow a distinction accordingto the above classes.


The phase equilibria are commonly shown in triangular coordinates which allow a reading of theconcentration of the constituents in at.%. In some cases mass% scaling is used for better data readability(see Figs. 3 and 4).

In the polythermal projection of the liquidus surface, monovariant liquidus grooves separate phaseregions of primary crystallization and, where available, isothermal lines contour the liquidus surface (seeFig. 3).

Isothermal Sections

Phase equilibria at constant temperatures are plotted in the form of isothermal sections (see Fig. 4).

Temperature – Composition Sections

Non-pseudobinary T-x sections (or vertical sections, isopleths, polythermal sections) show the phasefields where generally the tie lines are not in the same plane as the section. The notation employed for thelatter (see Fig. 5) is the same as that used for binary and pseudobinary phase diagrams.

Thermodynamics

Experimental ternary data are reported in some system reports and reference to thermodynamicmodelling is made.


Noteworthy physical and chemical materials properties and application areas are briefly reported if theywere given in the original constitutional and phase diagram literature.

Miscellaneous

In this section noteworthy features are reported which are not described in preceding paragraphs. Theseinclude graphical data not covered by the general report format, such as lattice spacing – composition data,p-T-x diagrams, etc.

XIV


MSIT®

Introduction

Fig

ure

2:

T

ypic

al r

eact

ion s

chem

e

Ag-T

lT

l-B

iB

i-A

gA

g-T

l-B

i

(Tl)

(h)

(T

l)(r

),(A

g)

23

4d

1l

(A

g)

+ (

Bi)

26

1e 5

(Ag

) +

(T

l)(h

) +

Tl 3

Bi

L +

Tl 3

Bi

(A

g)

+ (

Tl)

(h)

28

9U

1

l (

Ag

)+(T

l)(h

)

29

1e 3

l (

Tl)

(h)+

Tl 3

Bi

30

3e 1

l (

Bi)

+T

l 2B

i 3

20

2e 7

l T

l 3B

i+T

l 2B

i 3

19

2e 8

(Tl)

(h)

Tl 3

Bi+

(Tl)

(r)

14

4e 9

L (

Ag

) +

Tl 3

Bi

29

4e 2

(max

)

L (

Ag

) +

(T

l)(h

)

28

9e 4

(min

)

L (

Ag

) +

Tl 2

Bi 3

20

7e 6

(max

)

(Ag

)+(B

i)+

Tl 2

Bi 3

L (

Ag)+

(Bi)

+T

l 2B

i 31

97

E1

(Ag

)+(T

l)(r

)+T

l 3B

i

(Tl)

(h)

Tl 3

Bi

+ (

Tl)

(r),

(Ag

)1

44

D1

(Ag

)+T

l 3B

i+T

l 2B

i 3

L (

Ag

)+T

l 3B

i+T

l 2B

i 31

88

E2

seco

nd b

inar

y

eute

ctic

rea

ctio

nfi

rst

bin

ary e

ute

ctic

rea

ctio

n

(hig

hes

t te

mper

ature

)te

rnar

y m

axim

um

reac

tion

tem

per

ature

of

26

1°C

mo

no

var

iant

equil

ibri

um

sta

ble

do

wn

to

lo

w

tem

per

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s

seco

nd

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ary

eute

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reac

tion a

t 144°C

XV


MSIT®

Introduction

20

40

60

80

20 40 60 80

20

40

60

80

A B

C Data / Grid: at.%

Axes: at.%

δ700

p1

500

400

400°C

γ

300

U e1

700

500

β(h)

400

300

E

300

α400

e2

500°C isotherm, temperature is usualy in °C

liquidus groove to decreasing temperatures

estimated 400°C isotherm

limit of known region

ternary invariantreaction

binary invariantreaction

primary γ-crystallization

20

40

60

80

20 40 60 80

20

40

60

80

Al B

C Data / Grid: mass%

Axes: mass%

L+γ

γ+β(h)

L+γ+β(h)

β(h)

L+β(h)

L

L+α

α

phase field notation

estimated phase boundary

tie line

three phase field (partially estimated)

experimental points(occasionally reported)

limit of known region

phase boundary

γ

Fig. 3: Hypothetical liquidus surface showing notation employed

Fig. 4: Hypothetical isothermal section showing notation employed

XVI


MSIT®

Introduction

References

The publications which form the bases of the assessments are listed in the following manner:[1974Hay] Hayashi, M., Azakami, T., Kamed, M., “Effects of Third Elements on the Activity of Lead

in Liquid Copper Base Alloys” (in Japanese), Nippon Kogyo Kaishi, 90, 51-56 (1974) (Experimental,Thermodyn., 16)

This paper, for example, whose title is given in English, is actually written in Japanese. It was publishedin 1974 on pages 51- 56, volume 90 of Nippon Kogyo Kaishi, the Journal of the Mining and MetallurgicalInstitute of Japan. It reports on experimental work that leads to thermodynamic data and it refers to 16 cross-references.

Additional conventions used in citing are:# to indicate the source of accepted phase diagrams* to indicate key papers that significantly contributed to the understanding of the system.Standard reference works given in the list “General References” are cited using their abbreviations and

are not included in the reference list of each individual system.

60 40 200

250

500

750

A 80.00B 0.00C 20.00

A 0.00B 80.00C 20.00Al, at.%

Tem

pera

ture

, °C

L

32.5%L+β(h)

β(r) - room temperature

β(r)

L+α+β(h)

α+β(h)

α

L+α

phase field notation

concentration ofabscissa element

alloy compositionin at.%

β(h)

modification

β(h) - high temperaturemodification188

temperature, °C

Fig. 5: Hypothetical vertical section showing notation employed

XVII


MSIT®

Introduction

General References

[C.A.] Chemical Abstarts - pathways to published research in the world's journal and patentliterature - http://www.cas.org/

[Curr.Cont.] Current Contents - bibliographic multidisciplinary current awareness Web resource - http://www.isinet.com/products/cap/ccc/

[E] Elliott, R.P., Constitution of Binary Alloys, First Supplement, McGraw-Hill, New York(1965)

[G] Gmelin Handbook of Inorganic Chemistry, 8th ed., Springer-Verlag, Berlin [H] Hansen, M. and Anderko, K., Constitution of Binary Alloys, McGraw-Hill, New York

(1958) [L-B] Landolt-Boernstein, Numerical Data and Functional Relationships in Science and

Technology (New Series). Group 3 (Crystal and Solid State Physics), Vol. 6, Eckerlin, P.,Kandler, H. and Stegherr, A., Structure Data of Elements and Intermetallic Phases (1971);Vol. 7, Pies, W. and Weiss, A., Crystal Structure of Inorganic Compounds, Part c, KeyElements: N, P, As, Sb, Bi, C (1979); Group 4: Macroscopic and Technical Properties of

Matter, Vol. 5, Predel, B., Phase Equilibria, Crystallographic and Thermodynamic Data of

Binary Alloys, Subvol. a: Ac-Au ... Au-Zr (1991); Springer-Verlag, Berlin. [Mas] Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, ASM, Metals Park, Ohio (1986) [Mas2] Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International,

Metals Park, Ohio (1990) [P] Pearson, W.B., A Handbook of Lattice Spacings and Structures of Metals and Alloys,

Pergamon Press, New York, Vol. 1 (1958), Vol. 2 (1967) [S] Shunk, F.A., Constitution of Binary Alloys, Second Supplement, McGraw-Hill, New York

(1969) [V-C] Villars, P. and Calvert, L.D., Pearson's Handbook of Crystallographic Data for

Intermetallic Phases, ASM, Metals Park, Ohio (1985) [V-C2] Villars, P. and Calvert, L.D., Pearson's Handbook of Crystallographic Data for

Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)

1


MSIT®

Al–Fe–V

Aluminium – Iron – Vanadium

Gautam Ghosh

Literature Data

[1960Gup] studied the effect of additions of Al on the stability of the -phase (FeV). They prepared a

number of alloys, using electrolytic grade elements, in an induction furnace under He atmosphere. The

alloys were homogenized at 1175°C for 72 h. Metallographic observations and X-ray diffraction were

performed to identify the phases. [1987Sok] and [1988Sok] reported the phase equilibria in the Al-rich

ternary alloys containing up to about 50 at.% Fe. The alloys were prepared using the metals of following

purity: 99.95 mass% Al, 99.95 mass% Fe and electrolytic V. A number of alloys were prepared by arc

melting under Ar followed by homogenization at 500°C in evacuated silica capsules. The V-rich alloys (0 to

75 at.% Al) were heat treated for 1800 h at 1000°C followed by 600 h at 500°C, whereas Al-rich alloys (75

to 100 at.% Al) were annealed at 500°C for 1430 h [1988Sok]. The phase analysis was performed by means

of microstructural, thermal analysis, microhardness and X-ray diffraction techniques. Apart from

conventional casting, a number of ternary alloys were also subjected to rapid solidification by melt-spinning

which were subsequently annealed at 250 and 450°C for 50 h. An additional rapidly quenched alloy was

investigated by Mössbauer spectroscopy [1989Sok]. These results were assessed by [1992Gho] and

[1992Rag].

Recent experimental results are primarily related to phase separations [1989Zha, 1989Koz, 1993Miy,

1994Koz] and bcc-based ordering in Fe-rich alloys [1983Bus, 1985Okp, 1995Ant, 1997Nis1, 1997Nis2,

2001Nis1, 2001Nis2]. An update summarizing some of these results has been reported by [2002Rag].

Binary Systems

The Al-Fe and Al-V binary phase diagrams are accepted from [2003Pis] and [2003Sch], respectively. The

Al-Fe phase diagram has undergone slight modification due to recently established congruent melting

behavior of the Fe4Al13 phase [1986Len]. The Fe-V phase diagram is accepted from [1982Kub], which has

also been adopted in [Mas].

Solid Phases

The maximum equilibrium solid solubilities of V and Fe in (Al) are about 0.3 at.% at 660.4°C [1989Mur]

and 0.03 at.% at 652°C [1982Kub], respectively. However, by rapid solidification, the corresponding solid

solubilities can be enhanced up to about 1.25 at.% V and 4.4 at.% Fe [1976Mon] and in the ternary regime,

the solid solubility can be up to 0.5 at.% V and 2 at.% Fe [1987Sok]. The lattice parameter of supersaturated

(Al) containing about 4.4 at.% Fe is about 401.2 pm [1976Mon]. Also, the lattice parameter of (Al)

decreases linearly to 404.2 pm at 1.2 at.% V [1976Mon].

The substitution of Fe by V in Fe3Al increases both the D03 (Fe3Al) B2 (FeAl) and the B2 (FeAl) A2

( Fe) transition temperatures [1969Bul]. Recently, the effect of V on the D03 B2 ordering of Fe3Al has

been determined by several investigators [1997Nis1, 1997Nis2, 2001Nis1, 2001Nis2]. These results are

summarized in Fig. 1. The D03 B2 temperatures reported by [1969Bul] differ significantly from those of

Nishino and co-workers, as a result the data of [1969Bul] are not considered in Fig. 1.

Along the Fe3Al-V3Al section, solid solutions (Fe1-xVx)3Al have been prepared [2003Kaw1]. The D03

lattice of Fe3Al (x = 0) has three sublattices labeled Al (4 sites), FeI (4 sites) and FeII (8 sites). V has a strong

tendency to occupy the FeI sublattice as shown by X-ray absorption fine-structure [1997Nis1], and this

leads to the formation of Heusler phase at the ideal composition of VFe2Al [1976Vla, 1983Bus, 1985Okp,

1997Nis1, 1997Nis2, 2001Nis1, 2001Nis2]. While the addition of V in Fe3Al increases D03 B2 ordering

temperature, the Curie temperature of D03 decreases monotonically [2001Kan]. This is shown in Fig. 2.

Another consequence of substitution Fe by V is the decrease of lattice parameter of Fe3Al down to a

2


MSIT®

Al–Fe–V

minimum at the ideal Heusler composition of VFe2Al beyond which it increases [2001Nis1, 2001Nis2].

This behavior shown in Fig. 3.

[1969Bul] also reported the D03 B2 and B2 A2 ordering temperatures along Fe3Al-VFe3 section, both

showing increasing trend as V is substituted for Al as shown in Fig. 4. However, in view of the above

mentioned discrepancy, further measurements are needed to verify the results of [1969Bul].

As expected, V also increases D03 B2 ordering temperature of other Al-Fe alloys in the vicinity of Fe3Al.

For example, [1995Ant] prepared three alloys VFe73Al26, V2Fe72Al26 and V4Fe70Al26 and measured the

ordering temperature using DTA. The D03 B2 temperature transition of these alloys are 585, 624 and

695°C, respectively.

[1997And] determined site occupancy of V in V5Fe50Al45 ( 1) by ALCHEMI (Atom Location by

CHanneling Enhanced MIcroanalysis) in TEM. [1997And] observed that about 80% of the “Al-site” is

occupied by V, and the residual “Fe-site” is attributed to the kinetics of site-equilibrium mechanism.

The Fe4Al13 phase can dissolve about 5 at.% V at 500°C [1987Sok] and about 2 at.% V at room temperature

[1981Yin]. At 500°C, the VAl3, V4Al23, V7Al45 and V2Al21 phases can dissolve up to about 6.5, 2.0, 1.7

and 4.5 at.% Fe, respectively [1987Sok]. The V solubilities in Fe2Al5, FeAl2 and FeAl were reported to be

about 3, 1.7 and 10 at.% V, respectively [1988Sok]. However, [2000Sah] uses, in the Al-rich corner at

475°C, a diagram in which V4Al23 dissolves up to 4 at % Fe and Fe4Al13 dissolves up to 8 at.% V.

In contrast to the results of [1987Sok], Skinner et al. [1988Ski] reported that melt-spinning of Al-rich alloys

containing up to 16 at.% Fe and 10 at.% V gives rise to a quasicrystalline icosahedral phase. Also, [1988Ski]

suggested that the lattice parameter of such an icosahedral phase is dependent on the Fe:V ratio in the alloy.

Rapidly quenched alloys of the compositions 94Al-6Fe (at.%) and 95.3Al-4Fe-0.7V (at.%), which

consisted of (Al) + slight amounts of FeAl6 were investigated by Mössbauer spectroscopy. Two kinds of

coordination of Fe atoms in the Al lattice, a symmetric and an asymmetric one, were observed in the V

containing alloy. In contrast to this result the Al-Fe alloy had shown only one kind of coordination

[1989Sok]. The details of the crystal structures and lattice parameters of the solid phases are listed

in Table 1.

Isothermal Sections

[1960Gup] reported the phase boundaries involving and ( Fe) phases in the form of a partial isotherm at

1175°C. Al is a strong phase destabilizer; about 0.5 at.% Al at 1175°C is reported to be sufficient to

suppress the phase completely. [1994Koz] prepared ribbons of Fe rich alloys by melt-spinning. The

samples were annealed at 500C for 240ks, and were examined in a transmission electron microscope.

Figure 5 shows the partial Al-Fe-V isothermal section at 500°C from [1987Sok] and [1988Sok]. It should

be mentioned that the Al-V binary phases VAl6, VAl7 and VAl11 as designated by [1987Sok, 1988Sok],

correspond to V4Al23, V7Al45 and V2Al21 in the presently accepted Al-V phase diagram. [2000Sah]

presents, in the Al rich corner at 475°C a diagram in which the solubility of Fe in V2Al21 is very low so that

Al may be in equilibrium with VAl10 and V7Al45 phases, which contradicts the observations of [1987Sok,

1988Sok]. Figure 1 also includes the results of TEM analyses on Fe-rich samples annealed at 500°C

[1989Zha, 1994Koz]. Three types of phase separation sequences from the single phase regions of the , 1

and 2 phases into the + 1 phase region have been distinguished [1989Zha]. [1987Sok] also reported the

phases obtained in the as-melt-spun condition as well as after annealing at 250 and 450°C for 50 h. Their

results are summarized in Table 2. It was noted that, except for the ternary alloy containing more than 16.5

at.% Fe and 3.6 at.% V which was annealed at 450°C for 50 h, equilibrium was not reached in the rest of

the alloys after the annealing treatments used by the authors. For example, after annealing the binary Al-V

and Al-Fe melt-spun alloys at both 250 and 450°C, the authors obtained (Al+VAl3+V2Al21) and

(Al+Fe4Al13+FeAl6) phases, respectively. In the latter case, FeAl6 represents a metastable phase.

3


MSIT®

Al–Fe–V

Thermodynamics

[2003Kaw2] measured various thermophysical (dilatability, compressibility) and thermochemical

properties of VFe2Al, and proposes, for the heat capacity, the following expression:

Cp/J mol-1 K-1 = 229 - 0.328 T + 2.50 10-3 T 2 - 5.63 106 T - 2.

[1994Koz] constructed the free energies of A2, B2 and D03 phases by a statistical approach employing

Bragg-Williams-Gorsky approximation. They considered both atomic and magnetic interaction energies up

to second nearest neighbor. Based on the model description of free energies, they calculated isothermal

section at 500°C which is good agreement with the experimentally observed microstructures of Fe-rich

alloys.


Magnetic and electrical properties of V1-xFe2+xAl alloys have been studied extensively [1985Okp,

1997Nis1, 1997Nis2, 1998Kat, 1998Weh, 2000Kat, 2000Zar, 2001Fen, 2001Han, 2001Kan, 2001Lue,

2001Mak, 2001Nis1, 2001Nis2, 2001Sum, 2003Kaw1]. An important finding is that VFe2Al is nonmetallic

with respect to transport properties while it is metallic with respect to its thermodynamic properties. For

example, [1997Nis2] observed an anomalous negative temperature dependence of electrical resistivity such

that it behaves almost like a semiconductor. This is despite the fact that it has a large density of states at the

Fermi level as revealed by the photoemission valence-band spectra. VFe2Al is non-magnetic semimetal

with a sharp pseudogap at the Fermi level [2000Kat]. It has been reported that a strong hybridization of Fe-

and V-3d states causes a broadening of the d-states and their shift to the higher binding energy. As a result

long-range magnetic order disappears and a narrow energy gap near the Fermi level is formed [2000Zar].

The unusual electron transport is mainly attributed to the effect of strong spin fluctuations, in addition to the

existence of very low carrier concentrations [2000Kat].

[1962Min] studied the effect of V addition on the properties of Fe3Al. Addition of V increases hardness,

electrical resistivity and also improves the high temperature mechanical properties. [2001Nis1] reported the

mechanical properties of the (VxFe1-x)3Al alloys. In the composition range 0 x 0.38, the room

temperature yield stress exhibits a double-well behavior starting from 550 MPa for Fe3Al with a first

minimum at 150MPa for x = 0.02, a maximum at 300 MPa for x = 0.15 and a second minimum at 150 MPa

for x = 0.333 corresponding to the composition VFe2Al. Furthermore, [2001Nis1] observed a correlation

between the yield stress peak at higher temperature and the loss of D03 order. [2000Ino] reported a

significant increase in strength of rapidly solidified Al-Fe-V alloys containing nano-quasicrystalline phase.

Miscellaneous

From a preliminary investigation of the section Fe4Al13-V2Al21, a eutectic reaction was claimed to exist at

~610°C with an invariant composition at ~83 at.% Al [1988Sok].

References

[1960Gup] Gupta, K.P., Rajan, N.S., Beck, P.A., “Effect of Si and Al on the Stability of Certain

Phases”, Trans. Met. Soc. AIME, 218, 617-624 (1960) (Equi. Diagram, Experimental,

#, *, 18)

[1962Min] Mints, R.S., Samsonova, N.N., Malkov, Y. S., “The Effects of Elements of Group V in the

Periodic System (V, Nb, Ta) on the Properties of Fe3Al” (in Russian), Dop. Akad. Nauk

Ukrain. RSR, 144, 1324-1327 (1962) (Experimental, 1)

[1969Bul] Bulycheva, Z.N., Svezhova, S.I., Kondrat’ev, V.K., “Change in the Ordering Temperature

of Fe3Al on Adding a Third Element” (in Russian), Ukrain. Fiz. Zhur., 14, 1706-1708

(1969) (Crys. Structure, Experimental, 5)

[1976Mon] Mondolfo, L.F., “Aluminum-Vanadium System”, in “Aluminium Alloys: Structure and

Properties”, Butterworths, London, 392-394 (1976) (Review, 46)

4


MSIT®

Al–Fe–V

[1976Vla] Vlasova, E.N., Prokoshin, A.F., “Formation of L12 Substructure and Stratification in Solid

Fe-Cr Solutions Doped with Al and V” (in Russian), Dokl. Akad. Nauk SSSR, 231, 599-602


[1981Yin] Ying-Hong, Z., Jing-Qi, L., Jiang-Xuang, Z., Cheng, C.S., “A Room-Temperature Section

of the Phase Diagram of TiAl3-VAl3-MAl3 of the System Alloys of Al-Ti-V-M (M = Ni,

Fe)”, Acta Phys. Sin. (Chin. J. Phys.), 30, 972-975 (1981) (Crys. Structure, Experimental,

Equi. Diagram, 4)

[1982Kub] Kubaschewski, O., “Fe-V”, in “Iron-Binary Phase Diagrams”, Springer Verlag, Berlin,

160-164 (1982) (Equi. Diagram, #, 15)

[1983Bus] Buschow, K.H.J., van Engen, P.G., Jongebreur, R., “Magneto-Optical Properties of Metallis

Ferromagnetic Materials”, J. Magn. Magn. Mater., 38, 1-22 (1983) (Magn. Prop., Optical

Prop., 23)

[1985Okp] Okpalugo, D.E., Both, J.G., Faunce, C.A., “Onset of Ferromagnetism in 3d-Substituted

Fe-Al Alloys. I: Ti, V and Cr Substitutions”, J. Phys. F, Met. Phys., 15, 681-692 (1985)

(Crys. Structure, Experimental, 21)

[1986Len] Lendvai, A., “Phase Diagram of Al-Fe Sytem up to 45 mass% Iron”, J. Mater. Sci. Lett., 5,

1219-1220 (1986) (Equi. Diagram, Experimental, #, *, 7)

[1987Sok] Sokolovskaya, E.M., Badalova, L.M., Kazakova, E.F., “Phase Composition of Rapidly

Quenched Alloys of the System Al-Fe-V”, Izv. Akad. Nauk SSSR, Met., (5), 212-215 (1987)

(Equi. Diagram, Experimental, #, *, 7)

[1988Ski] Skinner, D.J., Ramanan, V.R.V., Zedalis, M.S., Kim, W.J., “Stability of Quasicrystalline

Phases in AlFeV Alloys”, Mater. Sci. Eng., 99, 407-411 (1988) (Crys. Structure,

Experimental, 8)

[1988Sok] Sokolovskaya, E.M., Badalova, L.M., Kazakova, E.F., Stroeva, N.V., “Interactions of

Intermetallic Compounds in the Ternary System Aluminum-Iron-Vanadium” (in Russian),

Vestn. Mosk. Univ., Ser. 2: Khim., 29(3), 303-306 (1988) (Experimental, 5)

[1989Koz] Kozakai, T., Zhao, P.Z., Miyazaki, T., “Phase Separations in Fe-Rich Fe-Base Ternary

Ordering Alloy Systems”, Met. Abstr. Light Metals and Alloys, 23, 32-33 (1989/1990)

(Crys. Structure, Equi. Diagram, Experimental, 0)

[1989Mur] Murray, J.L., “Al-V (Aluminum-Vanadium)”, Bull. Alloy Phase Diagrams, 10(4), 351-357

(1989) (Crys. Structure, Equi. Diagram, Review, 34)

[1989Sok] Sokolovskaya, E.M., Badalova, L.M., Kazakova, E.F., Reiman, S.I., Ryaskyi, G.K.,

Sorokin, A.A., Philipova, A.A., Chaldieva, G.M., “Investigation of Chemical Composition

Microcrystalline of an Al Alloys with Transition Metals” (in Russian), Vestn. Mosk. Univ.,

Ser. 2: Khim., 30(2), 162-165 (1989) (Crys. Structure, Experimental, 6)

[1989Zha] Zhao, P.Z., Kozakai, T., Miyazaki, T., “Phase Separation into A2+D03 Two Phases in

Iron-Aluminium-Vanadium Ternary Ordering Alloys” (in Japanese), Nippon Kinzoku

Gakkai Shi, 53(3), 266-272 (1989) (Crys. Structure, Equi. Diagram, Experimental, #, *, 23)

[1992Gho] Ghosh, G., “Aluminium-Iron-Vanadium”, MSIT Ternary Evaluation Program, in MSIT

Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH,

Stuttgart; Document ID: 10.19022.1.20, (1992) (Crys. Structure, Equi. Diagram,

Assessment, 14)

[1992Rag] Raghavan, V., “The Al-Fe-V (Aliminium-Iron-Vanadium) System”, in Phase Diagram of

Ternary Iron Alloys, Part 6A, Ind. Inst. Metals, Calcutta, 204-207 (1992) (Review, Equi.

Diagram, 7)

[1993Miy] Miyazaki, T., “Phase Diagrams of Iron-Base Ternary Ordering Alloy Systems”, Comput.

Aided Innovation New Mater. 2, Proc. Int. Conf. Exhib. Comput. Appl. Mater. Mol. Sci.

Eng., 2nd 1992 (Pub. 1993) (Pt.1), 707-712., 2ND1992 (1993) (Equi.Diagram)

[1994Koz] Kozakai, T., Miyazaki, T., “Experimental And Theoretical Investigations on Phase

Diagrams of Fe Base Ternary Ordering Alloys”, ISIJ Int., 34(5), 373-383 (1994)

(Calculation, Equi. Diagram, Experimental, Magn. Prop., #, *, 18)

5


MSIT®

Al–Fe–V

[1995Ant] Anthony, L., Fultz, B., “Effects of Early Transition Metal Solutes in the D03-B2 Critical

Temperature of Fe3Al”, Acta Metall. Mater., 43, 3885-3891 (1995) (Crys. Structure,

Experimental, 35)

[1997And] Anderson, I.M., “Alchemi Study of Site Distributions of 3d-Transition Metals in

B2-Ordered Iron Aluminides”, Acta Mater., 45(9), 3897-3909 (1997) (Calculation, Crys.

Structure, Experimental, Theory, 26)

[1997Nis1] Nishino, Y., Kumada, C., Asano, S., “Phase Stability of Fe3Al with Addition of 3d

Transition Elements”, Scr. Mater., 36, 461-466 (1997) (Crys. Structure, Equi. Diagram,

Experimental, 26)

[1997Nis2] Nishino, Y., Kato, M., Asano, S., Soda, K., Hayasaki, M., Mizutani, U.,

“Semiconductor-Like Bahavior of Electrical Resisitivity in Heusler-Type Fe2VAl

Compound”, Phys. Rev. Lett., 79(10), 1909-1912 (1997) (Crys. Structure, Experimental, 18)

[1998Kat] Kato, M., Nishino, Y., Asano, S. Ohara, S., “Electrical Resistance Anomaly and Hall Effect

in (Fe1-xVx)3Al Alloys” (in Japanese), J. Japan. Inst. Met., 62(7), 669-674 (1998) (Crys.

Structure, Experimental, 23)

[1998Weh] Weht, R., Pickett, W.E., “Excitonic Correlations in the Intermetallic Fe2VAl”, Phys. Rev. B,

58(11), 6855-6861 (1998) (Calculation, Crys. Structure, Mechan. Prop., 21)

[2000Ino] Inoue, A., Kimura, H.M., Zhang, T., “High-Strength Aluminium- and Zirconium-Based

Alloys Containing Nanoquasicrystalline Particles”, Mater. Sci. Eng. A, 294-296, 727-735

(2000) (Crys. Structure, Experimental, Mechan. Prop., 28)

[2000Kat] Kato, M., Nishino, Y., Mizutani, Y., Asano, S., “Electronic, Magnetic and Transport

Properties of (Fe1-xVx)3Al Alloys”, J. Phys.: Condens. Matter, 12, 1769-1779 (2000) (Crys.

Structure, Electr. Prop., Experimental, Magn. Prop., Phys. Prop., 33)

[2000Sah] Sahoo, K.L., Sivaramakrishnan, C.S., Chakrabarti, A.K., “Solidification Characteristics of

the Al-8.3Fe-0.8V-0.9Si Alloy”, Metall. Mater. Trans. A, 31A(6), 1599-1610 (2000)

(Experimental, #, 21)

[2000Zar] Zarek, W., Talik, E., Heimann, J., Kulpa, M., Winiarski, A., Neumann, M., “Electronic

Structure, Magnetic and Electrical Properties of Fe3-xVxAl Compounds”, J. Alloys Compd.,

297, 53-58 (2000) (Crys. Structure, Electr. Prop., Experimental, Magn. Prop., 15)

[2001Fen] Feng, Y., Rhee, J.Y., Wiener, T.A., Lynch, D.W., Hubbard, B.E., Sievers, A.J.,

Schlagel, D.L., Lograsso, T.A., Miller, L.L., “Physical Properties of Heusler-Like Fe2VAl”,

Phys. Rev. B, 63(16), 165109-1-165109-12 (2001) (Crys. Structure, Electr. Prop.,

Experimental, Magn. Prop., Phys. Prop., 30)

[2001Han] Hanada, Y., Suzuki, R.O., Ono, K., “Seebeck Coefficient of (Fe,V)3Al Alloys”, J. Alloys

Compd., 329, 63-68 (2001) (Electr. Prop., Experimental, 18)

[2001Kan] Kanomata, T., Sasaki, T., Hoshi, T., Narita, T., Harada, T., Nishihara, H., Yoshida, T.,

Note, R., Koyama, K., Nojiri, H., Kaneko, T., Motokava, M., “Magnetic and Electrical

Properties of Fe2+xV1-xAl”, J. Alloys Compd., 317-318, 390-394 (2001) (Crys. Structure,

Electr. Prop., Experimental, 19)

[2001Lue] Lue, C.S., Ross, J.H., Rathnayaka, Jr., K.D.D., Naugle, D.G., Wu, S.Y., Li, W.-H.,

“Supermagnetism and Magnetic Defects in Fe2VAl and Fe2VGa”, J. Phys.: Condens.

Matter, 13, 1585-1593 (2001) (Crys. Structure, Experimental, Magn. Prop., 25)

[2001Mak] Maksimov, I., Baabe, D., Klauss, H.H., Litterst, F.J., Feyerherm, R., Toebbens, D.M.,

Matsushita, A., Suellow, S., “Structure and Magnetic Order in Fe2+xV1-xAl”, J. Phys.:

Condens. Matter, 13, 5487-5501 (2001) (Crys. Structure, Experimental, Magn. Prop., 25)

[2001Nis1] Nishino, Y., “Electronic Structure and Transport Properties of Pseudogap System Fe2VAl”,

Mater. Trans., JIM, 42(6), 902-910 (2001) (Crys. Structure, Electr. Prop., Equi. Diagram,

Experimental, 58)

[2001Nis2] Nishino, Y., Makino, Y., “Effect of Vanadium Substitution on Strength Properties of

Fe3Al-Based Alloys”, Mater. Sci. Eng. A, 319-321, 368-371 (2001) (Equi. Diagram,

Experimental, Mechan. Prop., #, *, 29)

6


MSIT®

Al–Fe–V

[2001Sum] Sumi, H., Kato, M., Nishino, Y., Asano, S., Mizutani, U., “Electrical Resistivity Anomaly

and Magnetic Properties in Heusler-Type Fe2VAl Alloy” (in Japanese), J. Jpn. Inst. Met.,

65(9), 771-774 (2001) (Crys. Structure, Electr. Prop., Experimental, Magn. Prop.,

Thermodyn., 16)

[2002Rag] Raghavan, V., “Al-Fe-V (Aliminum-Iron-Vanadium) System”, J. Phase Equilib., 23,

439-440 (2002) (Equi. Diagram, Review, 7)

[2003Kaw1] Kawaharada, Y., Kurosaki, K., Yamanaka, S., “High Temperature Thermoelectric

Properties of (Fe1-xVx) 3Al Heusler Type Compounds”, J. Alloys Compd., 349(1-2), 37-40

(2003) (Electr. Prop., Experimental, Mechan. Prop., Phys. Prop., 27)

[2003Kaw2] Kawarahada, Y., Kurosaki, K., Zamanaka, S., “Thermophysical Properties of Fe2VAl”,

J. Alloys Compd., 352, 48-51 (2003) (Thermodyn., Phys. Prop., Mechan. Prop.,

Experimental, 22)

[2003Pis] Pisch, A., “Al-Fe (Aluminum-Iron)”, MSIT Binary Evaluation Program, in MSIT


Stuttgart; to be published, (2003) (Equi. Diagram, Review, 58)

[2003Sch] Schuster, J.C., “Al-V (Aluminium-Vanadium)”, MSIT Binary Evaluation Program, in MSIT


Stuttgart; to be published, (2003) (Equi. Diagram, Review, 31)

Table 1: Crystallographic Data of Solid Phases

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Al)

660.452

cF4

Fm3m

Cu

a = 404.96 pure Al at 25°C [Mas2]

, ( Fe)

912

cI2

Im3m

W

a = 286.65 pure Fe at 25°C [Mas2]

(V)

1910

cI2

Im3m

W

a = 302.40 pure V at 25°C [Mas2]

V5Al8 1408

cI52

I43m

Cu5Zn8

a = 923.0

a = 921.8

[2003Sch], Al-rich

[2003Sch], V-rich

solid solubility ranges

from 60.0 to 66.0 at.% Al

VAl3 1270

tI8

I4/mmm

TiAl3

a = 378.14

c = 832.2

a = 378.07

c = 830.9

[2003Sch], Al-rich limit

[2003Sch], V-rich limit

solubility ranges from 74 to 75 at.% Al

V4Al23

736

hP54

P63/mmc

V4Al23

a = 769.28

c = 1704.0

a = 769.9

c = 1705.3

[1989Mur]

[2003Sch]

7


MSIT®

Al–Fe–V

V7Al45

730

mC104

C2/m

V7Al45

a = 2540

b = 759

c = 1100

= 127

a = 2563.0

b = 763.7

c = 1108.8

= 128.83

[1989Mur]

[2003Sch]

V2Al21

690

cF184

Fd3m

V2Al21

a = 1449.2

a = 1452.1

[1989Mur, 2003Sch]

V3Al(r)

650

cP8

Pm3n

Cr3Si

a = 482.9 [2003Sch]

1, Fe3Al

547

cF16

Fm3m

BiF3

a = 578.86-579.30 [2003Pis], solid solubility ranges

from ~24 to ~37 at.% Al

2, FeAl

1310

cP2

Pm3m

CsCl

a = 289.76-290.78 [2003Pis], at room temperature



, Fe2Al31102 - 1232

cI16? a = 598.0 [2003Pis], solid solubility

ranges from 54.5 to 62.5 at.% Al

FeAl2 1156

aP18

P1

FeAl2

a = 487.8

b = 646.1

c = 880.0

= 91.75°

= 73.27°

= 96.89°

[2003Pis], at 66.9 at.% Al



Fe2Al5 1169

oC24

Cmcm

a = 765.59

b = 641.54

c = 421.84

[2003Pis], at 71.5 at.% Al



Fe4Al13

1160

mC102

C2/m

Fe4Al13

a = 1552.7-1548.7

b = 803.5-808.4

c = 1244.9-1248.8

= 107.7-107.99°

a = 1549.2

b = 807.8

c = 1247.1

= 107.69°

[2003Pis], 74.16 to 76.7 at.% Al



Sometimes called FeAl3 in the literature

[2003Pis], at 76.0 at.% Al

, VFe

1252

tP30

P42/mnm

CrFe

a = 895.6

c = 462.7

[V-C2], solid solubility

ranges from 33.5 to 64.0 at.% V

VFe2Al cF16

Fm3m

BiF3

a = 576.1

a = 576.16

a = 576.0

a = 576.19

[1983Bus] Heussler Alloy

[1998Kat]

[2001Lue]

[2001Nis1]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

8


MSIT®

Al–Fe–V

Table 2: Phases Present in the as-melt-spun Condition and after Annealing Treatments

Composition (at.%) As-melt-spun After annealing for 50 h, at [°C]

Al Fe V 250 450

98.0

96.0

94.0

91.0

87.0

98.0

95.0

92.0 e

86.0

97.5

92.5

90.0

88.0

-

-

-

-

-

2.0

5.0

8.0

14.0

2.0

6.5

8.5

10.0

2.0

4.0

6.0

9.0

13.0

-

-

-

-

0.5

1.0

1.5

2.0

(Al)

(Al)+VAl3(Al)+VAl3(Al)+VAl3(Al)+VAl3(Al)+FeAl6(Al)+FeAl6+Fe4Al13

(Al)+FeAl6+Fe4Al13

(Al)+FeAl6+Fe4Al13

(Al)

(Al)+FeAl6(Al)+FeAl6+Fe4Al13

(Al)+FeAl6+Fe4Al13

(Al)+VAl3+V2Al21

(Al)+VAl3+V2Al21

(Al)+VAl3+V2Al21

(Al)+VAl3+V2Al21

(Al)+VAl3+V2Al21

(Al)+FeAl6+Fe4Al13

(Al)+FeAl6+Fe4Al13

(Al)+FeAl6+Fe4Al13

(Al)+FeAl6+Fe4Al13

(Al)

(Al)+FeAl6+Fe4Al13

(Al)+FeAl6+Fe4Al13

(Al)+FeAl6+V2Al21

(Al)+VAl3+V2Al21

(Al)+VAl3+V2Al21

(Al)+VAl3+V2Al21

(Al)+VAl3+V2Al21

(Al)+VAl3+V2Al21

(Al)+FeAl6+Fe4Al13

(Al)+FeAl6+Fe4Al13

(Al)+FeAl6+Fe4Al13

(Al)+FeAl6+Fe4Al13

(Al)+Fe4Al13

(Al)+FeAl6+Fe4Al13

(Al)+FeAl6+Fe4Al13

(Al)+Fe4Al13+V2Al21

50 60 70500

600

700

800

900

1000

1100

V 30.00Fe 45.00Al 25.00

V 0.00Fe 75.00Al 25.00Fe, at.%

Tem

pera

ture

, °C

VFe2Al

B2(α1)

DO3(α2)

Fig. 1: Al-Fe-V.

Variation of D03 B2

ordering temperature

along V3Al-Fe3Al

section

9


MSIT®

Al–Fe–V

60 70-300

-200

-100

0

100

200

300

400

500

V 25.00Fe 50.00Al 25.00

V 0.00Fe 75.00Al 25.00Fe, at.%

Tem

pera

ture

, °C

Fig. 2: Al-Fe-V.

Variation of Curie

temperature of D03

phase along

V3Al-Fe3Al section

45 50 55 60 65 70 75

576.0

576.5

577.0

577.5

578.0

578.5

579.0

579.5

Al

Fe

V

25.0075.00

0.00

Al

Fe

V

25.0045.0030.00

Fe VAl2

Fe, at.%

Lattic

epara

mete

r,pm

Fig. 3: Al-Fe-V.

Variation of lattice

parameter of D03

phase along

Fe3Al-V3Al section

10


MSIT®

Al–Fe–V

20 10400

500

600

700

800

900

V 25.00Fe 75.00Al 0.00

V 0.00Fe 75.00Al 25.00V, at.%

Tem

pera

ture

, °C

BiF3-type

CsCl-type

(αFe)

(α2)

(α1)

Fig. 4: Al-Fe-V.

Variation of

order-disorder

reaction temperature

as a function of V

content along the

VFe3-Fe3Al section

20

40

60

80

20 40 60 80

20

40

60

80

V Fe

Al Data / Grid: at.%

Axes: at.%

α

α+α1

α1

σ

α2

FeAl2

Fe2Al5

Fe4Al13

VAl10V7Al45

V4Al23

VAl3

V5Al8

V3Al

(V)

(Al)

(V)+σ

Fig. 5: Al-Fe-V.

Partial isothermal

section at 500°C

11


MSIT®

Al–Fe–Y

Aluminium – Iron – Yttrium

Gabriele Cacciamani

Literature Data

The Al-Fe-Y phase equilibria have been systematically investigated by [1972Zar] in the 0-33 at.% Y

composition range.

Structural and magnetic properties of the Al-Fe-Y phases have been studied by several authors:

investigations mainly concerned the solid solutions at the Y2(Fe,Al)17 ratio [1976McN, 1996Kuc, 1998Che,

1998Kam, 2001Vor] and the Y(Fe1-xAlx)12 ternary phase [1966Zar, 1974Viv, 1976Bus, 1978Bus, 1980Fel,

1995Sch, 2000Sch, 2001Wae2]. Binary and ternary phases at the Y(Fe,Al)2 atomic ratio have been mainly

investigated by [1972Ryk, 1973Zar, 1975Bus, 1975Dwi, 1976Gro, 1977Mur, 1986Sec, 1988Cun,

2001Wae2]. The YFe2Al10 phase has been studied by [1998Thi, 2001Wae2].

Samples have been generally prepared by arc melting the pure elements (usually 99.9 mass% pure) under

an inert atmosphere. In a few cases other methods were used: synthesis in Al2O3 at 400 to 800°C [1998Thi]

or induction melting of Al-Fe master alloys with appropriate amounts of rare earth [1975Dwi]. Samples

were generally annealed at appropriate temperatures and then quenched.

This evaluation incorporates and continues the critical evaluation made by [1992Gri] considering new

published data.

Binary Systems

The binary systems Al-Fe and Al-Y are accepted from [2003Pis] and [2003Cor], respectively. The Fe-Y

phase equilibria are accepted from the assessment by [1992Zha].

Solid Phases

Crystal structure data are reported in Table 1. Al-Fe binary compounds and phases are not reported to

dissolve Y. Al-Y and Fe-Y phases may show more or less extended solubility ranges due to substitution

between Al and Fe.

The binary Laves phases YAl2 and YFe2 (isostructural, MgCu2 type) dissolve more than 20 at.% of the third

element. At intermediate compositions, however, a different Laves phase ( 1, MgZn2 type) is formed: the

solubility ranges have been studied by [1975Dwi] and crystal structures by [1972Ryk, 1972Zar, 1973Zar,

1975Bus, 1976Gro, 1977Mur, 1986Sec, 1988Cun].

The solid solutions at the Y2(Fe,Al)17 ratio have been studied by different authors [1976McN, 1996Kuc,

1998Che, 1998Kam, 2001Vor]: both Th2Ni17 and Th2Zn17 structures have been reported, but their

composition and temperature ranges of stability are still not well assessed.

The 2 phase has been studied by several authors either at the YFe4Al8, [1976Bus, 1978Bus], YFe6Al6[1980Fel, 1988Che] or different compositions [1995Sch, 1998Sch, 2000Sal, 2000Sch, 2001Wae2]. Also in

this case the solubility range seems to vary appreciably with temperature.

Finally, with the same Y(Fe,Al)12 ratio, a different ternary phase ( 3, at the composition HoFe2Al10) was

first identified by [1972Zar] and then studied by [1998Thi, 2001Wae2].

Isothermal Sections

The partial isothermal section at 500°C is reported in Fig. 1. Determined by [1972Zar], it has been adapted

considering the more recent indications concerning the solubility ranges of the solid solutions (homogeneity

ranges of YFe2 and YAl2 after [1975Dwi]) and the accepted binary systems. The 800°C isothermal section

has been recently investigated by [2001Wae2] in the 50-100 at.% Al region: it resulted to be consistent with

the section by [1972Zar].

12


MSIT®

Al–Fe–Y

Thermodynamics

Thermodynamic properties of the liquid phase have been studied by [1982Erm, 1983Erm1, 1983Erm2].


Mössbauer measurements have been carried out on the 1 [1975Dwi, 1977Mur], 2 [2000Wae, 2001Wae1,

2003Kal] and 3 [2001Wae2] phases.

Magnetic properties have been studied for 2 at different compositions: YFe4Al8 [1978Bus, 1998Hag,

1998Sch, 2000Sik, 2000Wae, 2001Pai, 2001Wae1], YFe6Al6 [1981Fel], YFe5Al7 [1995Sch], YFe7Al5[2000Sch] and variable composition [2000Wae, 2003Kal], and for 1 [1975Bus, 1976Gro, 1977Mur,

1986Sec], 3 [1998Thi], and the phases at the Y2(Fe,Al)17 ratio [1986Plu, 1996Kuc, 1998Che, 1998Kam,

1999Kuc, 2001Kny, 2001Vor]. [2001Kny] investigated also the optical properties of the Y2(Fe,Al)17 solid

solution.

[1988Cun] carried out resistivity measurements on 1 and [1992Joh, 1998All] studied the formation of

amorphous and nano-crystalline alloys in the system.

References

[1958Tay] Taylor, A., Jones, R.M., “Constitution and Magnetic Properties of Iron-Rich

Iron-Aluminium Alloys”, J. Phys. Chem. Solids, 6, 16-37 (1958) (Crys. Structure, Magn.

Prop., Experimental, 49)

[1961Lih] Lihl, F., Ebel, H., “X-Ray Examination fo the Constitution of Iron-Rich Alloys of the

Iron-Aluminium System” (in German), Arch. Eisenhuettenwesen, 32, 483-487, (1961)

(Crys. Structure, Magn. Prop., Experimental, 12)

[1966Zar] Zarechnyuk, O.S., “Ternary Compounds with a ThMn12 Superstructure in the Systems

Yttrium-Transition Metal-Aluminium”, Dop. Akad. Nauk Ukr. RSR, 6, 767-769 (1966)

(Crys. Structure, 2)

[1972Ryk] Rykhal, R.M., “Crystal Structures of the Ternary Compounds YFeAl and YCoAl” (in

Russian), Vestn. L'vov. Univ., Ser. Khim., 13, 11-14 (1972) (Crys. Structure, Experimental,

4)

[1972Zar] Zarechnyuk, O.S., Rikhal', R.M., Ryabov, V.R., Vivchar, O.I., “The Y-Fe-Al Ternary

System in the Region 0 - 33.3 at.% Y”, Izv. Akad. Nauk SSSR, Met., (1), 208 (1972) (Crys.

Structure, Equi. Diagram, Experimental, 12)

[1973Zar] Zarechnyuk, O.S., Rikhal, R.M., Vivchar, O.I., “Laves Phases in Ternary Systems of the

Type Rare-Earth Metal-Transition Metal-Al” (in Russian), Akad. Nauk Ukr. SSR,

Metallofizika, 46, 92-94 (1973) (Crys. Structure, Experimental, 22)

[1974Viv] Vivchar, O.I., Zarechnyuk, O.S., “Compounds of the ThMn12-Type Structure in R-Fe-Al

Systems” (in Russian), Tezisy Dokl. - Vses. Konf. Kristallokhim. Intermet. Soedin., Rykhal,

R.M. (Ed), Vol. 2, L'vov. Gos. Univ.: Lvov, USSR, 41 (1974) (Crys. Structure,

Experimental, 0)

[1975Bus] Buschow, K.H.J., “Crystal Structure and Magnetic Properties of YFe2xAl2-2x”, J.

Less-Common Met., 40, 361-363 (1975) (Crys. Structure, Experimental, 6)

[1975Dwi] Dwight, A.E., Kimball, C.W., Preston, R.S., Taneja, S.P., Weber, L., “Crystallographic and

Moessbauer Study of (Sc, Y, Ln)(Fe, Al)2 Intermetallic Compounds”, J. Less-Common

Met., 40, 285-291 (1975) (Crys. Structure, Moessbauer, Experimental, 8)

[1976Bus] Buschow, K.H.J., van Vucht, J.H.N., van den Haagenhof, W.W., “Note on the Crystal

Structure of the Ternary Rare Earth 3d Transition Metal Compounds of the Type RT4Al8”,

J. Less-Common Met., 50(1), 145-150 (1976) (Experimental, Crys. Structure, 2)

[1976Gro] Groessinger, R., Steiner, W., Krec, K., “Magnetic Investigations of Pseudobinary

RE(Fe,Al)2 Systems (RE = Y, Gd, Dy, Ho)” (in German), J. Magn. Magn. Mater., 2,

196-202 (1976) (Magn. Prop., Experimental, 20)

13


MSIT®

Al–Fe–Y

[1976Mcn] McNelly, D., Oesterreicher, H., “Structural and Low-Temperature Magnetic Studies on

Compounds Sm2Fe17 with Al Substitution for Fe”, J. Less-Common Met., 44, 183-193

(1976) (Crys. Structure, Magn. Prop., Experimental, 26)

[1977Mur] Muraoka, Y., Shiga, M., Nakamura, Y., “Magnetic Properties and Moessbauer Effects of

A(Fe1-xBx)2 (A = Y or Zr, B = Al or Ni) Laves Phase Intermetallic Compounds”, Phys.

Status Solidi A, 42A, 369-374 (1977) (Crys. Structure, Magn. Prop., Moessbauer,

Experimental, 15)

[1978Bus] Buschow, K.H.J., van der Kran, A.M., “Magnetic Ordering in Ternary Rare Earth Iron

Aluminium Compounds (RFe4Al8)”, J. Phys., F: Met. Phys., 8, 921-932 (1978)

(Experimental, Magn. Prop., 9)

[1980Fel] Felner, I., “Crystal Structures of Ternary Rare Earth-3d Transition Metal Compounds of the

RT6Al6 Type”, J. Less-Common Met., 72, 241-249 (1980) (Crys. Structure, 10)

[1981Fel] Felner, I., Seh, M., Rakavy, M., Nowik, I., “Magnetic Order and Hyperfine Interactions in

RFe6Al6 (R = Rare Earth)”, Phys. Chem. Solids, 42, 369-377 (1981) (Crys. Structure, Magn.


[1982Erm] Ermakov, A.F., Esin, Yu.O., Gel'd, P.V., “Partial and Integral Enthalpies of Formation of

Liquid Alloys of Iron Monoaluminide with Yttrium, Lanthanum and Cerium” (in Russian),

Izv. Akad. Nauk SSSR, Met., (5), 69-60 (1982) (Thermodyn., Experimental, 3)

[1983Erm1] Ermakov, A.F., Esin, Yu.O., Levin, E.S., Petrusevskij, M.S., “Estimation of the Enthalpy of

Formation of Liquid Ternary Alloys Fe-Y-Si and Fe-Y-Al from the Data of Characteristic

Boundaries of the Binary Systems” (in Russian), Fiz. Svoistva Met. Splavov (Sverdlovsk),

(4), 68-71 (1983) (Thermodyn., Experimental, 9)

[1983Erm2] Ermakov, A.F., Esin, Yu.O., Levin, E.S., Petrusevskij, M.S., “Assessment of the Enthalpy

of Formation of Iron, Yttrium, Silicon and Iron-Yttrium, Aluminum Liquid Ternary Alloys”

(in Russian), Fiz. Svoistva Met. Splavov (Sverdlovsk), (4), 71-74 (1983) (Experimental,

Thermodyn., 4)

[1985Gan] Gan, R.J., Littlewood, N.T., James, W.J., “Magnetic Structures of Y6(Fe1-xAlx)23

Compounds”, IEEE Trans., Magn., 21(5), 1984-1986 (1985) (Crys. Structure, Magn. Prop.,

Experimental)

[1986Plu] Plusa, D., Pfranger, R., Wyslocki, B., Mydlarz, T., “Magnetic Properties of Y2(Fe1-xAlx)17

Pseudobinary Compounds”, J. Less-Common Met., 120, 1-7 (1986) (Crys. Structure,

Experimental, 11)

[1986Sec] Sechovsky, V., Nozar, P., “Magnetic Phase Diagram of the System Yttrium - Iron -

Aluminum (Y(FexAl1-x)2)”, Acta Phys. Slovaca, 36(3), 210-211 (1986) (Magn. Prop., 3)

[1987Ric] Richter, R., Altounian, Z., Strom-Olsen, J.O., “Y5Al3, A New Y-Al Compound”, J. Mater.

Sci., 22, 2983-2986 (1987) (Experimental, Thermodyn., Crys. Structure, 7)

[1988Che] Chelkowska, G., Chelkowska, A., Winiarska, A., “Magnetic Susceptibility and Structural

Investigations of Rare Earth-Aluminium-Iron (REAl6Fe6) Compounds for RE = Yttrium,

Terbium, Dysprosium, Holmium, and Erbium”, J. Less-Common Met., 143, L7-L10 (1988)

(Crys. Structure, Magn. Prop., Experimental, 12)

[1988Cun] Da Cunha, S.F., Souza, G.P., Takeuchi, A.Y., “Electrical Resistivity of YFeAl

(Y(Fe1-xAlx)2) in the Spin Glass”, J. Magn. Magn. Mater., 73(3), 355-360 (1988) (Crys.

Structure, Electr. Prop., Experimental, 18)

[1989Gsc] Gschneidner Jr, K.A., Calderwood, F.W., “The Al-Y (Aluminium-Yttrium) System”, Bull.

Alloy Phase Diagrams, 10, 44-47 (1989) (Calculation, Equi. Diagram, Crys. Structure,

Review, #, 33)

[1992Gri] Grieb, B., “Aluminium-Iron-Yttrium”, MSIT Ternary Evaluation Program, in MSIT



Assessment, 18)

14


MSIT®

Al–Fe–Y

[1992Joh] Johnson, E., Johansen, A., Sarholt-Kristensen, L., “On Glass Formation in Rapidly

Solidified Aluminium-Based Alloys”, J. Mater. Res., 7(10), 2756-2764 (1992) (Crys.

Structure, Experimental, Phys. Prop., 35)

[1992Zha] Zhang, W., Liu, G., Han, K., “The Fe-Y (Iron-Yttrium) System”, J. Phase Equilib., 13(3),

304-308 (1992) (Equi. Diagram, Thermodyn., Review, #, 29)

[1993Kat] Kattner, U.R., Burton, B.P., “Al-Fe (Aluminum-Iron)“, in “Phase Diagrams of Binary Iron

Alloys”, Okamoto, H. (Ed), ASM International, Materials Park, Ohio 44073-0002, 12-28

(1993) (Equi. Diagram, Review, 99)

[1994Bur] Burkhardt, U., Grin, J., Ellner, M., Peters, K., “Structure Refinement of the Iron-Aluminium

Phase with the Approximate Composition Fe2Al5”, Acta Crystallogr., Sect. B: Struct.

Crystallogr. Crys. Chem., B50, 313-316 (1994) (Crys. Structure, Experimental, 9)

[1994Fol] Foley, J.C., Thoma, D.J., Perepezko, J.H., “Supersaturation of the Al2Y Laves Phase by

Rapid Solidification”, Metall. Mater. Trans. A, 25A, 230-233 (1994) (Crys. Structure,

Experimental, 8)

[1994Gri] Grin, J., Burkhardt, U., Ellner, M., Peters, K., “Refinement of the Fe4Al13 Structure and its

Relationship to Quasihomological Homotypical Structures”, Z. Kristallogr., 209, 479-487


[1995Sch] Schaefer, W., Kockelmann, W., Will, G., Fischer, P., Gal, J., “Neutron Diffraction on

YFe5Al7 as Reference of the f-Magnetism of Isostructural Rare Earth - Iron - Aluminium

Compounds”, J. Alloys Compd., 225, 440-443 (1995) (Crys. Structure, Experimental,

Magn. Prop., 17)

[1996Kuc] Kuchin, A.G., Kourov, N.I., Knyazev, Yu.V., Kleinerman, N.M., Serikov, V.V., Ivanova,

G.V., Ermolenko, A.S., “Electronic, Magnetic, and Structuralproperties of the Alloys

Y2(Fe1-xMx)17 where M = Al and Si”, Phys. Status Solidi A, A155, 479-483 (1996) (Crys.


[1997Kog] Kogachi, M., Haraguchi, T., “Quenched-in Vacancies in B2-Structured Intermetallic

Compound FeAl”, Mater. Sci. Eng. A, A230, 124-131 (1997) (Crys. Structure,

Experimental, 23)

[1998Ali] Aliravci, C.A., Pekgueleryuez, M.O., “Calculation of Phase Diagrams for the Metastable

Al-Fe Phases Forming in Direct-chill (DC)-Cast Aluminium Alloy Ingots”, Calphad, 22,

147-155 (1998) (Calculation, Equi. Diagram, 20)

[1998All] Allen, D.R., Foley, J.C., Perepezko, J.H., “Nanocrystal Development During Primary

Crystallization of Amorphous Alloys”, Acta Mater., 46(2), 431-440 (1998) (Calculation,

Crys. Structure, Equi. Diagram, Experimental, Thermodyn., 39)

[1998Che] Cheng, Z., Shen, B., Yan, Q., Guo, H., Chen, D., Gou, C., Sun, K., de Boer, F.R., Buschow,

K.H.J., “Strcuture, Exchange Interactions, and Magnetic Phase Transition of Er2Fe17-xAlxIntermetallic Compounds”, Phys. Rev. B, 57B(22), 14299-14309 (1998) (Crys. Structure,

Experimental, 35)

[1998Hag] Hagmusa, I.H., Brueck, E., de Boer, F.R., Buschow, K.H.J., “Magnetic Properties of

RFe4Al8 Compounds Studied by Specific Heat Measurements”, J. Alloys Compd., 278,

80-82 (1998) (Thermodyn., Magn. Prop., Experimental, 9)

[1998Kam] Kamimori, T., Koyama, K., Mori, Y., Asano, M., Kinoshita, K., Mochimaru, J., Konishi,

K., Tange, H., “Preferential Site Occupation of M Atoms and the Curie Temperature in

Y2Fe17-xMx (M = Al, Si, Ga)”, J. Magn. Magn. Mater., 177/181, 1119-1120 (1998) (Crys.


[1998Sch] Schobinger-Papamantellos, P., Buschow, K.H.J., Ritter, C., “Magnetic Ordering and Phase

Transitions of RFe4Al8 (R = La, Ce, Y, Lu) Compounds by Neutron Diffraction”, J. Magn.

Magn. Mater., 186, 21-32 (1998) (Crys. Structure, Experimental, Magn. Prop., 13)

[1998Thi] Thiede, V.M.T., Ebel, T., Jeitschko, W., “Ternary Aluminides LnT2Al10 (Ln = Y, La-Nd,

Sm, Cd-Lu and T = Fe, Ru, Os) with YbFe2Al10 Type Structure and Magnetic Properties of

the Iron-Containing Series”, J. Mater. Chem., 8(1), 125-130 (1998) (Crys. Structure, Magn.


15


MSIT®

Al–Fe–Y

[1999Dub] Dubrovinskaia, N.A., Dubrovinsky, L.S., Karlsson, A., Saxena, S.K., Sundman, B.,

“Experimental Study of Thermal Expansion and Phase Transformations in Iron-Rich Fe-Al

Alloys”, Calphad, 23(1), 69-84 (1999) (Equi. Diagram, Experimental, 15)

[1999Kuc] Kuchin, A.G., Medvedeva, I.V., Gaviko, V.S., Kazantsev, V.A., “Magnetovolume

Properties of Y2Fe17-xMx Alloys (M = Si or Al)”, J. Alloys Compd., 289, 18-23 (1999)


[2000Sal] Salamakha, P., Sologub, O., Waerenborgh, J.C., Goncalves, A.P., Godinho, M., Almeida,

M., “Systematical Investigation of the Y-Fe-Al Ternary System. Part 1. Single Crystal

Studies of the YFexAl12-x Compound”, J. Alloys Compd., 296, 98-102 (2000) (Crys.


[2000Sch] Schaefer, W., Barbier, B., Halevy, I., “ThMn12-Type Magnetic ErFe7Al5 and

Non-Magnetic YFe7Al5 Studied by X-ray and Neutron Diffraction”, J. Alloys Compd.,

303-304, 270-275 (2000) (Crys. Structure, Experimental, Magn. Prop., 7)

[2000Sik] Sikora, W., Schobinger-Papamantellos, P., Buschow, K.H.J., “Symmetry Analysis of the

Magnetic Ordering in RFe4Al8 (R = La, Ce, Y, Lu and Tb) Compounds (II)”, J. Magn.

Magn. Mater., 213, 143-156 (2000) (Calculation, Crys. Structure, Magn. Prop., 8)

[2000Wae] Waerenborgh, J.C., Salamakha, P., Sologub, O., Goncalves, A.P., Cardoso, C., Serio, S.,

Godinho, M., Almeida, M., “Influence of Thermal Treatment and Crystal Growth on the

Final Composition and Magnetic Properties of the YFexAl12-x (4 x 4.2) Intermetallics”,

Chem. Mater., 12, 1743-1749 (2000) (Crys. Structure, Experimental, Magn. Prop., 17)

[2001Ike] Ikeda, O., Ohnuma, I., Kainuma, R., Ishida, K., “Phase Equilibria and Stability of Ordered

BCC Phases in the Fe-Rich Portion of hte Fe-Al System”, Intermetallics, 9, 755-761 (2001)

(Equi. Diagram, Thermodyn., Experimental, 18)

[2001Kny] Knyazev, Yu.V., Kuchin, A.G., Kuz'min, Yu.I., “Optical Conductivity and Magnetic

Parameters of the Intermetallic Compounds R2Fe17-xMx (R = Y, Ce, Lu; M = Al, Si)”, J.

Alloys Compd., 327, 34-38 (2001) (Crys. Structure, Experimental, Magn. Prop., Optical

Prop., 23)

[2001Pai] Paixao, J.A., Silva, M.R., Waerenborgh, J.C., Concalves, A.P., Lander, G.H., Brown, P.J.,

Godinho, M., Burlet, P., “Magnetic Structures of MFe4+ Al8- (M = Lu, Y)”, Phys. Rev. B,

63B(5), 054410-1 - 054410-12 (2001) (Crys. Structure, Experimental, Magn. Prop., 29)

[2001Vor] Voronin, V.I., Berger, I.F., Kuchin, A.G., Sheptyakov, D.V., Balagurov, A.M., “Real

Disordered Crystal Structure and Curie Temperature of Intermetallic Compounds

Y2Fe17-xMx (M = Si or Al)”, J. Alloys Compd., 315, 82-89 (2001) (Crys. Structure,

Experimental, Magn. Prop., 17)

[2001Wae1] Waerenborgh, J.C., Salamakha, P., Sologub, O., Goncalves, A.P., Serio, S., Godinho, M.,

Almeida, M., “Fe Moessbauer Spectroscopy Study of the AFexAl12-x Intermetallics (A = Y,

Tm, Lu and U, 4 x 4.3)”, J. Alloys Compd., 318, 44-51 (2001) (Crys. Structure,

Experimental, Moessbauer, 21)

[2001Wae2] Waerenborgh, J.C., Salamakha, P., Sologub, O., Serio, S., Godinho, M., Goncalves, A.P.,

Almeida, M., “Y-Fe-Al Ternary System: Partial Isothermal Section at 1070 K Powder X-

Ray Diffraction and Moessbauer Spectroscopy Study”, J. Alloys Compd., 323-324, 78-82

(2001) (Crys. Structure, Experimental, Moessbauer, 9)

[2003Cor] Cornish, L., Cacciamani, G., Saltykov, P., “Al-Y (Aluminium-Yttrium)”, MSIT Binary

Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science

International Services GmbH, Stuttgart; Product ID: 20.14305.1.20, (2003) (Crys.

Structure, Equi. Diagram, Assessment, 23)

[2003Kal] Kalvius, G.M., Wagner, F.E., Noakes, D.R., Schreier, E., Waeppling, R., Zimmermann, U.,

Schaefer, W., Kockelmann, W., Halevy, I., Gal, J., “Magnetic Behavior of YFexAl12-x”,

Physica B, 326B(1-4), 460-464 (2003) (Experimental, Magn. Prop., Moessbauer, 7)



Stuttgart; to be published, (2003) (Equi. Diagram, Crys. Structure, Assessment, 58)

16


MSIT®

Al–Fe–Y


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

( Al) hP2

P63/mmc

Mg

a = 269.3

c = 439.8

at 25°C, 20.5 GPa [Mas2]

( Al)

< 660.452

cF4

Fm3mCu

a = 404.96 at 25°C [Mas2]

( Fe) hP2

P63/mmc

Mg

a = 246.8

c = 396.0

at 25°C, 13 GPa [Mas2]

( Fe)

1538 - 1394

cI2

Im3mW

a = 293.15 [Mas2]

( Fe)

1394 - 912

cF4

Fm3mCu

a = 364.67 at 915°C [V-C2, Mas2, 1993Kat]

dissolves up to 1.2 at.% Al

( Fe)

< 912

cI2

Im3mW

a = 286.65

a = 286.64 to 289.59

a = 286.60 to 289.99

a = 286.60 to 290.12

pure Fe at 25°C [Mas2]

dissolves up to 45.0 at.% Al at

1310°C

0 - 18.8 at.%Al, HT [1958Tay]

0 - 19.0 at.% Al, HT [1961Lih]

0 - 18.7 at.% Al, 25°C [1999Dub]

( Y)

1522 - 1478

cI2

Im3mW

a = 407 [Mas2]

( Y)

< 1478

hP2

P63/mmc

Mg

a = 364.82

c = 573.18

at 25°C [Mas2]

Fe4Al13

< 1160

mC102

C2/m

Fe4Al13

a = 1552.7 to 1548.7

b = 803.5 to 808.4

c = 1244.9 to 1248.8

= 107.7 to 107.99°

a = 1549.2

b = 807.8

c = 1247.1

= 107.69°

74.16 - 76.70 at.% Al [2003Pis]

also called FeAl3 in the literature

at 76.0 at.% Al [1994Gri]

Fe2Al5< 1169

oC24

Cmcm

-

a = 765.59

b = 641.54

c = 421.84

at 71.5 at.% Al [1994Bur]

17


MSIT®

Al–Fe–Y

FeAl2< 1156

aP18

P1

FeAl2

a = 487.8

b = 646.1

c = 880.0

= 91.75°

= 73.27°

= 96.89°

at 66.9 at.% Al [1993Kat]

1102 - 1232

cI16?

-

-

a = 598.0 at 61 at.% Al [1993Kat]

FeAl

< 1310

cP8

Pm3m

CsCl

a = 289.48 to 290.5

a = 289.53 to 290.9

a = 289.81 to 291.01

a = 289.76 to 190.78

34.5 - 47.5 at.% Al [1961Lih]

36.2 - 50.0 at.% Al [1958Tay]

39.7 - 50.9 at.% Al [1997Kog]

500°C quenched in water

room temperature

Fe3Al

< 547

cF16

Fm3mBiF3

a = 579.30 to 578.86

a = 579.30 to 578.92

~24 - ~37 at.% Al [2001Ike]

23.1 - 35.0 at.% Al [1958Tay]

24.7 - 31.7 at.% Al [1961Lih]

Fe2Al9 mP22

P21/c

Co2Al9

a = 869

b = 635

c = 632

= 93.4°

metastable

81.8 at.% Al [1993Kat]

FeAl6 oC28

Cmc21

FeAl6

a = 744.0

b = 646.3

c = 877.0

a = 744

b = 649

c = 879

metastable

85.7 at.% Al [1993Kat]

[1998Ali]

FeAl4+x t** a = 884

c = 2160

(0 < x < 0.4) metastable

[1998Ali]

YAl3980 - 654(?)

hR36

R3m

BaPb3

a = 620.4 0.2

c = 2118.4 0.7

[V-C2]

YAl3< 645(?)

hP8

P63/mmc

Ni3Sn

a = 627.6 0.2

c = 458.2 0.1

[V-C2]

Metastable phase?

Y(FexAl1-x)2

YAl2 < 1485

cF24

Fd3m

MgCu2

a = 783.4 to 768.9

a = 785.5 0.7

a = 778 to 786

0 x 0.41 [1975Dwi]

x = 0 - 0.25, T = 800°C [2001Wae2]

x = 0 [1989Gsc]

x = 0 [1994Fol]

YAl

< 1130

oC8

Cmcm

CrB

a = 388.4 0.2

b = 1152.2 0.4

c = 438.5 0.2

[V-C2]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

18


MSIT®

Al–Fe–Y

Y3Al2< 1100

tP20

P42/mnm

Zr3Al2

a = 823.9 0.3

c = 764.8 0.4

[V-C2]

Y5Al3 hP16

P63/mcm

Mn5Si3

a = 878.7

c = 643.5

Metastable [1987Ric] from

recrystallized rapidly quenched alloys

Y2Al

< 985

oP12

Pnma

Co2Si

a = 664.2 2

b = 508.4 1

c = 946.9 2

[V-C2]

Y(Fe1-xAlx)2

YFe2

< 1125

cF24

Fd3m

MgCu2

a = 735.5 to 751.0

a = 736.3

0 x 0.33 [1975Dwi]

at x = 0 - 0.30 annealed at 1000°C

[1977Mur]

at x = 0 [V-C2]

YFe3

1350

hR36

R3m

PuNi3

a = 513.7

c = 2461

[V-C2]

Y6(Fe1-xAlx)23

Y6Fe23

1300

cF116

Fm3m

Th6Mn23

a = 1208.4

a = 1208.4

at x = 0.09, refined at 250°C

[1985Gan]

at x = 0 [V-C2]

Y2(Fe1-xAlx)17(HT)

Y2Fe17(HT)

? < T < 1400

hP38

P63/mmc

Th2Ni17 a = 850.1 to 856.6

c = 831.2 to 833.7

a = 852.13 to 852.61

c = 832.86 to 833.44

a = 846.3

c = 828.2

0 x 0.24 (Th2Zn17 at x > 0.24)

[1998Kam]

at x = 0.06 - 0.18, annealed at 950°C

[1986Plu]

at x = 0 - 0.1, annealed at 1300°C

X-ray and neutron diffr. [2001Vor]

at x = 0.0 [V-C2]

Y2(Fe1-xAlx)17(RT)

Y2Fe17(RT)

hR19

R3m

Th2Zn17

a = 874.6 to 880.0

c = 1266.6 to 1274.9

a = 874.46

c = 1267.28

a = 860.4 to 872.4

c = 1256.8 to 1264.7

a = 846.0

c = 1241.0

0 x 0.45 at 500°C [1972Zar]

at x = 0.45 - 0.56, as cast [1976McN]

at Y2Fe9Al8, T = 10 K [1998Che]

at x = 0.23 - 0.41, annealed at 950°C

[1986Plu]

at x = 0 [V-C2]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

19


MSIT®

Al–Fe–Y

* 1, Y(Fe1-xAlx)2

YFeAl

hP12

P63/mmc

MgZn2

a = 536.5 to 540.2

c = 873.9 to 877.5

a = 541

c = 861

a = 534.1

c = 880.5

a = 541

c = 881

0.35 x 0.54 [1975Dwi]

at x = 0.40 - 0.50 [1975Bus]

at x = 0.5 [1973Zar]

at x = 0.4, annealed at 1000°C

[1977Mur]

at x = 0.33 [1972Ryk]

* 2, Y(FexAl1-x)12

YFe4Al8

YFe6Al6

tI26

I4/mmm

ThMn12

a = 872

c = 504

a = 872.2

c = 503.6

a = 874.0

c = 504.5

a = 864.6

c = 499.2

a = 873.2

c = 501.8

a = 871.2

c = 503.6

a = 871.6

c = 502.4

a = 869.83

c = 504.30

a = 868.7

c = 503.2

a = 882.6 to 871.6

c = 506.3 to 503.2

a = 864.67 to 876.04

c = 503.74 to 505.04

a = 861.7 to 862.9

c = 503.1 to 504.0

0.257 x 0.58


[1966Zar]

at x = 0.33 [1974Viv]

at x = 0.33 [1976Bus]

at x = 0.5 [1980Fel]

at x = 0.5 [1988Che]


[1988Che]

at x = 0.5, T = 210 K neutron

diffraction [1998Sch]

at x = 0.42 [1995Sch]

at x = 0.42 neutron diffraction

[1995Sch]

at x = 0.257 - 0.382 single crystal

[2000Sal]

at x = 0.33 - 0.46, T = 800°C

[2001Wae2]

at x = 0.58, T = 20 - 127°C [2000Sch]

* 3, YFe2Al10 oP52

Cmcm

YbFe2Al10

a = 896.9

b = 1015.6

c = 901.8

a = 896.49

b = 1015.68

c = 901.13

[1998Thi]

at T = 800°C [2001Wae2]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

20


MSIT®

Al–Fe–Y

20

40

60

80

20 40 60 80

20

40

60

80

Y Fe


Axes: at.%

YAl3

YAl2

YFe2

YFe3Y

6Fe

23

Y2Fe

17

Fe3Al

FeAl2

Fe2Al

5

Fe4Al

13τ3

τ2

τ1

FeAl

(αFe)

(αAl)Fig. 1: Al-Fe-Y.

Isothermal section at

500°C

21


MSIT®

Al–Fe–Zn

Aluminium – Iron – Zinc

Gautam Ghosh

Literature Data

Constitutional equilibria in the Al-Fe-Zn system is very important for the production of high quality

Zn-coatings in steels by a process commonly known as hot-dip galvanizing. As a result, a large number of

experimental studies have been carried out to determine the phase equilibria. The earlier results [1922Fue,

1934Fue, 1945May, 1947May, 1953Geb, 1953Ray, 1961Ren] on the phase equilibria were reviewed

several times [1943Mon, 1952Han, 1961Phi, 1969Wat, 1976Mon]. [1953Ray] studied the solidification

using about 150 ternary alloys, and also reported isothermal sections at 350 and 370°C. [1961Ren]

investigated the phase equilibria in alloys containing up to 20 mass% Al and 20 mass% Fe. They reported

isothermal sections at 600, 400°C and at room temperature. The most comprehensive study was carried out

by [1970Koe] and [1971Koe]. They investigated a large number of alloys containing up to 60 mass%

(Fe+Zn). The alloys were prepared using Armco-grade Fe and 99.99 mass% Al and Zn. The ternary alloys

were prepared by adding either Fe or Zn to a master alloy of Fe:Al 50:50 or to pure Al. The solidification

path and the isothermal sections were determined by means of thermal analysis, X-ray diffraction and

microstructural investigations. They presented a reaction scheme, liquidus surface, nine isothermal sections

in the temperature range of 250 to 700°C, and four temperature-composition sections. [1973Ure1]

investigated the partial isothermal section at 450°C by means of metallography and electron microprobe

analysis. They carried out equilibration experiments using solid Al-Fe intermetallic (FeAl, FeAl2, Fe2Al5,

or Fe4Al13) and either liquid Zn or Zn-1.71Al (mass%) alloy. Prepared samples in evacuated capsules were

held at 450°C for 800 h followed by quenching in iced water. These results were critically assessed by

[1992Gho] and [1992Rag].

Recently, there has been a renewed interest in the phase equilibria, particularly the Zn corner around 450°C,

due to very stringent quality control requirements of galvanized steel sheets for the automotive industry. As

a result, recent studies are focused primarily in experimental determination [1990Che, 1992Per, 1994Tan,

1995Tan2, 1996Tan, 1997Gyu, 1997Uwa1, 1999Tan] and CALPHAD modeling [1991Bel, 1992Per,

1999Cos, 2001Gio, 2002Bai] of phase equilibria of the Zn corner in the temperature range of 450 to 470°C.

Due to rapid interfacial reaction between steel and liquid Al-Zn alloys, the importance of metastable

equilibria [1991Bel, 1992Per, 2002Bai], diffusion path [1992Per, 1998Ada, 1998Uch1, 1998Uch2,

2002Bai], and the mechanism of phase transformations [1994Lin, 1995Lin1, 1995Lin2, 1995Tan1,

1995Yam2, 1997Mcd, 1997Mor, 1997Ser, 1998Ada, 1998Uch1, 1998Uch2, 1998Yam, 2002Bai] during

interfacial reaction have also been elucidated.

[1990Che] prepared three ternary alloys using Al, Fe and Zn powders of unspecified purity. The final heat

treatment of the alloys was annealing at 450°C for about 10 h. The phase equilibria were determined by

XRD and SEM/EDX techniques. [1991Bel] determined the stable and metastable solubility limits of Fe in

liquid (Zn) 447 to 480°C. [1992Per] determined the metastable and stable isothermal sections at 450°C

based on the interfacial reaction studies between solid Al-Fe and liquid Al-Zn alloys. They used Al-Fe

alloys containing 5, 29 and 36 at.% Al, and liquid Al-Zn alloys containing 0.12, 0.22, 0.39 and 11.2 at.%

Al. Both short time (less than 30 min) and long time (1000 h) experiments were carried out. The phase

compositions were determined by SEM/EDX technique. [1994Tan] reported an isothermal section of Zn

corner at 470°C. Tang [1995Tan2, 1996Tan] reported the phase equilibria at 450°C by combining the

results of [1990Che] and his experimental data of the Zn-corner. [1997Uwa1] prepared four ternary alloys

by dry ball milling. They used elemental powders of following purity: 99.5% Al, 99.9+% Fe and 99.9% Zn.

The ball milled powders were annealed at 300, 400 and 570°C for 3 h. They used DSC to study phase

transformations, and XRD to identify the phases. Some of the controversial results of [1997Uwa1] have

been the subject of extensive discussions [1997Tan, 1997Uwa2, 1998Tan, 1998Uwa]. [2000Tan]

determined the Fe solubilities in dilute liquid Al-Zn alloys in the temperature range of 450 to 480°C. He

prepared 16 ternary alloys containing up to 0.1 mass% Fe and up to 0.23 mass% Al using 99.5% pure Fe

22


MSIT®

Al–Fe–Zn

and Al, and special high grade Zn. The final equilibrations of encapsulated samples were carried out at 450,

465 and 485°C for 40 h followed by water quenching. The phase equilibria information were extracted from

SEM/EDX analysis. [2002Tan] re-investigated the phase equilibria of the Zn corner at 435°C using six

ternary alloys. They were annealed at 450°C for 15 days, and composition of phases were determined by

SEM-EDS analysis. [2002Bai] reported a calculated isothermal section at 450°C. These recent results have

been reviewed by [2003Rag].

Binary Systems

The Al-Fe, Al-Zn and Fe-Zn binary phase diagrams are accepted from [2003Pis], [2003Per] and [1982Kub],

respectively.

There are some differences between the presently accepted binary phase diagrams and those accepted by

the previous investigators [1953Ray, 1970Koe, 1971Koe]. For example, [1970Koe] and [1971Koe]

accepted an Al-Fe phase diagram in which all the order-disorder transitions involving ( Fe), 1 and 2

phases were considered to be first order, whereas in this assessment, ( Fe) 2 and 1 2 reactions have

been considered to be second order [1982Kub] reflected by the absence of the corresponding two-phase

fields. Furthermore, the Al-Fe phase diagram has undergone slight modification due to recently established

congruent melting behavior of the Fe4Al13 phase [1986Len].

In the case of the Fe-Zn phase diagram, [1953Ray, 1970Koe] and [1971Koe] considered the phase to be

stable between 672 and 620°C and the 1 phase to be stable below 640°C [1953Ray, 1970Koe, 1971Koe,

1973Ure1]. However, according to [1982Kub] the phase (which is the 1 phase as designated by the above

authors) is stable below 665°C. It is worth mentioning that [1970Koe] and [1971Koe] convincingly

established the phase at temperatures above the 1 phase field near the Zn corner, but later on [1973Ure1]

failed to identify the phase above the 1 phase field. Also, according to [1982Kub], the and phases

react to form the 1 phase at 550°C. This feature was also absent in the Fe-Zn phase diagram accepted by

the previous studies [1953Ray, 1970Koe, 1971Koe, 1973Ure1]. Very recent study of solid-state equilibria

of Zn rich alloys [2001Mit], and thermodynamic modeling of phase equilibria [2000Reu, 2001Su] are

consistent with the Fe-Zn phase diagram assessed by [1982Kub].

In the Al-Zn phase diagram, the phase designated by [1953Ray, 1970Koe, 1971Koe, 1973Ure1] is

identical to (Al) in the phase diagram given by [1983Mur]. All these features are taken into account in this

critical assessment of phase equilibria.

Solid Phases

Available data suggest that the solubility of Zn in (Fe,Al) is a function of time of heat treatment at 450°C,

with less Zn after shorter time compared to longer time. For example, [1990Che] gives 2 mass% Zn after

10 h at 450°C, while [1992Per] gives 2.26 mass% Zn after less than 30 min at 450°C and 4.85 mass% Zn

after 1000 h at 450°C.

The equilibrium solubility of Zn in Fe4Al13 at 450°C are 7 mass% [1973Ure1], 5.5 mass% [1990Che], 7.61

mass% [1992Per], while under metastable equilibrium Fe4Al13 can dissolve up to 13.92 mass% [1992Per]

and 15.2 mass% [1997Gyu]. [1953Ray] noted that the X-ray diffraction pattern of Zn containing Fe4Al13 is

slightly different from that of pure Fe4Al13 which might be due to the slight structural alteration caused by

the non-random occupation of the Zn atoms. [1992Per] reported that the presence of Zn in FeAl2 is hardly

detectable.

The solubility of Zn in Fe2Al5 ( ) has been determined several times by reacting Fe with liquid Al-Zn bath

containing varying amounts of Al [1971Ghu, 1973Har, 1973Ure1, 1973Ure2, 1984Nit, 1990Che, 1991Sai,

1992Per, 1997Gyu]. Available data fall in the range of 11 to 23 mass% Zn, and also show a systematic trend

that the Zn-content in Fe2Al5 ( ) is a function of reaction time. Due to rapid interdiffusion, the data after

short time reaction show higher solubility of Zn in Fe2Al5 compared to long time experiments. For example,

[1992Per] found 22.87 mass% Zn in Fe2Al5 after reaction at 450°C for less than 30 min compared to 18.7

mass% Zn after reaction at 450°C for 1000 h. [1971Ghu] noted a scatter of 14 to 17 mass% Zn in Fe2Al5after reaction at 600°C for 10s. It is important to note that while short time reaction data is relevant to

industrial galvanizing process, long time data is appropriate to construct the equilibrium phase diagram.

23


MSIT®

Al–Fe–Zn

Accordingly, we have accepted the solubility of 18.7 mass% Zn (11 at.%) at 450°C [1992Per] as

equilibrium value. X-ray diffraction and density measurement show that Zn atoms reside on the Fe site for

up to 6.7 at.% Zn giving the formula Fe4Zn10Al, and beyond this composition Zn atoms also reside on the

Al sites giving the formula Fe4Zn9Zn2 [2001Koe].

[1973Ure1] reported a solid solubility of 3.6 mass% Al in the phase (FeZn10) at 450°C, which is in

qualitative agreement with that of [1956Hor]. On the other hand, [1990Che] and [1992Per] reported solid

solubilities of 2.8, 3.71, and 1.84 mass% Al at 450°C. Since the latter value was obtained after long time

(1000 h) heat treatment, it is considered as equilibrium solid solubility while other values correspond to

metastable equilibria. The phase (FeZn13) dissolves 0.78 mass% Al at 450°C [1992Per], but [1961Ren]

gives a much lower value of 0.2 mass%. The solid solubilities of Al in and 1 phase at 450°C are similar

to that in phase [1992Per]. On the other hand, Tang’s [1996Tan] isothermal section at 450°C show much

higher solubility of Al in these two phases which may correspond to industrial galvanizing conditions.

[1992Per] reported two Phases, 1 (denoted as 2 by [1992Per]) and 2 (denoted as 3 by [1992Per]),

after equilibration for 1000 h at 450°C. However, [1973Ure1] did not detect any 2 after 800 h equilibration

at 450°C. On the hand, [1995Yam2] reported continuous solid solubility ( 1) and [1996Tan] reported

continuous solid solubility ( ´) in the isothermal sections at 440 and 450°C, respectively. It is possible that

these conditions are realized during galvanizing process, and may not represent equilibrium. Later,

[1998Yam] synthesized single phase alloys corresponding to 2 and 3 compositions of [1992Per], and

diffusion annealing (conditions are not specified) of mechanically pressed 2 and 3 did not show any

evidence of continuous solid solubility. Even though the crystallographic data of 2 is lacking, available

results suggest that it may be a ternary phase.

The details of the crystal structures and lattice parameters of the solid phases are listed in Table 1.


Based on the results of [1970Koe] and [1971Koe], the reaction scheme is summarized in Fig. 1. A number

of changes have been made to comply with the binary phase diagrams accepted here. The reaction scheme

proposed by [1970Koe] contained fourteen invariant reactions. However, three invariant reactions proposed

to occur at 485, 440 and 320°C [1970Koe, 1971Koe] are not considered in Fig. 1 as they are not compatible

with the presently accepted binary phase diagrams. The assessed reaction scheme is consistent with all the

phase diagram information available until now. [1961Ren] proposed a ternary U type invariant reaction

L+FeAl2 +Fe2Al5 at 592°C; however, subsequent detailed investigations by [1970Koe, 1971Koe] and

[1973Ure1] failed to detect this reaction.

Liquidus Surface

Figure 2 shows the liquidus surface from 20 to 70 mass% Al and 0 to 40 mass% Zn and Fig. 3 shows the

liquidus surface of the Zn corner, both according to [1970Koe] and [1971Koe]. Results of solidification

studies of Zn rich ternary alloys by [1945May, 1947May] and [1962May] and of Al/Zn rich alloys

[1953Geb] agree quantitatively with those of [1970Koe] and [1971Koe].

Isothermal Sections

Figures 4, 5 and 6 show the isothermal sections at 700, 575 and 500°C, respectively, after [1970Koe] and

[1971Koe]. Figure 7 shows the isothermal section of the Zn corner at 500°C [1970Koe, 1971Koe]. Figures

8 and 9 show partial isothermal section at 470 [1994Tan] and 460°C [2000Tan], respectively, depicting the

solubility limits of Fe in liquid-Zn with respect to (FeZn13), (FeZn10), and (Fe2Al5) phases.

The isothermal section at 450°C has been investigated several times. There is substantial agreement

between the earlier results of [1970Koe], [1971Koe] and [1973Ure1]. Recent significant results are due to

[1990Che, 1992Per, 1995Tan2, 1996Tan]. Except for [1992Per] and [1996Tan], others did not consider 1

phase in the 450°C isothermal section. Figure 10 shows the isothermal section at 450°C [1992Per]. Figure

11 shows the isothermal section of Zn corner depicting the phase fields involving liquid, , , 1 and 2

[1992Per]. [2002Tan] labelled as 2 phase T. Despite qualitative agreement between the results of

24


MSIT®

Al–Fe–Zn

[1992Per], [1996Tan] and [2002Tan] at 450°C, the isothermal section of [1992Per] is preferred because the

authors used much longer annealing time. Figure 12 shows the isothermal section at 450°C depicting the

saturation limits of Fe with respect to , , 2 and phase in liquid Zn [1996Tan]. Contrary to the

suggestion of [1962Cam] that the solubility of Fe should decrease with Al content in liquid Zn, [1973Ure1]

proposed that the solubility of Fe in liquid Zn at 450°C is 0.029 mass%, irrespective of the Al content. In

fact, [1991Bel] showed that when phase is in equilibrium with liquid Zn, indeed the Fe solubility

decreases with increasing Al content in liquid Zn which is seen in Figs. 8, 9 and 12. Thermodynamic

calculations also predict a similar behavior [2002Bai].

The isothermal section of the Zn-corner at 400°C [1970Koe, 1971Koe] is shown in Fig. 13. The

Fe4Al13-Al-Zn partial isothermal sections at 350, 330, 300 and 250°C are shown in Figs. 14, 15, 16, 17,

respectively according to [1970Koe] and [1971Koe]. A number of adjustments have been made in the

isothermal sections in order to comply with the binary phase diagrams.

[1961Ren] studied the isothermal sections of the Zn corner with up to about 20 mass% (Fe+Al) at 600°C,

450°C and room temperature. At 600°C, [1961Ren] observed three-phase fields L+ +FeAl2 and

L+Fe2Al5+FeAl2, and proposed a ternary U type invariant reaction L+FeAl2 +Fe2Al5 at 592°C.

However, more detailed investigations by [1970Koe, 1971Koe] and [1973Ure1] failed to observe these

features. The partial isothermal section at 450°C given by [1961Ren] agrees qualitatively with that of

[1973Ure1], but the exact locations of the phase boundaries differ significantly. Because of these reasons,

the results of [1961Ren] are not accepted here.


Figures 18, 19, 20 and 21 show isopleths at 30, 90, 95 and 98 mass% Zn, respectively [1970Koe, 1971Koe].

In Fig. 18, several changes have been made to comply with the accepted Al-Zn phase diagram.

Thermodynamics

[1995Yam1] reported the activity coefficient of Al in liquid Al-Zn alloys containing up to 10 mass% Zn,

and in liquid Al-Fe-Zn alloys containing up to 1 mass% Al at 480°C. [1995Yam2] determined the chemical

potential of Al in liquid Zn, in equilibrium with Fe4Al13, (Fe2Al5), 2, (FeZn10), and (FeZn13) in the

temperature range of 432 to 510°C. [1971Ghu] reported that the heat formation of Fe(Al,Zn)3 is much more

negative compared to the heat of formation of Fe4Al13 and Fe2Al5 phases; however, the actual values

reported by [1971Ghu] are very doubtful.

[2000Tan] reported that the solubility product of Fe2Al5 in liquid Zn can be expressed as

ln(mass% Al)5(mass% Fe)2 = 28.1 - 33066/T

where T is the temperature in Kelvin. Besides, [2000Tan] has also discussed a procedure to calculate the

solubility limits of Fe in liquid Zn with respect to saturation of , and phase. [1991Bel] reported

solubility products of Fe4Al13, Fe2Al5, FeAl2, FeAl and FeZn13 in liquid Zn. Using the experimental

solubility data, [2001Gio] has derived the Gibbs energy of formation of Fe2Al5Znx ( ). [2002Feu] measured

the standard enthalpy of formation of phase at Fe0.07Zn0.93 using solution calorimetry technique.

Several attempts have been made to calculate phase diagrams by CALPHAD method [1991Bel, 1992Per,

1999Cos, 2001Gio, 2002Bai]. Of particular interest is the prediction of solubility of Fe and Al in liquid Zn

around 450°C, and also the diffusion path during hot-dip galvanizing process. [1991Bel] calculated

metastable solubilities in liquid-Zn with respect to +Fe2Al5, +FeAl, +Fe4Al13 and +FeAl2 saturations

at 447 and 477°C, and did not consider the phase. On the other hand [1992Per] calculated the solubility

of Fe in liquid-Zn at 450°C considering all binary phases, and found a slightly higher solubility of Fe in

liquid-Zn compared to [1991Bel] due to participation of the phase.

[1999Cos] calculated the 465°C isothermal section, but only the Zn-corner to understand the limiting factor

controlling solubility of Fe in liquid Zn. They did not consider any ternary interaction parameter in the

liquid phase and also the ternary solubility of Fe-Zn intermetallics. Nonetheless, the calculated activity

coefficients of Al in liquid Zn-0.01 mass% Fe-xAl alloys are in good agreement with the experimental data

of Yamaguchi et al. [1995Yam1, 1995Yam2]. Even though their calculated solubility limit of Fe2Al5 is in

25


MSIT®

Al–Fe–Zn

good agreement with experiment, the calculated three phase equilibrium L+Fe2Al5+Fe4Al13 differs

significantly from the experimental data [1995Yam2, 1998Yam].

[2002Bai] calculated the entire isothermal section at 450°C, and it appears that they overestimated the solid

solubility of Al in , and 1 phase compared to the experimental data of [1992Per]. Also, they did not

consider the 2 phase. Nonetheless, their calculation clearly shows a decrease in solubility of Fe in liquid

Zn when it is in equilibrium with the phase (Fe2Al5).

Miscellaneous

The solubility of Fe in a liquid Zn-4Al (mass%) alloy, in the temperature range of 400 to 675°C, was

determined by [1963Fri]. The solubility can be expressed as

log(mass% Fe) = 3.6359 - 5149/T

log(at.% Fe) = 3.6825 - 5150/T

where T is the temperature in K.

Additions of Al to a liquid Zn bath inhibit the reaction between solid Fe and liquid Zn during the normal

galvanizing process. It is believed that Al causes the formation of an inhibition layer, consisting of Fe2Al5,

at the substrate/coating interface [1995Tan1]. However, detailed experiments using TEM/SEM/XRD

techniques clearly show that the inhibition layer actually consists of Fe2Al5 and Fe4Al13. The details of the

reactions and the formation sequences of the different binary intermetallic phases during the hot dip

galvanizing process have been reported by [1965Sou, 1971Ghu, 1973Har, 1973Ure2, 1975Gut, 1984Nit,

1991Sag, 1995Lin1, 1995Lin2, 1995Tan1, 1997Mcd, 1997Ser, 1998Uch1, 1998Uch2]. Addition of Si also

suppresses the rapid exothermic reaction between liquid Al-Zn and Fe by forming a solid reaction layer

[1989Sel] which acts as a diffusion barrier. A comprehensive review of physical metallurgy of the

galvanizing process has been presented by Marder [2000Mar].

[1998Akd] proposed that the value of activity coefficient of Al in (Fe,Al,Zn) alloys has a strong influence

on the formation and growth kinetics of interfacial diffusion layer. Besides, [2002Bai] compiled the

diffusion data in , , and 1 phases which were then used to model the mobility of components in these

phases within CALPHAD formalism.

[1977Sho] investigated the effect of pressure on the reaction kinetics between solid Fe and liquid Zn-1.5Al

(mass%) at 501°C. An applied pressure was found to cause the intermetallic compounds to become unstable

and change the overall reaction rate from linear to non-linear. The stability of phase, compared to other

phases, under pressure is markedly affected by the presence of the Al in the melt.

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26


MSIT®

Al–Fe–Zn

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27


MSIT®

Al–Fe–Zn

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[1996Tan] Tang, N.-Y., “450°C Isotherm of Zn-Fe-Al Phase Diagram Update”, J. Phase Equilib.,

17(5), 396-398 (1996) (Equi. Diagram, Experimental, #, *,13)

[1997Gyu] Gyurov, S., “The Reaction Between Solid Iron and Liquid Zn-Al Baths”, Z. Metallkd.,

88(4), 346-352 (1997) (Equi. Diagram, Experimental, Kinetics, 33)

28


MSIT®

Al–Fe–Zn

[1997Mcd] McDevitt E., Morimoto Y., Meshii M., “Characterization of the Fe-Al Interfacial Layer in

a Commercial Hot-Dip Galvanized Coating”, ISIJ Int., 37(8), 776-782 (1997)

(Experimental, 24)

[1997Mor] Morimoto Y., McDevitt E., Meshii M., “Characterization of the Fe-Al Inhibition Layer

Formed in the Initial Stages of Hot-Dip Galvannealing”, ISIJ Int., 37(9), 906-913 (1997)

(Experimental, 28)

[1997Ser] Sere, P.R., Culcasi, J.D., Elsner, C.J, Di Sarli, A.R., “Factors Affecting the Hot-dip Zinc

Coatings Structure” (in Spanish), Rev. de Metall., 33(6), 376-381 (1997) (Experimental,

Kinetics, 11)

[1997Tan] Tang, N.-Y., “Discussion of “Kinetics and Phase Transformation Evaluation of Fe-Zn-Al

Mechanically Alloyed Phases”, Metall. Mater. Trans. A, 28A(11), 2433-2434 (1997)

(Theory, 11)

[1997Uwa1] Uwakwen, O.N.C., Liu, Z., “Kinetics and Phase Transformation Evaluation of Fe-Zn-Al

Mechanically Alloyed Phases”, Metall. Mater. Trans. A, 28A(3), 517-525 (1997) (Equi.


[1997Uwa2] Uwakwen, O.N.C., Liu, Z., “Authors’ Reply”, Metall. Mater. Trans. A, 28A(11), 2434-2435

(1997) (Theory, 7)

[1998Ada] Adachi Y., Arai M., “Transformation of Fe-Al Phase to Fe-Zn Phase on Pure Iron During

Galvanizing Layer”, Mater. Sci. Eng. A, 254(1-2), 305-310 (1998) (Crys. Structure,

Experimental, 8)

[1998Akd] Akdeniz, M.V., Mekhrabon, A.O., “The Effect of Substitutional Impurities on the Evolution

of Fe-Al Diffusion Layer”, Acta Mater., 46(4), 1185-1192 (1998) (Calculation,

Thermodyn., 55)

[1998Tan] Tang, N.-Y., “Discussion of “Kinetics and Phase Transformation Evaluation of Fe-Zn-Al

Mechanically Alloyed Phases”, Metall. Mater. Trans. A, 29A(10), 2643-2644 (1998) (Equi.

Diagram, Theory, 9)

[1998Uch1] Uchida Y., Andoh A., Komatsu A., Yamakawa K., “Changing Process from Center Dot

Fe-Zn Phase to Al-Fe Intermetallic Compounds in Molten Zn-5mass%Al Alloy Bath” (in

Japanese), Tetsu to Hagane, 84(9), 632-636 (1998) (Experimental, 6)

[1998Uch2] Uchida Y., Andoh A., Komatsu A., Yamakawa K., “Changing Process from Center Dot

Fe-Zn Phase to Al-Fe Intermetallic Compounds in Molten Zn-5mass%Al Alloy Bath” (in

Japanese), Tetsu to Hagane, 84(9), 637-642 (1998) (Experimental, 4)

[1998Uwa] Uwakwen, O.N.C., Liu, Z., “Authors’ Reply”, Metall. Mater. Trans. A, 29A(10), 2644-2645

(1998) (Equi. Diagram, Theory, 5)

[1998Yam] Yamaguchi, S., “Thermochemical Stability and Precipitation Behavior of Dross Phases in

CGL Bath” in Proc. Galvatech’98, Chiba, Japan, The Iron and Steel Institute of Japan,

84-88 (1998) (Experimental, Thermodyn., *, 8)

[1999Cos] Costa e Silva, A., Avillez, R.R., Marques, K., “A Preliminary Assessment of the Zn-rich

Corner of the Al-Fe-Zn System and Its Implications in Steel Coating”, Z. Metallkd., 90(1),

38-43 (1999) (Calculation, Equi. Diagram, Thermodyn., *, 25)

[1999Tan] Tang, N.-Y., “Characteristics of Continuous-Galvanizing Baths”, Metall. Mater. Trans. B.,

30(1), 144-148 (1999) (Equi. Diagram, *, 26)

[2000Mar] Marder, A.R., “The Metallurgy of Zinc-Coated Steel”, Prog. Mater. Sci., 45, 191-271

(2000) (Equi. Diagram, Phys. Prop., Review, 188)

[2000Reu] Reumont, G., Perrot, P., Fiorani, J.M., Hertz, J., “Thermodynamic Assessment of the Fe-Zn

System”, J. Phase Equilib., 21(4), 371-378 (2000) (Thermodyn., *, 26)

[2000Tan] Tang, N.-Y., “Determination of Liquid-Phase Boundaries in Zn-Fe-Mx Systems”, J. Phase

Equilib., 21(1), 70-77 (2000) (Equi. Diagram, Experimental, Thermodyn., #, *, 29)

[2001Gio] Giorgi, M.-L., Guillot, J.-B., Nicolle, R., “Assessment of the Zinc-Aluminium-Iron Phase

Diagrams in the Zinc-Rich Corner”, Calphad, 25(3), 461-474 (2001) (Equi. Diagram,

Thermodyn., *, 36)

29


MSIT®

Al–Fe–Zn

[2001Koe] Koester, M., Schuhmacher, B., Sommer, D., “The Influence of the Zinc Content on the

Lattice Constants and Structure of the Intermetallic Compound Fe2Al5”, Steel Res., 72(9),

371-375 (2001) (Crys. Structure, Experimental, 29)

[2001Mit] Mita, K., Ikeda, T., Maeda, M., “Phase Diagram Study of Fe-Zn Intermetallics”, J. Phase

Equilib., 22(2), 122-125 (2001) (Experiment, Equi. Diagram, #, *, 14)

[2001Su] Su, X., Tang, N.-Y., Toguri, J.M., “Thermodynamic Evaluation of the Fe-Zn System”, J.

Alloys Compd., 325(9), 129-136 (2001) (Thermodyn., *, 49)

[2002Bai] Bai, K., Wu, P., “Assessment of the Zn-Fe-Al System for Kinetic Study of Galvanizing”, J.

Alloys Compd., 347, 156-164 (2002) (Equi. Diagram, Thermodyn., Kinetics, *, 40)

[2002Feu] Feutelais, Y., Legendre, B., de Avillez, R. R., “Standard Enthalpy of Formation of the

-Phase in the Fe-Zn System at 298 K”, J. Alloys Compd., 346, 1-2 (2002) (Experimental,

Thermodyn., Kinetics, *, 20)

[2002Tan] Tang, N.Y., Su, P., “Assessment of the Zn-Fe-Al System for Kinetic Study of Galvanizing”,

J. Alloys Comp., 347, 156-164 (2002) (Equi. Diagram, Experimental, #, *, 16)

[2003Per] Perrot, P., “Al-Zn (Aluminium-Zinc)”, MSIT Binary Evaluation Program, in MSIT


Stuttgart; to be published, (2003) (Equi. Diagram, Assessment, Crys. Structure, 41)



Stuttgart; to be published, (2003) (Equi. Diagram, Assessment, Crys. Structure, 58)

[2003Rag] Raghavan, V., ”Al-Fe-Zn (Aluminum-Iron-Zinc)”, J. Phase Equilib., 24, 546-550 (2003)

(Equi. Diagram, Review, *, 33)


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Al) cF4

Fm3m

Cu

a = 404.88

a = 403.52

a = 403.29

a = 403.14

pure Al at 24°C [V-C]

at 63.0 at.% Zn and 360°C [1983Mur]

at 64.8 at.% Zn and 360°C [1983Mur]

at 70.1 at.% Zn and 360°C [1983Mur]

( Fe) cI2

Im3m

W

a = 286.65 pure Fe at 20°C [V-C]

(Zn) hP2

P63/mmc

Mg

a = 266.46

c = 494.61

pure Zn at 22°C [V-C]

1, Fe3Al

552.5

cF16

Fm3m

BiF3

a = 578.86 to 579.3 [2003Pis], solid solubility

ranges from 22.5 to 36.5 at.% Al

2, FeAl

1310

cP2

Pm3m

CsCl

a = 289.76 to 290.78 [2003Pis], at room temperature



, Fe2Al31102 - 1232

cI16? a = 598.0 [2003Pis], solid solubility ranges


30


MSIT®

Al–Fe–Zn

FeAl2 1156

aP18

P1

FeAl2

a = 487.8

b = 646.1

c = 880.0

= 91.75°

= 73.27°

= 96.89°

[2003Pis], at 66.9 at.% Al



, Fe2Al5 1169

oC24

Cmcm

a = 765.59

b = 641.54

c = 421.84

a = 764.14

b = 642.76

c = 421.87

a = 762.23

b = 646.25

c = 423.00

[2003Pis], at 71.5 at.% Al


from 71.0 to 72.5 at.% Al.

Equilibrium solubility is up to 11 at.%

Zn at 450°C [1992Per].

[2001Koe], at Fe4Al10Zn

[2001Koe], at Fe4Al9Zn2

Fe4Al13

1160

mC102

C2/m

Fe4Al13

a = 1552.7 to 1548.7

b = 803.5 to 808.4

c = 1244.9 to 1248.8

= 107.7 to 107.99°

a = 1549.2

b = 807.8

c = 1247.1

= 107.69

[2003Pis], 74.16 to 76.7 at.% Al



[2003Pis], at 76.0 at.% Al

sometimes called FeAl3 in the literature

, Fe3Zn10

782

cI52

I43m

Fe3Zn10 ?

Cu5Zn8

a = 897.41

a = 901.8

[V-C], solid solubility ranges

from 68.0 to 82.5 at.% Zn

1, Fe11Zn39

550

cF408

F43m

Fe11Zn39

a = 1796.3 [V-C2], solid solubility ranges

from 75.5 to 81.0 at.% Zn

, FeZn10

665

hP555

P63mc

FeZn10

a = 1283.0

b = 5770.0

[V-C], solid solubility ranges from

86.5 to 92.0 at.% Zn.

Equilibrium solubility is up to 4.3 at.%

Al at 450°C [1992Per].

, FeZn13

530

mC28

C2/m

CoZn13

a = 1342.4

b = 760.8

c = 506.1

= 127.3°

[V-C], solid solubility ranges from

92.5 to 94.0 at.% Zn.

Equilibrium solubility is up to 1.85 at.%

Al at 450°C [1992Per].

2, AlFe14Zn1.5

450 (?)

- - [1992Per, 1998Yam]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

31


MSIT®

Al–Fe–Zn

Fig

. 1:

A

l-F

e-Z

n.

Rea

ctio

n s

chem

e

Al-

Fe

Fe-

Zn

A-B

-C

l +

α2

ε1

232

p1

l +

(αF

e)Γ

78

2p

3

Lδ

+ (Z

n)

ca.4

25

max

Al-

Fe-

Zn

L +

α2

(αF

e) +

εca

.12

00

U1

Al-

Zn

l (

Al)

+ (

Zn)

38

1e 5

lε

+ η

11

65

e 1

lη

+ F

e 4A

l 13

11

60

e 2

ε +

η F

eAl 2

11

56

p2

εα 2

+F

eAl 2

11

02

e 3

l +

Γδ

66

5p

4

55

0p

5

Γ +

δΓ 1

l +

δ

ζ5

30

p6

l +

ζ (

Zn)

42

5p

7

l (

Al)

+ F

e 4A

l 13

66

5e 4

(Al´

) (

Al´

´) +

(Z

n)

27

7e 6

L +

ε (

αFe)

+ η

11

30

U2

ε +

η(α

Fe)

+ F

eAl 2

10

65

U3

ε(α

Fe)

+ α

2 +

FeA

l 21

038

E1

L +

Γ(α

Fe)

+ δ

ca.6

60

U4

L +

(αF

e)δ

+ η

55

3U

5

Lδ

+ (

Zn

) +

ζca

.42

0E

2

L +

δη

+ (

Zn

)4

18

U6

L +

η F

e 4A

l 13 +

(Z

n)

40

9U

7

L F

e 4A

l 13+

(Al)

+(Z

n)

37

9E

3

(Al´

)(A

l´´)

+F

e 4A

l 13+

(Zn

)2

74

E4

L+

(αF

e)+

α 2

?

(αF

e)+

α 2+

ε(α

Fe)

+L

+ε

η+F

e 4A

l 13+

(Zn

)

L+

(αF

e)+

η

(αF

e)+

η+ε

(αF

e)+

FeA

l 2+

η(α

Fe)

+F

eAl 2

+ε

(αF

e)+

α 2+

FeA

l 2

Γ+(α

Fe)

+δ

L+

(αF

e)+

δ

(αF

e)+

δ+η

?

L+

δ+η

L+

η+(Z

n)

δ+(Z

n)+

ζ

L+

Fe 4

Al 1

3+

(Zn

)

Fe 4

Al 1

3+

(Al)

+(Z

n)

Fe 4

Al 1

3+

(Al´

)+(A

l´´)

ca.3

51

(Zn

)+F

e 4A

l 13+

(Al)

η+δ+

(Zn

)

32


MSIT®

Al–Fe–Zn

30

40

50

10 20 30

50

60

70

Fe 60.00Zn 0.00Al 40.00

Fe 20.00Zn 40.00Al 40.00

Fe 20.00Zn 0.00Al 80.00 Data / Grid: at.%

Axes: at.%

e2

e1

p1

U1

U2

Fe4Al

13

η

(αFe)

α2

ε

?

Fig. 2: Al-Fe-Zn.

Partial liquidus

surface

Fe 10.00Zn 90.00Al 0.00

Zn


Axes: at.%

Γ δ

ζ

(Zn)

(αFe)

η

Fe4 A

l13

U7

U5

U6

E2

U4

p4

from p3

from U2

from e2

to E3

p6

p7

Fig. 3: Al-Fe-Zn.

Liquidus surface of

the Zn corner

33


MSIT®

Al–Fe–Zn

20

40

60

80

20 40 60 80

20

40

60

80

Fe Zn


Axes: at.%

(αFe)

Γ

L

Fe4Al

13

ηFeAl

2

α2

L+Γ+(αFe)

L+(αFe)+η

L+Fe4Al

13+η

(αFe)+FeAl2+η

20

40

60

80

20 40 60 80

20

40

60

80

Fe Zn


Axes: at.%

(αFe)

Γ δ

L

(Al)

Fe4Al

13

ηFeAl

2

α2

(αFe)+Γ+δ L+δ+(αFe)

L+(αFe)+η

L+Fe4Al

13+η

L+(Al)+Fe4Al

13

(αFe)+FeAl2+η

(αFe)+α2

Fig. 4: Al-Fe-Zn.


700°C

Fig. 5: Al-Fe-Zn.


575°C

34


MSIT®

Al–Fe–Zn

10

20

80 90

10

20

Fe 30.00Zn 70.00Al 0.00

Zn


Axes: at.%

Γ1

δ

ζ

L

L+Fe4Al

13+η

L+δ+η

L+ζ+δ

δ+η+(αFe)

(αFe)+Γ+δ

Γ+Γ1+δ

20

40

60

80

20 40 60 80

20

40

60

80

Fe Zn


Axes: at.%

(αFe)

Γ Γ1

δ

ζ

L

(Al)

Fe4Al

13

ηFeAl

2

α2α

1

L+(Al)+Fe4Al

13

L+Fe4Al

13+η

L+δ+Fe+η

(αFe)+δ+η

(αFe)+FeAl2+η

(αFe)+α1

Γ+δ+(αFe)

Fig. 7: Al-Fe-Zn.

Partial isothermal

section at 500°C

Fig. 6: Al-Fe-Zn.


500°C

35


MSIT®

Al–Fe–Zn

Fe 0.55Zn 99.45Al 0.00

Zn


Axes: at.%

L

L+η

L+ζ

L+δ

L+ζ+δ

L+δ+η

Fe 0.55Zn 99.45Al 0.00

Zn


Axes: at.%

L

L+η

L+η+δ

L+ζ

L+δL+δ+ζ

Fig. 8: Al-Fe-Zn.

Partial isothermal

section at 470°C

Fig. 9: Al-Fe-Zn.

Partial isothermal

section at 460°C

36


MSIT®

Al–Fe–Zn

20

40

60

80

20 40 60 80

20

40

60

80

Fe Zn


Axes: at.%

(αFe)

Γ Γ1

δζ

L

(Al)

Fe4Al

13

FeAl2

α2

α1

(αFe)+δ

Γ2

L+Fe4 A

l13

L+Fe4 Al

13 +(Al)

(αFe)+δ+η

(Al)+Fe4Al

13

η

(αFe)+Γ

10

90

10

Fe 20.00Zn 80.00Al 0.00

Zn


Axes: at.%

Γ1

δζ

L

η+δ

(αFe)+δ

L+ζ

η+L

L+Fe4Al

13

δ+L

Γ2

η+Γ2

Γ1+δ

η+Γ2+L

Γ2+L

Fig. 10: Al-Fe-Zn.


450°C

Fig. 11: Al-Fe-Zn.

Partial isothermal

section at 450°C

37


MSIT®

Al–Fe–Zn

Fe 0.55Zn 99.45Al 0.00

Zn


Axes: at.%

L

L+η

L+ζ+δ

L+δ+Γ2

L+Γ2+η

L+Γ2L+δ

L+ζ

10

20

30

70 80 90

10

20

30

Fe 40.00Zn 60.00Al 0.00

Zn


Axes: at.%

Γ1

δ

ζ

(Zn)

L

L+(Al)

(Al)

L+Fe4Al

13+(Al)

L+Fe4Al

13+(Zn)

(Zn)+η+Fe

4 Al13

(Zn)+η+δ

(αFe)+η+δ

(αFe)+Γ+δ

Γ+Γ1+δ

(Zn)+δ+ζ

Fig. 12: Al-Fe-Zn.

Partial isothermal

section at 450°C

Fig. 13: Al-Fe-Zn.

Partial isothermal

section at 400°C

38


MSIT®

Al–Fe–Zn

20

40

60

80

20 40 60 80

20

40

60

80

Fe Zn


Axes: at.%

(Zn)

(Al')

(Al")

Fe4Al

13

(Al')+(Al")

(Al')+(Zn)+Fe4Al

13

TK

(Al')+(Al")+Fe4Al

13

20

40

60

80

20 40 60 80

20

40

60

80

Fe Zn


Axes: at.%

(Zn)

(Al')

(Al")

Fe4Al

13

(Al')+(Al")+Fe4Al

13

(Al')+(Zn)+Fe4Al

13

Fig. 15: Al-Fe-Zn.

Partial isothermal

section at 330°C

Fig. 14: Al-Fe-Zn.

Partial isothermal

section at 350°C

39


MSIT®

Al–Fe–Zn

20

40

60

80

20 40 60 80

20

40

60

80

Fe Zn


Axes: at.%

(Zn)

(Al')

(Al")

Fe4Al

13

(Al')+(Zn)+Fe4Al

13

(Al')+(Al")+Fe4Al

13

20

40

60

80

20 40 60 80

20

40

60

80

Fe Zn


Axes: at.%

(Zn)

(Al)

Fe4Al

13

(Al)+(Zn)+Fe4Al

13

Fig. 16: Al-Fe-Zn.

Partial isothermal

section at 300°C

Fig. 17: Al-Fe-Zn.

Partial isothermal

section at 250°C

40


MSIT®

Al–Fe–Zn

60 70 80

200

300

400

500

600

700

800

900

1000

Fe 21.67Zn 18.51Al 59.82

Fe 0.00Zn 15.03Al 84.97Al, at.%

Te

mp

era

ture

, °C

379°C409°C

274°C

L+(Al)

(Al)

(Al')+(Zn)

Fe4Al13+L

L

η+F

e 4A

l 13+L

η+F

e4A

l 13+

(Zn)

Fe

4Al 13

+(Z

n)

Fe 4Al13+(Al')+(Zn)

+(Zn)

Fe4Al13+(Al)

Fe4Al13+(Al)+L

Fe4Al13+(Al)

Fe4Al13+(Al')+(Al'')

10 20

200

300

400

500

600

700

800

900

1000

Zn 88.49Fe 11.51Al 0.00

Zn 78.78Fe 0.00Al 21.22Al, at.%

Te

mp

era

ture

, °C L+η

L

409°

379°C

274°Cη+(Z

n)

418°C

553°C

δ+(Zn)

δ δ+ζ

660°C

L+Γ

L+δ

L+(Al)

(Al)+(Zn)

(Al")+(Zn)

420

L+(αFe)

L+η+(αFe) L+Fe4Al13

Fe4Al13+(Al'')+(Zn)

Fe4Al3+(Zn)

Fe4Al13+(Al)+(Al'')

η+F

e 4Al 1

3+(Z

n)

η+δ+(Zn)

L+η+δ

L+(αFe)+δ

L+Γ+(αFe)

δ+ζ+(Zn)

L+δ+ζ

L+(Al)+(Zn)

(Al)+(Al'')+(Zn)

L+Fe4Al13+(Zn)

L+η+Fe4Al13

Fig. 18: Al-Fe-Zn.

Section at a constant

Zn-content

of 30 mass%

Fig. 19: Al-Fe-Zn.

Vertical section at a

constant Zn-content

of 90 mass%

41


MSIT®

Al–Fe–Zn

10

200

300

400

500

600

700

800

900

1000

Zn 94.20Fe 5.80Al 0.00

Zn 88.69Fe 0.00Al 11.31Al, at.%

Te

mp

era

ture

, °C

L

L+Fe4Al13

L+η

553°C

418°C420°C

L+Γ

660°C

409°

379°C

274°C

δ+(Zn)

ζ+(Zn)

L+(αFe)

L+(Zn)

(Al)+(Zn)

(Al")+(Zn)

δ+η+(Zn)

Fe4Al13+(Zn)

L+(αFe)+η

L+δ+η

L+(αFe)+Γ

L+(αFe)+δ

L+δ

L+ζ

L+Γ+δ

L+ζ+δ

δ+ζ+(Zn)

L+η+Fe4Al13

η+(Zn)

η+Fe4Al13+(Zn)

L+Fe4Al13+(Zn)

L+δ+ζ

Fe4Al13+(Al'')+(Zn)

L+η+(Zn)

L+(Zn)+(Al)

(Al)+(Al'')+(Zn)

Fe4Al13+(Al)+(Zn)

L+(Zn)+δ

200

300

400

500

600

700

800

Zn 97.67Fe 2.33Al 0.00

Zn 95.29Fe 0.00Al 4.71Al, at.%

Te

mp

era

ture

, °C

L

L+δ

L+(αFe)

553°C

418°C420°C

L+ζ

ζ+(Zn)ζ+δ+(Zn)

409°C 379°C

274°C

L+(Zn)

(Al)+(Zn)

η+(Zn)

(Al")+(Zn)

Fe

4Al 13

+(Z

n)

L+η

L+η+(αFe)

L+η+δ

Fe4Al13+(Al'')η+F

e 4A

l 13+(Z

n)

η+(Zn)+δ

+(Zn)

δ+(Zn)

L+ζ+δ

Fe4Al13+(Al)

+(Zn)

L+(Al)+(Zn)

(Al)+(Al'')+(Zn)

L+(αFe)+δ

L+ζ+(Zn)L+η+(Zn)

Fig. 20: Al-Fe-Zn.


constant Zn-content

of 95 mass%

Fig. 21: Al-Fe-Zn.


constant Zn-content

of 98 mass%

42


MSIT®

Al–Fe–Zr

Aluminium – Iron – Zirconium

Zoya M. Alekseeva, updated by Viktor Kuznetsov

Literature Data

[1966Mar] investigated alloys along the section ZrAl2-ZrFe2 by X-ray diffraction; the alloys studied were

prepared by arc-melting and annealed at 900°C for 20 d. Two ternary Laves type compounds 1 and 2 were

found with extended homogeneity regions along the section studied.

[1968Gru] investigated, essentially by metallography, alloys in the Zr corner along the sections with Al to

Fe ratios 2:1 and 1:2 up to 14 mass% Al and 14 mass% Fe, using alloys that were quenched from 1350,

1200, 1100, 900, 800 and 700°C. Partial isothermal sections at 1200, 900, 800 and 700°C were constructed.

However, in the isothermal sections below 1200°C the existence of the ternary compound Zr6FeAl2reported by [1969Bur] has not been taken into account. Isothermal sections at 900 and 800°C do not contain

the binary compound Zr2Fe and the binary compound Zr3Fe is missing in the isothermal sections at 800 and

700°C.

[1969Bur] investigated, mainly by X-ray diffraction, 116 alloys which were prepared by arc melting and

annealed at 900°C for 2100 h. Two more ternary compounds have been found in addition to 1 and 2

reported earlier: (i) a "line" compound ZrFe7-4Al5-8 with Al content varying from 37 to 61 at.% and (ii) a

stoichiometric compound Zr6FeAl2. An isothermal section at 900°C has been constructed.

[1970Kri] established the crystal structure of Zr6FeAl2 compound; the structure was later that refined by

[1997Yan].

[1973Ath] investigated (by EMPA, X-ray and electron diffraction) the ternary compound ZrFe3.3Al1.3

occurring in a two-phase alloy (the other phase was Fe3Al) which was prepared by substituting 5 at.% Zr

for Fe in the alloy Fe76Al24. The alloy studied was annealed at 950°C for 24 h.

[1974Dwi] investigated the ternary equiatomic compound ZrFeAl which was prepared by arc melting and

annealed in a Vycor capsule.

[1974Kuz] prepared alloys from the elements with a purity of 99.99% and annealed them at 500°C for 50

days. By the measurement of the lattice parameters they determined the existing phases and their solubility

ranges on the section ZrAl2-ZrFe2. These results are in agreement with [1966Mar].

[1977Mur] studied the crystal structure, magnetic properties and Fe Moessbauer effect on the Laves phase

Zr(Fe1-xAlx)2 in the stoichiometric range x = 0 to 0.4.

[1987Bla] investigated the solubility of Al in ZrFe2 by means of X-ray powder diffraction and

measurements of microhardness on alloys melted and heat treated for at least 24 h in the range 800 to

1500°C. The samples were either quenched or cooled at 1.7 K·min-1. The substitution of Al for Zr changed

the unit cell parameter from 706.8 pm for ZrFe2 to 702.3 pm for (Zr0.87Al0.13)Fe2.

A brief review of the system mainly concerning intermetallics formation may be found in [1990Kum].

[1991Des] found no evidence for the presence of L12 phase in mechanically alloyed sample with

composition of Al-12.5Fe-25Zr (at.%).

[1991Sok] studied partial section from Al corner with Zr to Fe ratio being 1:3, isopleth of 25 at.% Zr and

partial isothermal section at 500°C for Al < 25 at.%. They used Al 99.9% purity, iodide-purified Zr 99.9%

and Fe 99.9%. Alloys were prepared in arc furnace with water-cooled Cu bottom in Ti gettered Ar

atmosphere with subsequent annealing in evacuated silica tubes at 500°C for 1000 h and water quenching.

Samples were studied by DTA, metallography and X-ray analysis. No ternary phases were found in the

region studied.

[1992Sle] investigated temperature dependence of lattice spacing at 0 to 300 K and magnetic susceptibility

at 80 to 600 K of the Laves phase with composition of ZrFe1.2Al0.8.

43


MSIT®

Al–Fe–Zr

[1993Nov], [1994Isr] and [1997Isr] studied bonding characteristics in Laves phases Zr(AlxFe1–x)2 with

various x values. Experimental techniques included Moessbauer spectroscopy,

nuclear-resonant-photon-scattering and neutron diffraction; all were used to determine effective Debye

temperature which measured bonding strength. Minimal value of that at x = 0.2 was found to coincide with

maximal hydrogen absorption power.

[1994Kle] measured standard enthalpies of formation calorimetrically for Zr(Fe(1-x)Alx phase at x = 0,

0.0833, 0.2, 0.5, 0.7 and 1 by measuring heat of dissolution in acid mixture (HF+HNO3). [1996Gon] used

these data (among much others) to test their generalization of well-known Miedema model to ternary

intermetallics with moderate success.

[1999Zav] investigated structural changes of Zr6FeAl2 under hydrogen treatment.

[1999Mek] performed ab initio calculation of interatomic potentials and influence of Zr additions on the

ordering in intermetallics of the Al-Fe system.

[2000Biz] studied in great detail kinetics of crystallization of Al-Fe and Al-Fe-Zr rapidly solidified alloys.

In particular, a number of kinetic models were tried.

Mechanical alloying of sample with Zr3Fe7Al90 composition was studied by [2001Rod] who found a

mixture of amorphous and unspecified nanocrystalline phases and studied their crystallization behavior

using DSC and X-ray techniques.

This evaluation incorporates and continues the critical evaluation made by [1992Ale] considering new

published data.

Binary Systems

For the Al-Fe and Al-Zr binary systems recently updated versions of [2003Pis] and [2003Sch] were

accepted, respectively. Fe-Zr system is from [Mas2].

Solid Phases

Five ternary compounds have been found in the system. ZrFe2 extends into the ternary to about 10 at.% Al.

The existence of an additional ternary phase with AuCu3 structure was claimed by [1989Sch] at the

composition Zr25Fe5.5Al and 1100°C; the temperature and composition range of existence is still unknown,

so it could not be included in the phase diagram.

Crystallographic data of all the phases are listed in Table 1.


The partial vertical section from Al corner with Zr to Fe ratio of 1:3 [1991Sok] crosses a plane of invariant

reaction between L, (Al), Fe4Al13 and Al3Zr phases at about 650°C (see below Fig. 4), but neither its nature,

nor phase compositions are provided (the temperature value was taken by present author from small-scale

figure).

Isothermal Sections

The partial isothermal section at 1200°C, presented in Fig. 1, is based on the results of [1968Gru]. To bring

that into agreement with accepted version of Fe-Zr binary, the boundary ( Zr)+L/L was shifted; also some

modification of position of L corner of ( Zr)+L+Zr5Al3 tie-triangle was necessary. These changes

necessitate certain boundaries given in the original work as uncertain.

Figure 2 displays the isothermal section at 900°C based on the results of [1969Bur]. In both isothermal

sections the phase of [1989Sch] is not included since [1968Gru, 1969Bur] did not detect this phase. To

adapt to the accepted binary systems, changes were made as following: the three-phase field

Zr+ Zr+Zr3Al was inserted; a liquid single-phase field in the Al corner and the corresponding two- and

three-phase fields were added. The ternary compound Zr18Fe59Al23 [1973Ath] was also included with the

corresponding three-phase fields. Extension of the 2 phase field is shown according to the stoichiometry

reported in [1966Mar, 1969Bur, 1974Kuz]. It should be noted that in the isothermal section reported by

[1969Bur] it has been shown up to 60 at.% Al, which however, contradicts the tabulated results of

44


MSIT®

Al–Fe–Zr

[1966Mar], where a sample with 60 at.% Al contained a second phase, ZrAl2. Some minor shifts of position

of tie-lines of equilibria with that phase, which do not contradict to real phase compositions of the alloys

studied, had to be done.

Figure 3 presents the partial isothermal section at 500°C in the Al rich corner [1991Sok].


Figure 4 displays partial vertical section from Al corner with Zr to Fe ratio of 1:3 [1991Sok]. Figure 5

presents the isopleth at 75 at.% Al, taken from the same source. According to the accepted Al-Fe binary, the

L/L+Fe4Al13 boundary line (given in [1991Sok] as dashed line) must approach the temperature axis a bit

higher than point c, but this may hardly be seen in the scale of original figure.

Thermodynamics

[1994Kle] measured standard enthalpies of reactions: 2x Al + 2(1-x)Fe + Zr = Zr(AlxFe(1-x))2 using acid

-solution calorimetry at 25°C. The results are: for x = 0 H = –71 8 kJ, for x = 0.0833 H = –74 9 kJ, for

x = 0.2 H = –83 10 kJ, for x = 0.5 H = –125 13 kJ, for x = 0.7 H = –232 21 kJ, and for x = 1

H = –154 13 kJ.

Theoretical results of [1996Gon] are not in very good agreement with these.

Miscellaneous

[1988Vig] compared the microstructural stability of Al-8Fe and Al-8Fe-1.5Zr (mass%) alloys. The ribbons

used were produced by melt spinning and were about 40 to 60 m thick and 4 to 5 mm wide. Fine ZrAl3precipitates appear in the Al matrix during ageing at 200 to 400°C along with FeAl6.

The substitution of Al for Fe rapidly reduces the Fe magnetic moment of the compound ZrFe2 [1977Mur]

and the substitution of Al for Zr reduces microhardness values of the compound from 8329 to 6818 N·mm-2

[1987Bla].

[1991Sik] studied possible techniques of industrial treatment of Fe3Al intermetallic, including that with Zr

additions.

[1999Mek] performed ab initio calculation of influence of a number of elements (including Zr) on ordering

in FeAl compound. It has been shown that Zr atoms substitute preferentially for Fe sublattice sites in FeAl

compound.

References

[1961Now] Nowotny, H., Schob, O., Benesovsky, F., “The Crystal Structure of Zr2Al and Hf2Al” (in

German), Monatsh. Chem., 92, 1300-1304 (1961) (Crys. Structure, Experimental, 10)

[1966Mar] Markiv, V.Ya., Kripyakevich, P.I., “Compounds of the Type R(X', X'')2 in Systems with

R = Ti, Zr, Hf; X' = Fe, Co, Ni, Cu and X'' = Al, Ga and Their Crystal Structures”, Sov.

Phys.-Crystallogr. (Engl. Transl.), 11, 733-738 (1967), translated from Kristallografiya, 11,


[1968Gru] Gruzdeva, N.M., Zagorskaya, T.N., Raevskii, I.I., “Structure and Properties of Alloys in the

Zirconium Corner of Al-Fe-Zr System” (in Russian), in: Fiziko-Khimiya Splavov

Tsirkoniya (Physical Chemistry of Zirconium Alloys), Moscow: Nauka, 5-9 (1968) (Equi.

Diagram, Experimental, #, 3)

[1969Bur] Burnashova, V.V., Markiv, V.Ya., “Study of Al-Fe-Zr System”, Dopov. Akad. Nauk Ukr.

RSR, A, (4), 351-353 (1969) (Crys. Structure, Equi. Diagram, Experimental, *, 16)

[1970Kri] Kripyakevich, P.I., Burnashova, V.V., Markiv, V.Ya., “Crystal Structure of the Compounds

Zr6FeAl2, Zr6CoAl2, and Zr6NiAl2”, Dopov. Akad. Nauk Ukr. RSR A, (9), 828-831 (1970)

(Crys. Structure, Experimental)

[1973Ath] Athanassiadis, G., Dirand, M., Rimlinger, L., “X-Ray Diffraction and Electron Diffraction

Study of the Compound of Al1.3Fe3.5Zr” (in French), C. R. Seances Acad. Sci. (Paris), 277,

C915-C917 (1973) (Crys. Structure, Experimental, 3)

45


MSIT®

Al–Fe–Zr

[1974Dwi] Dwight, A.E., “Alloying Behavior of Zr, Hf and the Actinides in Several Series of

Isostructural Compounds”, J. Less-Common Met., 34, 279-284 (1974) (Crys. Structure,

Experimental, 6)

[1974Kuz] Kuz’menko, P.P., Suprunenko, P.A., Markiv, V.Ya., Butsik, T.M., “Magnetic Properties of

Laves Phases in the Zr-Fe-Al and Zr-Co-Al Systems” (in Russian), Akad. Nauk Ukr. SSR,

Metallofizika, 52, 58-61 (1974) (Crys. Structure, Equi. Diagram, Experimental, 10)

[1977Mur] Muraoka, Y., Shigas, M., Nakamura, Y., “Magnetic Properties and Mössbauer Effect of

A(Fe1-xBx)2 (A =Y or Zr, B = Al or Ni) Laves Phase Intermetallic Compounds”, Phys.

Status Solidi, 42A, 369-374 (1977) (Crys. Structure, Experimental, 15)

[1987Bla] Blarzina, Z., Trojko, R., “On Friauf-Laves Phases in the Zr1-xAlxT2, Zr1-xSixT2 and

Zr1-xTixT2 (T = Mn, Fe, Co) Systems”, J. Less Common Met., 133, 277-286 (1987) (Crys.


[1988Vig] Vigier, E., Ortez-Mendez, U., Merles, P., Thaller, G., Fouguet, F., “Microstructural

Stability of Rapidly Quenched Al, Fe Alloys: Influence of Zirconium”, Mater. Sci. Eng., 98,

191-195 (1988) (Experimental, 11)

[1989Ale] Alekseeva, Z.M., Korotkova, N.V., “Phase Diagram of the Fe-Zr System” (in Russian), Izv.

Akad. Nauk SSSR, Met., (4), 202-208 (1989) (Crys. Structure, Equi. Diagram,

Experimental, #, 21)

[1989Sch] Schneibel, J.H., Porter, W.D., “High Temperature Order Intermetallic Alloys III”, Mater.

Res. Soc. Symp. Proc., Stoloff, N.S. (Ed.), 335-340 (1989) (Crys. Structure)

[1990Kum] Kumar, K.S., “Ternary Intermetallics in Aluminium-Refractory Metal-X Systems (X = V,

Cr, Mn, Fe, Co, Ni, Cu, Zn)”, Int. Mat. Rev., 35, 293-327 (1990) (Crys. Structure, Equi.


[1991Des] Desch, P.B., Schwarz, R.B., Nash, P., “Formation of Metastable L12 Phases in Al3Zr and

Al-12.5% X-25% Zr (X = Li, Cr, Fe, Ni, Cu)”, J. Less-Common Met., 168, 69-80 (1991)


[1991Sik] Sikka, V.K., “Production of Fe3Al-Based Intermetallic Alloys”, Mater. Res. Soc. Symp.

Proc.: High-Temp. Ordered Intermetallic Alloys IV, 213, 907-912 (1991) (Experimental, 2)

[1991Sok] Sokolovskaya, E.M., Kazakova E.F., Grigorovitch E.V., Matveyev I.N., “Phase Equilibria

in Alloys of the Al-Fe-Zr System”(in Russian), Vestn. Mosk. Univ., Ser. 2: Khim., 32,

478-481 (1991) (Equi. Diagram, Experimental, *, #, 7)

[1992Ale] Alekseeva, Z.M., “Aluminium - Iron - Zirconium”, MSIT Ternary Evaluation Program, in

MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services

GmbH, Stuttgart; Document ID: 10.16088.1.20, (1992) (Crys. Structure, Equi. Diagram,

Assessment, 15)

[1992Sle] Slebarski, A., Hafez, M., Zarek, W., “Spin Fluctuations in ZrM1.2Al0.8 with Transition

Metal M of the 3d Type”, Solid State Commun., 82(1), 59-61 (1992) (Crys. Structure,

Experimental, 12)

[1993Nov] Novik I., Yacob B., March R., “Moessbauer Study of Crystallographic and Magnetic Phase

Transitions, Phonon Softening, and Hyperfine Interactions in Zr(AlxFe1–x)2”, Phys. Rev. B,

47, 723-726 (1993) (Phys. Prop., Experimental)

[1994Isr] Israel A., Yacob I., March R., Shanal O., Wolf A., Fogel M., “Correlation Between

Anomalous Hydrogen Absorption and 56Fe-Bonding Strength in the Zr(AlxFe1-x)2 System”,

Phys. Rev. B, 50, 3564-3569 (1994) (Phys. Prop., Experimental, 29)

[1994Kle] Klein, R., Jacob, I., O'Hare, P.A.G., Goldberg, R.N., “Solution-Calorymetric Determination

of the Standard Molar Enthalpies of Formation of the Pseudobinary Compounds

Zr(AlxFe(1-x))2 at the Temperature 298.15 K”, J. Chem. Thermodyn., 26, 599-608 (1994)

(Thermodyn., Experimental, 22)

[1996Gon] Goncalves, A.P., Almeida, M, “Extended Miedema Model: Predicting the Formation

Enthalpies of Intermetallic Phases with More than Two Elements”, Physica B (Amsterdam),

228, 289-294 (1996) (Thermodyn., Theory, 19)

46


MSIT®

Al–Fe–Zr

[1997Yan] Yanson, T.I., Manyako, M.B., Bodak, O.I., Cerny, R., Pacheko, J.V., Yvon, K., “Crystal

Structure of Zirconium Iron Aluminide, Zr6FeAl2”, Z. Kristallogr. NCS, 212, 504 (1997)


[1997Isr] Israel, A., Jacob, I., Soubeyroux, J.L., Fruchart, D., Pinto, H., Melamud, M., “Neutron

Diffraction Study of Atomic Bonding Properties in the Hydrogen-Absorbing Zr(AlxFe1-x)2

System”, J. Alloys Compd., 253-254, 265-267 (1997) (Phys. Prop., Experimental, 12)

[1999Mek] Mekhrabov, A.O., Akdeniz, M.V., “Effect of Ternary Alloying Elements Addition on

Atomic Ordering Characteristics of Fe-Al Intermetallics”, Acta Mater., 47, 2067-2075

(1999) (Thermodyn., Theory, 63)

[1999Zav] Zavaliy, I.Yu., Pecharsky, V.K., Miller, G.J., Akselrud, L.G., “Hydrogenation of Zr6MeX2

Intermetallic Compounds (Me=Fe, Co, Ni, X=Al, Ga, Sn): Crystallographic and Theoretical

Analysis”, J. Alloys Compd., 283, 106-116 (1999) (Crys. Structure, Experimental, 31)

[2000Biz] Bizjak, M., Kosec, L., “Phase Transformations of Al-Fe and Al-Fe-Zr Rapidly Solidified

Alloys”, Z. Metallkd., 91, 160-164 (2000) (Kinetics, Electr. Prop., Experimental, 12)

[2001Rod] Rodriguez, C.A.D., Botta F., W.J., “High-Energy Ball Milling of Al-Based Alloys”, Key

Eng. Mater., 189-191, 573-578 (2001) (Crys. Structure, Experimental, 10)



Stuttgart, to be published, (2003) (Equi. Diagram, Assessment, Crys. Structure, 58)

[2003Sch] Schuster, J.C., “Al-Zr (Aluminium-Zirconium)”, MSIT Binary Evaluation Program, in


GmbH, Stuttgart; submitted for publication (2003) (Crys. Structure, Equi. Diagram,

Assessment, 151)


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Al)

< 660.452

cF4

Fm3m

Cu

a = 404.88 pure Al [V-C]

( Fe)

1538 - 1394

cI2

Im3m

W

a = 293.15 [Mas2]

( Fe)

1394 - 912

cF4

Fm3m

Cu

a = 364.67 at 915°C [V-C2, Mas2]

( Fe)

< 912

cI2

Im3m

W

a = 286.65 pure Fe at 20°C [V-C]

( Zr)(h)

1855 - 863

cI2

Im3m

W

a = 362 [P]

( Zr)(r)

< 863

hP2

P63/mmc

Mg

a = 323.2

c = 514.7

[V-C]

47


MSIT®

Al–Fe–Zr

Fe4Al13

(FeAl3.2, FeAl3)

1160

mC102

C2/m

Fe4Al13

a = 1552.7 - 1548.7

b = 803.5 - 808.4

c = 1244.9 - 1248.8

= 107.7 - 107.99°

a = 1549.2

b = 807.8

c = 1247.1

= 107.69

[2003Pis], 74.16 to 76.7 at.% Al



[2003Pis], at 76.0 at.% Al

Fe2Al5 1169

oC24

Cmcm

a = 765.59

b = 641.54

c = 421.84

[2003Pis], at 71.5 at.% Al



FeAl2< 1156

aP18

P1

FeAl2

a = 487.8

b = 646.1

c = 880.0

= 91.75°

= 73.27°

= 96.89°

[V-C]

65.5 to 67 at.% Al [Mas]

2 Fe100-xAlx< 1310

cP2

Pm3m

CsCl

a = 290.9

28.0 x 52.5 at 900°C

at x = 50 [V-C]

ZrAl3< 1580

tI16

I4/mmm

ZrAl3

a = 399.93 0.05

c = 1728.3 0.02

[2003Sch]

ZrAl2< 1660

hP12

P63/mmc

MgZn2

a = 528.24

c = 874.82

[2003Sch]

Zr2Al3< 1590

oF40

Fdd2

Zr2Al3

a = 960.1 0.2

b = 1390.6 0.2

c = 557.4 0.2

[2003Sch]

ZrAl

< 1275 25

oC8

Cmcm

CrB

a = 335.9 0.1

b = 1088.7 0.3

c = 427.4 0.1

[2003Sch]

Zr4Al3 1030

hP7

P6/mmm

Zr4Al3

a = 543.3 0.2

c = 539.0 0.2

[2003Sch]

Zr3Al2< 1480

tP20

P42/mnm

Zr4Al3

a = 763.0 0.1

c = 699.8 0.1

[2003Sch]

Zr5Al3(h)

< 1400

tI32

I4/mcm

W5Si3

a = 1104.4

c = 539.1

[2003Sch]

Zr2Al

< 1350

hP6

P63/mmc

Ni2In

a = 489.39 0.05

c = 592.83 0.05

[2003Sch]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

48


MSIT®

Al–Fe–Zr

Zr3Al

< 1019

cP4

Pm3m

Cu3Au

a = 437.2 0.3 [2003Sch]

,

(Fe1-xAlx)2-z(Zr1-yAly)1+z

cF24

Fd3m

Cu2Mg a = 706.8

a = 707.4

a = 709.4

a = 713.5

a = 712.4

a = 702.3

a = 701.0 0,3

a = 704.0 0.3

0 x 0.20, 0 y 0.133,

-0.17 z 0.03

at x = 0, y = 0, z = 0

[1987Bla]

at x = 0, y = 0, z = 0

[1977Mur]

at x = 0.1, y = 0, z = 0

[1977Mur]

at x = 0.15, y = 0, z = 0

[1966Mar]

at x = 0.2, y = 0, z = 0

[1977Mur]

at x = 0, y = 0.133, z = 0

[1987Bla]

at x = 0, y = 0, z = -0.17

[1989Ale]

at x = 0, y = 0, z = 0.03

[1989Ale]

* 1, Zr(Fe1-xAlx)2 hP12

P63/mmc

MgZn2

a = 508.7

c = 827.7

a = 524.3

c = 852.5

0.375 x 0.75 [1966Mar]

at x = 0.375 [1974Kuz]

at x = 0.75 [1974Kuz]

* 2, Zr(Fe1-xAlx)2 cF24

Fd3m

Cu2Mg

a = 743.0

a = 746.1

0.15 x 0.175

x = 0.175 [1966Mar,1974Kuz]

x = 0.15

* Zr6FeAl2 hP9

P62m

K2UF6

a = 792.1 0.2

c = 336.03 0.09

[1970Kri], [1997Yan]

* Zr18Fe59Al23 tI52

I4/mcm

a = 837

c = 998

[1973Ath]

* , Zr(Fe1-xAlx)12 tI26

I4/mmm

ThMn12

a = 859.5

c = 496.7

a = 849.3

c = 488.9

0.416 x 0.667 [1969Bur]

at x = 0.416

at x = 0.667

* , Zr25Fe5.5Al cP4

Pm3m

AuCu3

claimed by [1989Sch]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

49


MSIT®

Al–Fe–Zr

70

80

90

10 20 30

10

20

30

Zr Zr 60.00Fe 40.00Al 0.00

Zr 60.00Fe 0.00Al 40.00 Data / Grid: at.%

Axes: at.%

(βZr)

(βZr)+LL

L+Zr5Al

3

(βZr)+L+Zr5Al

3

(βZr)+Zr5Al

3

(βZr)+Zr2Al

(βZr)+Zr2Al+Zr

5Al

3

Zr5Al

3

Zr2Al

Fig. 1: Al-Fe-Zr.

Partial isothermal

section at 1200°C

20

40

60

80

20 40 60 80

20

40

60

80

Zr Fe


Axes: at.%

(αFe)

α2

FeAl2

Fe2Al5

Fe4Al13

L

ZrAl3

ZrAl2

Zr2Al3

ZrAl

Zr4Al3

Zr3Al2

Zr2Al

Zr3Al

(αZr)

(βZr) Zr2Fe ZrFe3

Zr18Fe59Al23

γ

L+ZrAl3

Zr6FeAl2

λ1

λ2

ZrFe2

Fig. 2: Al-Fe-Zr.


900°C

50


MSIT®

Al–Fe–Zr

10

20

10 20

80

90

Zr 25.00Fe 0.00Al 75.00

Zr 0.00Fe 25.00Al 75.00


Axes: at.%(Al)

(Al)+Fe4Al13+ZrAl3

Fe4Al13

ZrAl3

ZrAl3+Fe4Al13

(Al)+Fe

4 Al

13

ZrA

l 3+(

Al)

400

500

600

700

800

900

1000

1100

1200

Al Zr 6.25Fe 18.75Al 75.00Zr, at.%

Te

mp

era

ture

, °C

L

L+Fe4Al13+ZrAl3

(Al)+Fe4Al13+ZrAl3

Fe4Al13+ZrAl3

(Al)+Fe4Al13

(Al)

(Al)+L+Fe4Al13

(Al)+L

L+Fe4Al13

Fig. 3: Al-Fe-Zr.

Partial isothermal

section at 500°C

Fig. 4: Al-Fe-Zr.

Vertical section from

Al corner with

Zr/Fe=1:3 (in at.%)

51


MSIT®

Al–Fe–Zr

10 20

500

750

1000

1250

1500

Zr 0.00Fe 25.00Al 75.00

Zr 25.00Fe 0.00Al 75.00Zr, at.%

Tem

pera

ture

, °C

Fe4Al13+ZrAl3

ZrAl3

L+ZrAl3

L

1580°C

c

Fe4Al13

L+Fe4Al13+ZrAl3

L+Fe4Al13

Fig. 5: Al-Fe-Zr.

Vertical section at 75

at.% Al

52


MSIT®

Al–Ge–Li

Aluminium – Germanium – Lithium

Oksana Bodak

Literature Data

Studies on the Al-Ge-Li ternary system are confined to the identification and characterization of a few

ternary compounds. Literature data up to 1986 was reported by [1989Goe] and discussed in [1995Pav]. The

first report of a ternary phase emanated from [1952Boo] who added Li to a hypereutectic Al-Ge alloy giving

a ternary alloy with 38.30 mass% Ge, 6.22 mass% Li. Metallographic analysis clearly indicated the presence

of an unidentified phase, probably the ternary 5 compound (LiAlGe). This compound was synthesized by

[1960Now] who heated stoichiometric mixtures of the elements in an Fe crucible at temperatures between

800 and 950°C, and found that at 800°C the reaction was incomplete. At higher temperatures the compound

LiAlGe was identified, Table 1, together with a very small amount of an unidentified phase of lower

crystallographic symmetry.

By the same way [1976Sch] prepared the compound LiAlGe, heating stoichiometric amounts of the

elements in a tantalum crucible under argon for 15 min at 1000°C. The sample was subsequently annealed

for 24 h at 600°C, cooled slowly to room temperature and then the crystal structure was characterized by

neutron diffraction analysis, Table 1. The chemical analysis of the compound, 6.6Li-25.2Al-68.3Ge

(mass%), agreed well with the calculated values 6.52Li-25.33Al-68.15Ge (mass%) for the composition

LiAlGe.

A second ternary compound was identified as Li2AlGe by [1974Boc] using the same preparation technique

as [1978Ble].

[1978Ble] used 99.98 % Li, 99.999 % Al and Ge, to prepare a third ternary compound whose composition

was given as Li5.3Al0.7Ge2 with 1 formula unit in the elementary cell. This compound showed superlattice

reflections, which were ascribed to the presence of a phase of the same composition containing 3 formula

units in the elementary cell with enlargement of the “a” axis by 3. Due to the reactivity of the alloys high

temperature X-ray diffraction analysis could not be employed to determine whether Li16Al2Ge6( 1´) with

3 formula units, is a low temperature polymorph of Li5.3Al0.7Ge2 ( 1).

[1981Kis] examined three compositions on the section Li(Al1-xGex) with x = 0.02, 0.066 and 0.11. Alloys

were prepared by melting 99.999 % Al, 99.9 % Li and an Al-Ge master alloy under argon. The ingot was

encapsulated in a Pyrex glass ampoule under 0.5 atm Ar for annealing it 7 days at 500°C and then cooling

it slowly down to room temperature. Metallographically the alloys showed a eutectic structure dispersed

throughout the sample. X-ray diffraction analysis showed the presence of LiAl in the alloy with x = 0.02

and Li-rich ternary fcc-phase with a = 620 pm. It is the ternary compound 3, Table 1, with a = 616.3 pm

according to [1974Boc].

Alloys of nominal weight composition Al-2Li-0.2Ge [1986Cas] were solution heat-treated, quenched and

aged for various holding-times at 200°C. The microstructure and deformation behavior were compared for

two alloys revealing that the solubility of lithium was increased when germanium was in solid solution,

however, lithium decreased the solubility of germanium at 200°C resulting in small germanium precipitates

which were homogeneously distributed throughout the matrix. These precipitates had a very positive effect

on the deformation behavior and ductility of the alloy.

[1992Pav, 1993Pav1, 1993Pav2, 1996Dmy] constructed an isothermal section at 200°C. They prepared

their alloys in an electric arc furnace under an argon atmosphere (1.1 105 Pa) and determined the crystal

structures of the compounds. The purity of lithium was 98 mass%, the purity of silicon and aluminum was

better than 99.9 mass%. After melting all alloys were homogenized in evacuated quartz ampoules, at 200°C

for 500 h and subsequently quenched into ice water. X-ray powder analysis was used. The authors

confirmed the composition and the structure of Li5.3Al0.7Ge2 and LiAlGe compounds, determined the

crystal structure of new ternary compounds Li2AlGe and LiAl2Ge and concluded that further new

compounds Li9Al2Ge3 and Li6Al3Ge with unknown structures do exist.

53


MSIT®

Al–Ge–Li

[1994Hos] studied the effect of some ternary additions (among them Ge) in the L12 type metastable LiAl3phase calculating the heat of formation by Miedema semi-empirical formula.

Binary Systems

The description of the Al-Li phase diagram is accepted as given by [2003Gro], that of Al-Ge as given by

[Mas2]. For the system Ge-Li it is necessary to note the following. Since long there is a contradiction in

number of compounds reported in the phase diagram by [Mas2] and results of X-ray investigations on the

crystal structure of compounds. The authors [1997San] made an attempt to resolve this contradiction by

compiling the available data and constructing a hypothetic phase diagram, which subsequently was

published as a confirmed one by [2000Oka]. [1997San] however, missed the work of [1982Gru], in which

the phase diagram has been constructed in detail, using DTA and X-ray investigations. The investigations

on the Al-Ge-Li isothermal section by [1993Pav1, 1993Pav2, 1996Dmy] confirm the binary diagram given

by [1982Gru], which hence is accepted in the present evaluation and shown in Fig. 1. Remaining

discrepancies concern the composition of the Li-richest compound (Table 1) may be due to the difficult Li

refinement in the compounds during the X-ray investigation. In the present evaluation the composition

Li4Ge is accepted, as given by [1982Gru].

Solid Phases

Crystallographic data for the solid phases of this system are presented in Table 1.

Isothermal Sections

The isothermal section at 200°C shown in Fig. 2 is based on [1993Pav1, 1993Pav2, 1996Dmy]. However

the homogeneity regions of the Al-Li binary phases are adjusted to match the accepted binary diagrams.

Solubilities of a third component in the binary and unary phases were not determined by [1993Pav1,

1993Pav2, 1996Dym], and hence are not reproduced in Fig. 2 for this evaluation. The same applies for the

homogeneity ranges of the ternary phase presented by [1996Dmy].

Thermodynamics

Thermodynamic calculations of Li vapor pressures over Al-Li and Al-Li-Me, (Me=Ag, Zn, Cd, Ga, In, etc.)

are reported by [1986Lee].

References

[1952Boo] Boom, E.A., “New in the Systems Aluminium-Germanium-Sodium and

Aluminium-Germanium-Lithium” (in Russian), Dokl. Akad. Nauk SSSR, 84(4), 697-699

(1952) (Equi. Diagram, Experimental, 4)

[1960Now] Nowotny, H., Holub, F., “Investigation of Metallic System with Fluorspar Phases” (in


[1974Boc] Bockelmann, W., Schuster, H.-U., “Crystallographic Aspects of Ternary Phases of Li with

Group III A and IVA Elements in Ionic and Non-Ionic Compositions” (in German),

Z. Anorg. Allg. Chem., 410, 241-250 (1974) (Crys. Structure, Experimental, 5)

[1976Sch] Schuster, H.-U., Hinterhauser, H.-W., Schäfer, W., Will, G., “Neutron Diffraction

Investigations of the Phases LiAlSi and LiAlGe” (in German), Z. Naturforsch. B, 31,


[1978Ble] Blessing, J., “Synthesis and Studies of Ternary Phases of Li with Elements of the 3 and 4

Sub Groups” (in German), Thesis, Univ. Cologne, 167 pp. (1978) (Crys. Structure,

Experimental, 87)

[1981Kis] Kishio, K., Brittain, J.O., “Phase Stability of Doped -LiAl”, Mater. Sci. Eng., 49, P1-P6


54


MSIT®

Al–Ge–Li

[1981Gru] Gruttner, A., Nesper, R., Schnering, H.G., “New Phases in the Li-Ge System: Li7Ge12,

Li12Ge7, Li14Ge6”, Acta Crystallogr., 37A, 161 (1981) (Crys. Structure, Experimental, 5)

[1982Gru] Gruttner, A., “About the Lithium-Germanium System and Formation of Metastable

Germanium-Modifications from Li-Germanides” (in German), Diss. Dokt. Naturwiss.,

Chem. Fak. Univ. Stuttgart, 1-102 (1982) (Equi. Diagram, Crys. Structure, Experimental)

[1986Lee] Lee, J.J., Sommer, F., “Thermodynamic Properties of Lithium in Liquid Aluminium

Alloys” (in Korean), Tachan Kunsok Hakhoechi, 24(10), 1185-1189 (1986) (Thermodyn.,

Theory, 19)

[1986Gas] Cassada, W.A., Shiflet, G.J., Starke, Jr, E.A., “The Effect of Germanium on the

Precipitation and Deformation Behavior of Al-2Li Alloys”, Acta Metall., 34(3), 367-378

(1986) (Crys. Structure, Equi. Diagram, Experimental, 25)

[1987Eve] Evers, V.J., Oehlinger, G., Sextl, G., Becker, H.-O., “High Pressure LiGe with Layers of

Two- and Four-Bond Germanium Atoms” (in German), Angew. Chem., 99(1), 69-71 (1987)


[1989Goe] Goel, N.C., Cahoon, J. R., “The Al-Li-X Systems (X = Ag, As, P, B, Cd, Ge, Fe, Ga, H, In,

N, Pb, S, Sb and Sn)”, Bull. Alloy Phase Diagrams, 10(5), 546-548 (1989) (Review, 25)

[1992Pav] Pavlyuk, V.V., Dmytriv, G.S., Starodub, P.K., “Crystal Structure of the Compounds of the

Li-M-X (M = Mg, Al; X = Si, Ge, Sn) Systems” (in Russian), Cryst. Chem. Inorg. Coord.

Compounds, VI Conf. (Abstact), L’viv, 210 (1992) (Crys. Structure, Experimental, 6)

[1993Pav1] Pavlyuk, V.V., Dmytriv, G.S., Bodak, O.I., “Phase Equilibria in the Li-Al-Ge System at

470 K” (in Ukrainian), Dop. Akad. Nauk Ukrainy, (8), 84-86 (1993) (Equi. Diagram,

Experimental, #, 6)

[1993Pav2] Pavlyuk, V.V., “Synthesis and Crystal Chemistry of Lithium Intermetallic Compounds” (in

Ukrainian), Summary of the Thesis for Doctor Science Degree, L’viv Univ., 1-35 (1993)

(Crys. Structure, Experimental, Review, 49)

[1994Hos] Hosoda, H., Sato, T., Tezuka H., Mishima Y., Kamio A., “Substitution Behaviour of

Additional Elements in the L12-Type Al3Li Metastable Phase in Al-Li Alloys”, J. Jpn. Inst.

Met., 58(8), 865-871 (1994) (Calculation, 26)

[1995Pav] Pavlyuk, V., Bodak, O., MSIT Ternary Evaluation Program, in MSIT Workplace,

Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart;

Document ID: 10.14593.1.20, (1995) (Crys. Structure, Equi. Diagram, Assessment, 15)

[1996Dmy] Dmytriv, G.S., “Phase Equilibria and Crystal Structure of Compounds in Mg-Li-Si,

Ca-Li-{Si, Ge}, Al-Li-{Si, Ge, Sn}, Zn-Li-{Al, Sn}” (in Ukrainian), Summary of the Thesis

for the Degree of Candidate of Science, 1-23 (1996) (Crys. Structure, Equi. Diagram,

Experimental, 10)

[1997San] Sangster, J., Pelton, A.D., “The Ge-Li (Germanium-Lithium) System”, J. Phase Equilib.,

18(3), 289-294 (1997) (Calculation, Crys. Structure, Review, Thermodyn., 31)

[2000Oka] Desk Handbook: Phase Diagrams for Binary Alloys, Okamoto, H., (Ed.), ASM (2000)

(Equi. Diagram, Crys. Structure, Review)

[2001Gow] Goward, G.R., Taylor, N.J., Souza, D.C.S., Nazar, L.F., “The True Crystal Structure of

Li17M4 (M = Ge, Sn, Pb) - Revised from Li22M5”, J. Alloys Compd., 329, 82-91 (2001)


[2003Gro] Groebner, J., “Al-Li (Aluminium-Lithium)”, MSIT Binary Evaluation Program, in MSIT



Assessment, 21)

55


MSIT®

Al–Ge–Li


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Li)

< 180.6

cI2

Im3m

W

a = 351.0 pure Li at 25°C

[V-C2]

(Al)

< 660.452

cF4

Fm3m

Cu

a = 404.96 pure Al at 25°C [Mas2]

Dissolves up to 15 at.% Li

(Ge)

< 938.3

cF8

Fd3m

C (diamond)

pure Ge at 25°C [Mas2]

, Li9Al4< 347 - 275

mC26

C2/m

Li9Al4

a = 1915.51

b = 542.88

c = 449.88

= 107.671°

[2003Gro]

`, Li9Al4< 275

? ? [Mas2]

, Li3Al2< 520

hR15

R3m

Li3Al2

a = 450.8

c = 1426

[2003Gro]

60 to 61 at.% Li [Mas2]

, LiAl

< 700

cF16

Fd3m

NaTl

a = 637 at 50 at.% Li [2003Gro]

45 to 55 at.% Li [Mas2]

46 to 52 at.% Li at 200°C [1993Pav1]

`, LiAl3< 190 - ~120

cP4

Pm3m

Cu3Au

a = 403.8 Metastable [2003Gro]

Li7Ge12

< 510

oP*

Pnm21

Li7Ge12

a = 1154.1 0.3

b = 807.3 0.2

c = 1535.9 0.4

[1981Gru, 1982Gru]

LiGe

< 540

tI32

I41/a

MgGa

tI24

I41/amd

LiGe

a = 975 2

c = 578 2

a = 981.0 0.3

c = 580.7 0.2

a = 405.29 0.01

c = 2328.2 0.3

[1987Eve]

[1982Gru]

high pressure phase [1987Eve]

Li12Ge7

< 510

oP152

Pnma

Li12Si7

a = 876.3

b = 2011.5

c = 1464

[1981Gru, 1982Gru]

Li9Ge4

< 740

oC52

Cmcm

Na9Sn4

h**

a = 449

b = 787

c = 2444

a = 449

c = 2444

[V-C2, 1982Gru]

[1982Gru]

Li14Ge6

< 770

hR21

R3m

Li14Si6

a = 449.4 0.1

c = 1843.9 0.4

[1981Gru, 1982Gru]

56


MSIT®

Al–Ge–Li

Li13Ge4

< 780

oP34

Pbam

Li13Si4

a = 924

b = 1321

c = 463

[V-C2, 1982Gru]

Li15Ge4

< 720

cI76

I43d

Cu15Si4

a = 1072

a = 1082.5

[V-C2]

[1982Gru]

Li4Ge

< 640

cF416

F43m

Li20Si5cF432

F23

Li22Pb5

cF419

F43m

Li17Ge4

a = 1892.9 0.1

a = 1886

a = 1875.6 0.2

[1982Gru]

Li22Ge5 [V-C2, Mas2]

Li17Ge4 [2001Gow]

* 1, Li5.3Al0.7Ge2

* 1´, Li16Al2Ge6

hP8

P63/mmc

Na3As

hP24

a = 438.0

c = 816.2

a = 438.0

c = 816.2

a = 758.6

c = 816.2

[1978Ble]

m = 2.42 g·cm-3

x = 2.46 g·cm-3

[1993Pav1]

[1978Ble]

* 2, Li9Al2Ge3 ? ? [1993Pav1]

* 3, Li2AlGe cF

F43m

CuHg2Ti

a = 616.3

a = 597.5

[1974Boc]

m = 2.848 g·cm-3

[1992Pav, 1993Pav1]

* 4, Li6Al3Ge ? ? [1993Pav]

* 5, LiAlGe cF16

F43m

LiAlSi

cF16

Fd3m

NaTl

a = 598.9

a = 598.9

a = 597.7

[1976Sch]

m = 3.27 g·cm-3

x = 3.29 g·cm-3

[1992Pav, 1993Pav1]

[1960Now]

[1981Kis]

* 6, LiAl2Ge cF16

Fm3m

MnCu2Al

a = 599.8 [1992Pav, 1993Pav1]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

57


MSIT®

Al–Ge–Li

80 60 40 20

0

100

200

300

400

500

600

700

800

900

1000

Li Ge

Li, at.%

Te

mp

era

ture

, °C

Li4Ge

LiGeLi15Ge4Li13Ge4

Li14Ge6

Li9Ge4

Li12Ge7Li7Ge12

L

530

510

180

530

640

510540

500

730

770780

720740

690

Fig. 1: Al-Ge-Li.

Phase diagram of the

Ge-Li system after

[1982Gru]

20

40

60

80

20 40 60 80

20

40

60

80

Li Al

Ge Data / Grid: at.%

Axes: at.%

τ1

τ2

τ3

βγδ´

Li4Ge

Li15

Ge4

Li13

Ge4

Li9Ge

4

LiGe

Li12

Ge7

Li14

Ge6

Li7Ge

12

(Al)

(Ge)

τ4

τ5

τ6

Fig. 2: Al-Ge-Li.

Partial triangulation

of the Al-Ge-Li

ternary system

58


MSIT®

Al–H–Li

Aluminium – Hydrogen – Lithium

Oksana Bodak, Pierre Perrot

Literature Data

Two ternary hydrides have been prepared and characterized, LiAlH4 and Li3AlH6. The hydride LiAlH4 is

available as a commercial product. Crystal structure data for Li3AlH6, obtained by the reaction of LiAlH4,

LiH and Al(C2H5)3 in C6H5CH3, were given by [1966Chi], Table 1. The crystal structure data for the

Li3AlH6 are given in [1985Bas2]. The thermal stability of LiAlH4 was studied by [1970Bra] using DTA,

by [1972Dil] using DTA and thermogravimetric analysis and by [1985Bas1] using DSC. The first critical

review of literature data, published until 1990, was made by [1993Fer, 1995Pav], followed by the present

evaluation. The influence of mechanochemical processing of polycrystalline LiAlH4 was studied in

[1999Zal, 2000Bal]. The enthalpy of formation LiAlHx was calculated using the Miedema’s model

[2002Her].

Binary Systems

The Al-Li system reported by [2003Gro] and the Al-H system as described by [2003Per] are accepted as

terminal descriptions of the ternary Al-H-Li phase diagram. The H-Li is accepted from [Mas2].

Solid Phases

All authors completely agree that two hydrides, LiAlH4 and Li3AlH6, are formed in this system. Their

crystal structures were reported by [1967Skl, 1970Gor, 1985Bas1, 1985Bas2, 2000Bal] and are given in

Table 1. [1967Skl] proposed a unit cell with an “a” parameter only half of what was adopted by the other

workers. For the remaining cell parameters there is good agreement between the reported data.

Mechanochemical processing of polycrystalline LiAlH4 revealed good stability of this complex

aluminohydride during high-energy ball-milling in a helium atmosphere for up to 35 h. The decomposition

of lithium aluminohydride into Li3AlH6, Al and H2 is observed during prolonged mechanochemical

treatment for up to 110 h and is most likely associated with the catalytic effect of a third material, iron,

which is introduced into the hydride as a contaminant during mechanical treatment [2000Bal]. According

to [2000Bal] the attempts to solve the crystal structure of Li3AlH6 by X-ray powder diffraction data were

unsuccessful because of the strong pseudosymmerty found in this compound. The unit cell volume of the

rhombohedral lattice is 1.5 times greater than that of both primitive and base centered monoclinic lattices.

Isothermal Sections

LiAlH4 has a melting point of 163.7°C and decomposes at 160-180°C [1999Zal] according to the reaction:

3LiAlH4(liquid) Li3AlH6(solid) + 2Al + 3H2.

The standard Gibbs energy of this reaction at 298K was assessed to be -27.7 kJ·mol-1 [2000Bal].

At temperatures above 250°C the hydride Li3AlH6 decomposes:

Li3AlH6 3LiH + Al + 3/2H2.

According to [2000Bal] the temperature of decomposition is in the range 207-260°C which is in good

agreement with data of [1999Zal]. The analogous ternary deuteride, LiAlD4, has its melting point at

167.5°C and decomposes at 195°C [1985Bas1]. The phase stability diagram at 500°C calculated by

[1988Cro] is based on the assumption that only the LiAl phase occurs in the Al-Li binary system. The three-

phase regions identified were: Li+LiH+LiAl, H+LiH+LiAl and Al+H+LiAl. It should be pointed out that,

if the hydrides are considered to be unstable at 500°C [1985Bas1], Al would react with LiH following the

reaction: Al + LiH LiAl + 1/2H2 and the tie line LiAl-H of the Al-Li-H stability diagram would be stable.

However, hydrides are stable under large hydrogen pressure and the existence of the LiAl-H tie line

contradicts the decomposition of LiAlH4 into Li3AlH6+Al and the subsequent decomposition of Li3AlH6

into LiH+Al [1999Zal]. Figure 1 shows a stability diagram taking into account experimental observation.

59


MSIT®

Al–H–Li

Each of the triangles numbered 1 to 5 is characterized by a hydrogen pressure depending on the given

temperature and decreasing from p1 to p5:

p1 = p(AlH3/Al),

p2 = p(LiAlH4/Li3AlH6 + Al),

p3 = p(Li3AlH6/LiH + Al),

p4 = p(LiH + Al/AlLi),

p5 = p(LiH/Li).

In Fig. 1 the dashed lines correspond to tie lines never observed experimentally. Between p4 and p5 one

should actually observe the following equilibria: LiH + LiAl/Li3Al2 and LiH + Li3Al2/Li9Al4.

The solubility of hydrogen in equiatomic LiAl alloys was measured at 500°C as a function of hydrogen

pressure between 204 and 716 mbar (204 102 and 716 102 Pa) by [1976Tal]. Sieverts’ Law was obeyed,

with an average value of Sieverts’ constant of 2.20 104 0.15 mbar1/2/atomic fraction H2 (Table 2).

[1988Any] determined the solubility of H2 in molten Al-Li alloys containing 1, 2 and 3 mass% Li (3.8, 7.4

and 10.7 at.% Li, respectively) from 670°C to 800°C and from 5.3 104 Pa to 10.7 104 Pa. Sieverts’ Law was

obeyed for all three alloys; the solubility of H2 increases with increasing Li content (Table 3). [1990Fed]

quoted data for the solubility of H2 in the Al-2Li (mass%) alloy. At 700°C the data are in good agreement

with [1988Any]. The solubility of H2 in molten Al-Li alloys containing up to 4 mass% Li was measured by

[1989Lin] for temperatures of 700, 800, 900 and 1000°C. At 700°C the calculated H2 solubilities are lower

than determined by [1988Any, 1990Fed]. Interaction parameters for Al-H-Li melts were calculated for

927°C by [1986Lee]. A more general expression of the first order interaction parameter of Li upon H has

been proposed by [2003Ma]: eH(Li) = (d ln H/ d (mass% Li)) = -0.138 - 158.2/T.

A negative value of the interaction parameter means that the presence of Li increases the solubility of H in

liquid Al; this result is already confirmed experimentally by [1988Any] and theoretically by [1989Lin].

Thermodynamics

The molar heat capacity of LiAlH4 [1978Cla, 1979Bon, 1985Bas1], of LiAlD4 [1985Bas1] and Li3AlH6

[1978Cla, 1979Bon] at 298.15 K are given in Table 4. According to [2002Her] the calculated enthalpy of

formation using the Miedema’s model was -69 kJ·mol-1 for LiAlH4 and -86 kJ·mol-1 for Li3AlH6.


Besides its well-known application as a reducing agent in organic synthesis, LiAlH4 contains 10.5 mass%

H, which is one of the highest values among hydrides. Thus LiAlH4 is of considerable interest as potential

ultra-high capacity hydrogen storage solid.

Miscellaneous

[1982Wak] determined the electrical resistance of LiAlH4 at pressures up to 125 kbar. The resistance

decreases with applied pressure up to 75 kbar and remains virtually constant from 75 to 125 kbar.

Adsorption and desorption of hydrogen in Al and Al-Li alloys were presented and discussed by [1988Wat].

References

[1965Ame] Amendola, A., Index Inorganic to the Powder Diffraction File 1965, American Society for

Testing and Materials, Philadelphia, Pa, n.12473, p.469 (1965) as quoted in [1970Gor]

[1966Chi] Chini, P., Baradel, A., Vacca, C., “The Reaction of Aluminum with Hydrogen and

Natriumfluoride” (in Italian), La Chimica e l’Industria, Special, 48(6), 596-601 (1966)


[1967Skl] Sklar, N., Post, B., “The Crystal Structure of LiAlH4”, Inorg. Chem., 6, 669-671 (1967)


[1970Bra] Brachet, F.-G., Etienne, J.-J., Mayet, J., Tranchant, J., “Structure and Properties of LiAl

Hydrides. III. Differential Thermal Analysis and Isothermal (70 and 130°C) Thermal

60


MSIT®

Al–H–Li

Decomposition of LiAlH” (in French), Bull. Soc. Chim. Fr., (11), 3799-3807 (1970)

(Experimental, 14)

[1970Gor] Gorin, P., Marchon, J. C., Tranchant, J., Kovacevic, S., Marsault, J. P., “Structure and

Properties of LiAl Hydrides. II. Structure of LiAlH4 in the Crystalline State and in Diethyl

Ether Solutions” (in French), Bull. Soc. Chim. Fr., (11), 3790-3799 (1970) (Crys. Structure,

Experimental, 27)

[1972Dil] Dilts, J.A., Ashby, E.C., “A Study of the Thermal Decomposition of Complex Metal

Hydrides”, Inorg. Chem., 11(6) 1230-1236 (1972) (Experimental, 27)

[1976Tal] Talbot, J. B., Smith, F. J., Land, J. F., Barton, P., “Tritium Sorption in Li-Bi and Li-Al

Alloys”, J. Less-Common Met., 50, 23-28 (1976) (Experimental, 10)

[1978Cla] Claudy, P., Bonnetot, B., Letoffe, J.M., Turck, G., “Determination of Thermodynamic

Constants of Simple Hydrides of Aluminium. IV. Enthalpy of Formation of LiAlH4 and

Li3AlH6” (in French), Thermochim. Acta, 27, 213-221 (1978) (Thermodyn.,

Experimental, 11)

[1979Bon] Bonnetot, B., Claudy, P., Diot, M., Letoffe, J.M., “LiAlH4 and Li3AlH6: Molar Heat

Capacity and Thermodynamic Properties from 10 to 300K”, J. Chem. Thermodyn., 11,

1197-1202 (1979) (Thermodyn., Experimental, 8)

[1981Gor] Gorbunov, V.E., Gavrichev, K.S., Bakum, S.I., “Thermodynamic Properties of LiAlH4 in

the Temperature Range 12-300 K”, Russ. J. Inorg. Chem. (Engl. Transl.), 26, 168-169

(1981) (Thermodyn., Experimental, 8)

[1982Wak] Wakamori, K., Sawaoka, A., Filipek, S.M., Baranowski, B., “Electrical Resistance of Some

Alkaline Earth Metal Hydrides and Alkali Metal Al Hydrides and Borohydrides Under High

Pressure”, J. Less-Common Met., 88, 217-220 (1982) (Experimental, 6)

[1985Bas1] Bastide, J.-P., Bonnetot, B., Letoffe, J.-M., Claudy, P., “Comparative Study of LiAlH4 and

LiAlD4. I. Preparation, Crystallography and Thermal Behaviour- Evidence for a Metastable

Form of LiAlD4”, Mater. Res. Bull., 20, 999-1007 (1985) (Crys. Structure,

Experimental, 16)

[1985Bas2] Bastide, J.-P., Bonnetot, B., Letoffe, J.-M., Claudy, P., “Structural Chemistry of Some

Complex Hydrides of Alkaline Metals”, Stud. Inorg. Chem., 3, 785-788 (1983) (Crys.


[1986Lee] Lee, J.J., Sommer, F., “Thermodynamic Properties of Li in Liquid Aluminum Alloys” (in

Korean), Taehan Kumsok Hakhoechi, 24(10), 1185-1189 (1986) (Thermodyn., Theory,

Experimental, 23)

[1988Any] Anyalebechi, P.N., Talbot, D.E., Granger, D.A., “The Solubility of H2 in Liquid Binary Al-

Li-Alloys”, Metall. Trans. B, 19, 227-232 (1988) (Thermodyn., Experimental, 24)

[1988Cro] Crouch-Baker, S., Huggins, R.A., “Phase Behaviour in the Li-Al-O-H System at

Intermediate Temperatures”, Solid State Ionics, 28-30, 611-616 (1988) (Equi. Diagram,

Thermodyn., Theory, 21)

[1988Wat] Watson, J.W., “Hydrogen in Aluminum and Aluminum-Lithium Alloys”, Thesis,

Northwestern University, 1-366 (1988) (Experimental, 147)

[1989Lin] Lin, R.Y., Hoch, M., “The Solubility of Hydrogen in Molten Aluminum Alloys”, Metall.

Trans. A, 20(9), 1785-1791 (1989) (Equi. Diagram, Thermodyn., Calculation, Theory, 31)

[1990Fed] Fedosov, A.S., Danilkin, V.A., Makarov, G.S., “The Interaction of Al-Li Alloy Metals with

Hydrogen” (in Russian), Tsvetn. Met., (8), 88-90 (1990) (Experimental, 4)

[1993Fer] Ferro, R., Saccone, A., Delfino, S., “Aluminium-Hydrogen-Lithium”, in “Ternary Alloys:

A Comprehensive Compendium of Evaluated Constitutional Data and Phase Diagrams”

Petzow, G., Effenberg, G. (Eds.), Vol. 6, VCH, Weinheim, 111-112 (1993) (Crys. Structure,

Review, 9)

[1995Pav] Pavlyuk, V., Bodak, O., “Aluminium-Hydrogen-Lithium”, MSIT Ternary Evaluation

Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International

Services GmbH, Stuttgart; Document ID: 10.12744.1.20, (1995) (Crys. Structure, Equi.

Diagram, Assessment, 17)

61


MSIT®

Al–H–Li

[1999Zal] Zaluski, L., Zaluska, A., Ström-Olsen, J.O., “Hydrogenation Properties of Complex Alkali

Metal Hydrides Fabricated by Mechano-Chemical Synthesis”, J. Alloys Compd., 290, 71-78


[2000Bal] Balema, V.P., Pecharsky, V.K., Dennis, K.W., “Solid State Transformations in LiAlH4

during High-Energy Ball-Milling”, J. Alloys Compd., 313, 69-74 (2000) (Equi. Diagram,

Crys. Structure, Experimental, 22)

[2002Her] Herbst, J.F., “On Extending Miedema’s Model to Predict Hydrogen Content in Binary and

Ternary Hydrides”, J. Alloys Compd., 337, 99-107 (2002) (Thermodyn., Calculation, 20)



Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 21)

[2003Ma] Ma, Z., Janke, D., “Solution Behawior of Hydrogen in Aluminium and ist Alloys Melts”,

Metall, 57(9), 552-556 (2003) (Thermodyn., Calculation, Review, 14)

[2003Per] Perrot, P., “Al-H (Aluminium-Hydrogen)”, MSIT Binary Evaluation Program, in MSIT




Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Li)

< 180.6

cI2

Im3m

W


[V-C2]

(Al)

< 660.45

cF4

Fm3m

Cu

a = 404.96 pure Al at 25°C

[Mas2]

dissolves up to 15 at.% Li

Li9Al4< 347 - 275

mC26

C2/m

Li9Al4

a = 1915.51

b = 542.88

c = 449.88

= 107.671°

[2003Gro]

Li9Al4 ( ´)

< 275

? ? [Mas2]

Li3Al2 ( )

< 520

hR15

R3m

Li3Al2

a = 450.8

c = 1426

[2003Gro]

60 to 61 at.% Li

[Mas2]

LiAl ( )

< 700

cF16

Fd3m

NaTl

a = 637 at 50 at.% Li [2003Gro]

45 to 55 at.% Li [Mas2]

LiAl3 ( ´)

< 190 - ~120

cP4

Pm3m

Cu3Au

a = 403.8 metastable [2003Gro]

LiH cF8

Fm3m

NaCl

a = 408.3 [V-C2]

AlH3

< 110

hR24

R3c

a = 445.6

c = 1183

[2003Per]

metastable

62


MSIT®

Al–H–Li

AlH3

< 80

[2003Per] metastable “Aluminum

hydrogenoaluminate”

Al(AlH4)3

* LiAlH4 mP48

mP48

mP48

mP24

a = 960

b = 786

c = 790

= 112.5°

a = 967.9

b = 781.0

c = 792.5

= 112.53°

a = 967.9

b = 788.1

c = 791.2

= 111.88°

a = 484.5

b = 782.6

c = 791.7

= 112.5°

[1965Ame]

x = 0.908 g·cm-3

m = 0.917 g·cm-3

[1970Gor]

x = 0.911 g·cm-3

m = 0.95 g·cm-3

[1985Bas1]

x = 0.900 g·cm-3

m = 0.907 g·cm-3

[1967Skl]

x = 0.904 g·cm-3

m = 0.92 g·cm-3

* Li3AlH6 m** a = 571.5

a = 539.1

c = 569.4

= 91.33°

[1966Chi]

* -Li3AlH6 mP*

P21/m

LiAlSi2O6

mP*

P21/c

mC*

C2/m

hR*

R3m

a =790.5

b = 812.5

c = 567.5

= 92.7°

a = 566.7 0.1

b = 810.7 0.2

c = 791.7 0.2

= 92.07 0.01°

a = 791.7 0.2

b = 810.7 0.2

c = 566.7 0.1

= 92.07 0.01°

a = 811.3 0.1

c = 957.0 0.1

[1985Bas2]

high-pressure phase, 500°C, 50 kbar,

pseudo-cubic

[2000Bal]

prepared mechano-chemically

[2000Bal]

[2000Bal]

* -Li3AlH6 o*

Li3Al2 (LiF4)3

a = 1114

b = 1145

c = 1034

[1985Bas2]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

63


MSIT®

Al–H–Li

Table 2: Solubility of H2 in LiAl at 500°C [1976Tal]

Table 3: Solubility of H2 in Molten Al-Li Alloys. S is the solubility expressed in cm3 H2 measured at 273

K and 101.325 Pa; S° is the standard value: S° = 1cm3 measured at 273 K and 101.325 Pa; p is the

pressure expressed in Pa; p° is the standard pressure: p° = 101.325 Pa

Table 4: Molar Heat Capacity of LiAlH4, LiAlD4 and Li3AlH6

H2 Pressure, (p)

[mbar]

H2 Concentration, (N)

[atomic fraction]

Sieverts’ Constant,

(p/N)1/2 [mbar]-1/2 atomic fraction H2)

204

307

420

716

6.04 10-4

8.07 10-4

9.94 10-4

12.2 10-4

2.36 104

2.17 104

2.06 104

2.19 104

1 mass% Li:

2 mass% Li:

3 mass% Li:

log(S/S°) - 1/2 log(p/p°) = -2113/T + 2.568

log(S/S°) - 1/2 log(p/p°) = -2997/T + 3.329

log(S/S°) - 1/2 log(p/p°) = -2889/T + 3.508

Phase Molar Heat Capacity, Cp, [J K-1 mol-1] at 296.15 K Reference

LiAlH4 89.2

83.19

83.01

82.60

[1978Cla]

[1979Bon]

[1981Gor]

[1985Bas1]

LiAlD4 92.70 [1985Bas1]

Li3AlH6 131.0

127.75

[1978Cla]

[1979Bon]

20

40

60

80

20 40 60 80

20

40

60

80

Li Al

H Data / Grid: at.%

Axes: at.%

LiAlH4

Li3AlH

6

AlH3

LiH

Li9Al

4Li

3Al

2LiAl

12

3

4

5

(Al)

Fig. 1: Al-H-Li.

Stability diagram.

The triangles 1 to 5

are characterized by

hydrogen pressure at

equilibrium

decreasing from

p1(AlH3/Al) to

p5(LiH/Li) (see text)

64


MSIT®

Al–H–Mg

Aluminium – Hydrogen – Magnesium

Lazar Rokhlin, updated by Volodymyr Ivanchenko

Literature Data

The solubility of hydrogen in Al-Mg alloys was measured for different temperatures and composition

ranges using a range of different experimental techniques. [1973Hua] used a modified Sieverts apparatus

for determination of solubility of hydrogen in pure magnesium and its alloys including Al-Mg system. It

was shown that alloying of magnesium with 10 at.% Al lowered the solubility of hydrogen at 700°C and

pH2 = 105 Pa from 50 cm3 H2/100 g to 40 cm3 H2/100 g (hydrogen volumes measured at 273 K under

101325 Pa). These values are very close to the values calculated by [1965Bur]. [1974And] studied the

solubility of hydrogen in (Al) solid solution with 0.45 and 4.75 at.% Mg at 500°C using saturation and

vacuum extraction and showed that alloying with Mg raised the hydrogen solubility from

0.012 cm3 H2/100 g (for pure Al) to 0.04 0.01 (for 0.45 at.% Mg) and to 0.06 cm3 H2/100 g (for 4.75 at.%

Mg). These results are significantly lower than those presented by [1976Wat]. [1974Gab] studied the

solubility of hydrogen in phase (Mg2Al3) in temperature interval from 380 to 560°C using high pressure

Sieverts apparatus and high temperature vacuum extraction. Under crystallization the hydrogen solubility

in Mg2Al3 dropped from 5.9 cm3 H2/100 g to 1.45 cm3 H2/100 g. [1976Lev] studied the porosity of Al-Mg

alloys which is caused by hydrogen. [1977Che] investigated permeability, diffusivity and solubility of

hydrogen at temperatures from 650 to 800°C in liquid Al-Mg alloys containing up to 16 mass% Al.

[1981Tuc] studied the hydrogen saturation of Al-Mg alloys exposed to water-vapor saturated air at elevated

temperature.

Reversible hydrogen storage in magnesium alloys was reviewed by [1978Gui]. They reported that phase

(Mg17Al12, sometimes designated as Mg3Al2) did not hydride at 350°C under hydrogen pressure from 3 to

5 MPa. These results are in contradiction with those presented by [1980Min, 1981Gav], who studied the

reactions of hydrogen with Mg2Al3 and Mg17Al12 and reported their main features: hydrogenation of the

intermetallic Al-Mg compounds resulting in disproportionation; namely, for Mg17Al12 the reaction may be

written as: Mg2Al3+2H2 2MgH2+3Al; while for Mg3Al2, the reaction may be written:

Mg17Al12+9H2 9MgH2+4Mg2Al3 [1983Sem] pointed that Al-Mg alloys dissolved only a very small

quantity of hydrogen due to very low rate of process.

Differential scanning calorimetry and gas chromatography were used to investigate and quantify the

reactions occurring when Al-5Mg (mass%) alloy, previously exposed to water-vapor saturated air, were

heated from ambient temperature to 600°C.

[1984Lue] measured the equilibrium hydrogen pressure at 142 and 170°C of the three phase fields

MgH2+(Mg)+ , MgH2+ + , and MgH2+ +(Al). The H was introduced into Al-Mg alloys by electrolysis

in an organometallic melt, NaAlEt4, containing dissolved Na+H- as electrolyte. [1985Lue1, 1985Lue2]

discussed the results thermodynamically.

[1978Cla1] prepared a ternary hydride Mg(AlH4)2 by reaction of NaAlH4 with MgCl2 dissolved in

tetrahydrofurane and measured its heat capacity at room temperature, to be 136 J (mol K)-1. Since

[1985Lue1, 1985Lue2] did not find this ternary hydride, it may possibly be stable only under high hydrogen

pressure, an assumption supported by the method of sample preparation.

A new theoretical method of describing and investigating metal hydrides has been developed by [1987Lue].

It involves thermodynamics and interprets the hydrogenation reaction by ternary phase diagrams.

[1987Lue] showed that intermetallic compounds formed by elements of the boron group with magnesium

form two phase regions with MgH2 in the ternary phase diagrams. Thus the hydrogen pressures of the

resulting three phase equilibria will be higher than or equal to the value of Mg/MgH2 equilibrium. In these

systems, including Al-H-Mg, ternary hydrides are not taken into account.

The solubility of hydrogen in molten aluminium alloys containing magnesium has been calculated from the

solubility of hydrogen in pure metals and binary metal-metal interaction parameters by [1989Lin].

65


MSIT®

Al–H–Mg

The structure and hydrogen absorption properties of Al-Mg alloys prepared by high-energy ball milling

were studied over the whole composition range in their as milled and Al-leached forms by [2000Bou]. The

latter were obtained from the milled materials by leaching out of Al in a 1N NaOH solution. Their results

on the interaction of intermetallic phases with hydrogen are in good agreement with those of [1980Min,

1981Gav].

[2002Her] used Miedema’s model to predict the hydrogen content and the enthalpy of formation of

hypothetical ternary hydrides in the Al-H-Mg system.

Binary Systems

The binary systems Al-H [2002Per], Al-Mg [2003Luc] and H-Mg [2001Per] are accepted to present the best

boundary systems for the Al-H-Mg ternary system.

Solid Phases

One ternary phase has been reported, Mg(AlH4)2, which is stable under high hydrogen pressure. The phase

AlH3 is known to have two polymorphic modifications which are both metastable [1978Cla1, 1978Cla2,

1979Cla]. Chemically AlH3 is stable at room temperature and decomposes when heated at 110°C

[1980Her]. Under high hydrogen pressures (2 GPa at 300°C and 6 GPa at 600°C), it is possible to synthesize

AlH3 reversibly [1992Kon]. All solid phases are listed in Table 1.

Isothermal Sections

Figure 1 shows the isothermal section between 140 and 170°C [1984Lue, 1985Lue1, 1985Lue2]. The

section is corrected to the accepted homogeneity ranges of the Al-Mg phases: ( Al), (Mg), and . Figure 1

shows that MgH2 is in equilibrium with , , , and (Mg) phases. Mg(AlH4)2 and AlH3 hydrides are stable

phases at these temperatures under hydrogen pressure higher than 100 kPa.

Figure 2 shows the solubility of hydrogen in liquid Al-Mg alloys at 500, 700 and 800°C. It is taken from

[1976Wat, 1989Lin] with small corrections to match the solubility in Al given in the Al-H system by

[2002Per]. From the activity coefficients of hydrogen in molten Al-Mg alloys at 827°C [1989Lin], the

interaction coefficient of Mg upon H in liquid Al may be assessed: H(Mg) = (dln H/dxMg) = -8.12 at 827°C.

This negative value means that Mg in liquid Al increases the solubility of H.

Al-Mg alloys show liquid-solid two-phase fields at 500°C. There, the hydrogen solubility must be

represented by a straight line. The pressure-composition isotherms of the Mg2Al3-H system for temperature

interval from 335 to 410°C are presented in Fig. 3 [1980Min]. These isotherms are in fair agreement with

the measurements of [2000Bou] at 350°C who observed a plateau towards 0.8 MPa for Mg75Al25,

corresponding to the Mg-MgH2 equilibrium and a plateau towards 1 MPa for Mg58Al42, corresponding to

the equilibrium / +MgH2.

Figure 3 shows that under 5 MPa H2 the global composition of the hydride is Mg2Al3H7. The corresponding

point lies inside the (Al)-MgH2-MgAl2H6 triangle in Fig. 1, which confirms the formation of the ternary

phase under high hydrogen pressures.

Thermodynamics

The dependence of the equilibrium pressure on temperature for the disproportionation reaction

1/2Mg2Al3+H2=MgH2+3/2Al was reported by [1980Min, 1981Gav] as log10(p/Pa) = -3306/T+11.47.

Hydrogen activity and Gibbs energy changes for the three-phase reactions in the Al-H-Mg system, as

measured electrochemically at 142°C were presented by [1985Lue1, 1985Lue2, 1987Lue], as

Mg17Al12 - Mg - MgH2 aH2 = 2.7 10-3 G = - 20080 J (mol H2)-1

Mg2Al3 - Mg17Al12 - MgH2 aH2 = 1.1 10-2 G = - 15481 J (mol H2)-1

Al -Mg2Al3 - MgH2 aH2 = 2.3 10-2 G = - 12970 J (mol H2)-1.

These values are about 1.5 kJ lower than the accepted values. For instance, the first figure ( G = -20080

J (mol H2)-1) which corresponds to the Mg/MgH2 equilibrium has to be compared with the value (-18833

J (mol H2)-1) accepted by [2001Per] at 142°C. The last value ( G = -12970 J (mol H2)-1) has to be

66


MSIT®

Al–H–Mg

compared with ( G = -11870 J (mol H2)-1) calculated from the expression of [1980Min, 1981Gav] given

above.

[2002Her] predicted the enthalpy of formation of some virtual hydrides as

Mg17Al12Hx xcalc = 29.64 Hcalc (Xcalc) = - 65000 J (mol f.u.)-1

Mg2Al3Hx xcalc = 5.83 Hcalc (Xcalc) = - 47000 J (mol f.u.)-1

MgAl2Hx xcalc = 3.37 Hcalc (Xcalc) = - 43000 J (mol f.u.)-1

The molar heat capacity of the hydride Mg(AlH4)2 has been measured at 25°C by means of a Calvet

microcalorimeter as Cp = 136 J mol-1 K-1.


The Al-Mg system is of great importance for developing of many of the Al based and Mg based

multicomponent light alloys used in avionic and space industry. The Al-Mg alloys are also of potential

interest as materials for hydrogen storage.

Miscellaneous

The alloying of Al with Mg dramatically raises the absorption capacity of Al [1976Lev]. [1981Tuc] showed

evidence for the formation of MgH2 on the grain boundaries of Al-Mg alloys when exposed to water-vapor

saturated air at 70°C and for about 50 days. These authors suggest that its presence plays a prominent role

in the pre-exposure embrittlement and stress-corrosion cracking of Al-Mg alloys.

The diffusion of hydrogen in liquid Al-Mg alloys at temperatures from 650 to 800°C is slowly changed by

raising Al contents up to 5.5 mass% Al [1977Che]. The faster rise of DH was observed in concentration

interval of 5.5 to 12 mass% Al; after that DH raised slowly up to 16 mass% Al. For pure magnesium

DH(650°C) = 1.5 10-8 m2 s-1 and activation energy is ED = 31380 1670 J mol-1. For Al-Mg alloys:

DH( 5.5 mass% Al, 650°C) = 1.7 10-8 m2 s-1 and ED(5.5 mass% Al) = 34730 1670 J mol-1; DH(12 mass%

Al, 650°C) = 8 10-8 m2 s-1 and ED(12 mass% Al) = 33470 1670 J mol-1; DH(16 mass% Al, 650°C) =

1 10-7 m2 s-1 and ED(16 mass% Al) = 33470 1670 J mol-1. But at 7.5 mass% the Al activation energy has

a maximum at ED(7.5 mass% Al) = 50210 1670 J mol-1.

[2000Bou] showed that the measured hydrogen capacity of the as milled material decreases with Al content,

from H/M = 1.74 for pure un-milled Mg, to 1.38 for Mg/Al = 90/10, and then to 1.05 for Mg/Al = 75/25. In

each case, there is a further 10-15% decline of the hydrogen absorption capacity after leaching. In the case

of Mg/Al = 58/42, which basically contains a nanocrystalline Mg17Al12 intermetallic phase, only, hydriding

leads to the formation of MgH2 and Al. This reaction is totally reversible and Mg17Al12 is recovered upon

dehydriding. In each case, there is an increase in the kinetics of hydrogen absorption and desorption

following leaching. This change in the sorption kinetics is thought to arise as a consequence of the presence

of Al solutes in the hexagonal structure of Mg, rather than to be due to purely geometric effects, such as the

increase of the surface area.

References

[1965Bur] Burylev, B.P., “The Solubility of Hydrogen in Magnesium Alloys” (in Russian), Liteynoe

Proizvod., 9(1), 25-26 (1965) (Thermodyn., Theory, 11)

[1973Hua] Huang, Y.C., Watanabe, T., Komatsu, R., “Hydrogen in Magnesium and its Alloys”, Proc.

4th Internat. Conf. Vacuum Metallurgy, 176-179 (1973, published 1974) (Experimental, 8)

[1974And] Andreev, L.A., Levchuk, B.V., Gel’man, B.G., Danilkin,V.A., Kharin, P.A., Myagkov,

E.A, “The Solubility of H in Al-Mg Alloys” (in Russian), Tekhnol. Legk. Splavov, Nauch.

Byul. VILSa., (4), 58-62 (1974) (Experimental, 8)

[1974Gab] Gabidullin, R.M., Shvetsov, I.V., Kolachev, B.A., Archakov, Yu.I., “The Solubility of

Hydrogen in Intermetallic Compounds of Aluminium with Magnesium, Copper,

Manganese, Titanium and Zirconium” (in Russian), in “Constitution, Properties and

Application of Metallides”, Kornilov I.I., Matveeva N.M., (Eds.), Nauka, Moscow, 188-190


67


MSIT®

Al–H–Mg

[1976Lev] Levchuk, B.V., Andreev, L.A., “Interaction of Al-Mg Alloys with H” (in Russian),

Metalloved. Term. Obrab. Met., (7), 23-27 (1976) (Experimental, 10)

[1976Wat] Watanabe, T., Tachihara, T., Huang, Y.C., Komatsu, R., “The Effect of Various Alloying

Elements on the Solubility of Hydrogen in Magnesium” (in Japanese), J. Jpn. Inst. Light

Met., 26(4), 167-174 (1976) (Experimental, 25)

[1977Che] Chernega, D.F., Gotvyanskii, Yu.Ya., Prisyazhnyuk, T.N., “Permeability, Diffusion and

Solubility of Hydrogen in Magnesium-Aluminum Alloys” (in Russian), Liteinoe Proizvod.,

(12), 9-10 (1977) (Experimental, 4)

[1978Cla1] Claudy, P., Bonnetot, B., Letoffe, J.M., Turck, G., “Determination of the Thermodynamic

Constants of Simple and Complex Al Hydrides. II. Measurements of Molar Heat Capacities

at 298 K” (in French), Thermochim. Acta, 27, 199-203 (1978) (Thermodyn.,

Experimental, 10)

[1978Cla2] Claudy, P., Bonnetot, B., Letoffe, J.M., “Determination of Thermodynamic Constants of

Simple and Complex Aluminium Hydrides. III. Enthalpy of Formation of AlH3 and

AlH3” (in French), Thermochim. Acta, 27, 205-211 (1978) (Thermodyn.,

Experimental, 12)

[1978Gui] Guinet, P., Halotier, D., Perroud, P., “Hydrogen Sorage by Means of Reversible Magnesium

Alloys”, Eur. Communities Rep., EUR 1978, EUR 6085. Semin. Hydrogen Energy Vector:

Prod., Use, Transp., 373-391 (1978) (Experimental, 20)

[1979Cla] Claudy, P., Bonnetot, B., Letoffe, J.M., “Preparation, Physicochemical Properties and

Enthalpy of Formation of Aluminium Hydride -AlH3” (in French), J. Therm. Anal., 16(1),

151-162 (1979) (Thermodyn., 16)

[1980Her] Herley, P.J., Christofferson, O., Todd, J.A., “Microscopic Observations on the Thermal

Decomposition of -Aluminum Hydride”, J. Solid State Chem., 35, 391-401 (1980)

(Experimental, 15)

[1980Min] Mintz, M.H., Gavra, Z., Kimmel,G., “The Reaction of Hydrogen with Magnesium Alloys

and Magnesium Intermetallic Compounds”, J. Less-Common Met., 74, 263-270 (1980)


[1981Gav] Gavra, Z., Hadari, Z., Mintz, M.H., “Effects of Nickel and Indium Ternary Additions on the

Hydrogenations of Mg-Al Intermetallic Compounds”, J. Inorg. Nucl. Chem., 43, 1763-1768

(1981) (Thermodyn., Review, 11)

[1981Tuc] Tuck, C.D.S., “Evidence for the Formation of Magnesium Hydride on the Grain Boundaries

of Al-Mg and Al-Zn-Mg Alloys During their Exposure to Water Vapour”, in “Hydrogen Eff.

Met. ”, Proc. 3rd Int. Conf., 1980 (Publ. 1981), 503-511 Bernstein I.M., Thompson, A.V.,

(Eds.), Metall. Soc. AIME, Warrendale, USA, (1981) (Experimental, 23)

[1982Mur] Murray, J.L., “The Al-Mg (Aluminum-Magnesium) System”, Bull. Alloy Phase Diagrams,

3, 60-74 (1982) (Review, Equi. Diagram, Thermodyn., 112)

[1983Sem] Semenenko, K.N., Verbettskii, V.N., Kotchukov, A.V., Sytnikov, A.N., “Reaction of

Magnesium Containing Intermetallic Compounds and Alloys with Hydrogen” (in Russian),

Vestn. Mosk. Uni., Ser. 2: Khim., 24(1), 16-27 (1983) (Thermodyn., Review, 46)

[1984Lue] Luedecke, C.M., Deublein, G., Huggins, R.A., “Use of Electrochemical Methods to Study

and Control Hydrogen Storage in Solid Metal Hydrides”, Adv. Hydrogen Energy, 4,

(Hydrogen Energ. Prog. 5, Vol. 3) 1421-1431 (1984) (Equi. Diagram, Thermodyn,

Experimental, #)

[1985Lue1] Luedecke, C.M., Deublein, G., Huggins, R.A., “Electrochemical Investigation of Hydrogen

Storage in Metal Hydrides”, J. Electrochem. Soc.: Electrochem. Sci. Techn., 132(1), 52-56


[1985Lue2] Luedecke, C.M., Deublein, G., Huggins, R.A., “Investigation of Metal Hydrides with

Thermodynamic Calculations and Electrochemical Experiments”, Hydrogen Syst. Pap. Int.

Symp Meeting Date, 1, 363-377 (1985) (Equi. Diagram, Thermodyn., Experimental, #, 18)

68


MSIT®

Al–H–Mg

[1987Lue] Luedecki, C.M., Deubleiun,G., Huggins, R.A., “Thermodynamic Characterization of Metal

Hydrogen Systems by Assessment of Phase Diagrams and Electrochemical Measurements”,

Int. J. Hydrogen Energy, 12(2) 81-88 (1987) (Equi. Diagram, Thermodyn., Review, #, 18)

[1989Lin] Lin, R.J., Hoch, M., “The Solubility of Hydrogen in Molten Aluminium Alloys”, Metall.

Trans. A, 20(9), 1785-1791 (1989) (Theory, Thermodyn., 31)

[1992Kon] Konovalov, S.K., Bulchev, B.M., “High Pressures in the Chemistry of Beryllium and

Aluminium Hydrides”, Russ. J. Inorg. Chem., 37(12), 1361-1365 (1992), translated from

Zh. Neorg. Khim., 37, 2640-2646 (1992) (Equi. Diagram, Experimental, 16)

[1992San] San Martin, A., Manchester, F.D., “The Al-H (Aluminum-Hydrogen) System”, J. Phase

Equilib., 13(1), 17-21 (1992) (Equi. Diagram, Review, 45)

[1998Lia] Liang, P., Su, H.-L., Donnadieu, P., Harmelin, M.G., Quivy, A., Ochin, P., Effenberg, G.,

Seifert, H.J., Lukas, H.L., Aldinger, F., “Experimental Investigation and Thermodynamic

Calculation of the Central Part of the Mg-Al Phase Diagram”, Z. Metallkd., 89(8), 536-540

(1998) (Experimental, Assessment, Calculation, Equi. Diagram, Thermodyn., 33)

[2000Bou] Bouaricha, S., Dodelet, J.P., Guay, D., Huot, J., Boily, S., Schulz, R., “Hydriding Behavior

of Mg-Al and Leached Mg-Al Compounds Prepared by High-Energy Ball-Milling”,

J. Alloys Compd., 297, 282-293 (2000) (Equi. Diagram, Crys. Structure, Experimental, 27)

[2001Per] Perrot, P., Schmid-Fetzer, R., “Hydrogen-Magnesium”, in “Ternary Alloys: A

Comprehensive Compendium of Evaluated Consitutional Data and Phase Diagrams”,

Effenberg, G., Aldinger, F., Rogl, P. (Eds.), Vol. 18, MSI, Materials Science International

Services GmbH, Stuttgart, 3-4 (2001) (Thermodyn., Assessment, Equi. Diagram, #, 6)

[2002Her] Herbst, J.F., “On Extending Miedema’s Model to Predict Hydrogen Content in Binary and

Ternary Hydrides”, J. Alloys Compd., 337, 99-107 (2002) (Calculation, Thermodyn., 20)



Stuttgart; Document ID: 20.14832.1.20, (2002) (Equi. Diagram, Crys. Structure,

Assessment, 21)

[2003Luk] Lukas, H.-L., Lebrun, N., “Al-Mg (Aluminium-Magnesium)”, MSIT Binary Evaluation


Services, GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram,

Assessment, 49)


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

( Al) hP2

P63/mmc

Mg

a = 269.3

c = 439.8

at 25°C, 20.5 GPa [Mas2]

( Al)

< 660.452

cF4

Fm3m

Cu

a = 404.96 at 25°C [Mas2]

100 to 81.4 at.% Al at 450°C [1982Mur]

(Mg)

< 650

hP2

P63/mmc

Mg

a = 320.94

c = 521.07

at 25°C [Mas2]

0 to 11.5 at.% Al at 437°C [1982Mur]

, Mg17Al12

458

cI58

I43m

Mn

a = 1054.38 at 41.4 at.% Al [V-C2]

39.5 to 51.5 at.% Al [1998Lia]

40 to 52 at.% Al [2003Luk]

69


MSIT®

Al–H–Mg

, Mg2Al3 452

cF1168

Fd3m

Mg2Al3

a = 2816 to 2824 60-62 at.% Al [2003Luk]

1168 atoms on 1704 sites per unit cell

[2003Luk]

, Mg23Al30

410 - 250

hR159

R3

Mn44Si9

a = 1282.54

c = 2174.78

54.5-56.5 at.% Al [2003Luk] Structure:

159 atoms refer to hexagonal unit cell

[2003Luk]

AlH3

< 110

hR24

R3c

a = 445.6

c = 1183

[1992San], metastable

AlH3

< 80

- - Metastable “Aluminum

hydrogenoaluminate”

Al(AlH4)3 [1978Cla2]

MgH2 tP6

P42/mnm

TiO2

a = 451.68

c = 302.05

[P]

* Mg(AlH4)2 - - [1978Cla1], stable above 5 MPa H2 at

410°C [1980Min]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

20

40

60

80

20 40 60 80

20

40

60

80

Mg Al

H Data / Grid: at.%

Axes: at.%

(Mg) γ β (αAl)

MgH2

(Mg)+γ+MgH2

β+γ+MgH

2

(αAl)+β+MgH2

AlH3

MgAl2H

8

Fig. 1: Al-H-Mg.


temperatures between

140 and 170°C

70


MSIT®

Al–H–Mg

20 40 60 80

2

4

6

8

10

12

(H,at.

%)10

-3�

Mg, at.% MgAl

800°

C

500°C

700°C

100

0

0

10

1

P(M

Pa

)H

2

0.5 1.0 1.5 2.0

H/Mg

335°C

350°C

375°C

410°C

00

Fig. 3: Al-H-Mg.

Pressure-composition

isotherms of the

Mg2Al-H system

Fig. 2: Al-H-Mg.

Hydrogen solubility

in liquid Al-Mg alloys

under 1 bar at 500,

700 and 800°C

71


MSIT®

Al–H–Ti

Aluminium – Hydrogen – Titanium

Viktor Kuznetsov

Literature Data

The major works on this system has been concentrating on H in Ti-rich phases to investigate H

embrittlement and related phenomena and using Al-Ti alloys for hydrogen storage. [1981Ive] mentioned

Ti3Al as one of most promising candidate systems for hydrogen storage. Unfortunately the equilibrium

usually was obtained only between H2 gas and metal surface, if at all, but not within the metal sublattice,

which corresponds to paraequilibrium conditions. True phase equilibria were achieved and investigated

very rarely and the information on them remains very limited.

[1958Ber] studied by metallography the H embrittlement of Ti and alloys with 2.5, 5 and 7 mass% Al

prepared between 675 and 940°C and concluded that Al increases the H solubility. Later [1971Pat]

re-investigated this using alloys of iodide-purified Ti with 1, 3 and 10 at.% Al. Using resistometrical

methods and direct observation of hydride formation by electron microscopy, he showed the increase of H

solubility to be due to self-stresses around the hydride particles; plastic flow of the matrix causes a strong

hysteresis. This hysteresis state is rather stable and the apparent equilibrium is not disturbed for several

weeks from 20 to approximately 150°C. This was confirmed by [1976Che] who showed that such

supersaturated solutions of H in Ti-4Al alloys do decompose, giving TiH2 after annealing for 40 d under

stress conditions.

[1974Sch1, 1974Sch2] investigated in great detail the solubility of H in Ti and its alloys with 5, 7 and 10

at.% Al from 800 to 900°C and described the coexisting phase configurations in three partial isothermal

sections. The main impurities were up to 0.03 mass% Fe, 0.04% C and 0.4% O. Analogous work was

conducted by [1981Buk] from 500 to 800°C, but the results at 800°C agree rather poorly.

[1981Buk] also displayed the position of three-phase triangles - 2- . The claim of the three-phase state of

the products of hydridation of alloys with Al content from 7.5 up to 18.4 mass% based on metallography

are corroborated to some degree by the observations which [1989Ili] made on the formation of the 2 phase

in hydridated alloys with more than 7 mass% Al.

[1972Gab] measured a H2 solubility in TiAl3 at 500 and 600°C extrapolating data which were obtained for

H2 pressures of 0.4 to 0.6 kbar to a H2 pressure of 1.01 bar. Only the solubilities at 1.01 bar were given. For

both temperatures the solubilities were found to be 1.4 to 1.6 ml H2 per 100 g of alloy. [1977Rud]

investigated the solubility and the thermodynamics of solution of H in Ti3Al from 450 to 800°C and for H2

pressures lower than 1.333 bar. The hydrogen solubilities at room temperature under hydrogen pressure of

5 MPa were measured by [2001Has, 2002Has, 2002Ito] around the composition Ti3Al, as well as the

temperature at which 50 % hydrogen is desorbed in the whole interval of compositions.

[1972Sch] studied the influence of the temperature on the rate of thermal decomposition of hydridation

products for Ti alloys with 1.2, 3.0 and 5.9 mass% Al. [2000Sor2] studied the temperatures and details of

kinetics of thermal decomposition of two hydrides, obtained by hydridation of Ti3Al under H2 pressure of

3.8 MPa at room temperature. [2002Ito] studied hydrogen adsorption isotherms for another hydridation

conditions (127°C, 0.001 to 10 MPa).

[1978Rud] studied the interactions of Ti3Al with H2 at higher H2 pressures than [1977Rud] but using the

same specimen and found three metastable phases analogous to hydrides of Ti. The phases exist up to 150°C

and decompose at 200°C giving TiH2. Based on metallographic investigations [1975Buk] suggested the

existence of a hydride distinguishable from TiH2 after slow cooling from 800°C to room temperature, the

Al content being more than 4 to 5 at.%. [1981Kol] confirmed this by X-ray studies. A fragment of the

diffraction pattern of the two-phase mixture of TiH2 and a new phase “TiAlHx” is given, but no structural

data were extracted; the real composition of that phase is also not known. [1991Sch] obtained a Ti3AlH

compound by a reaction of Ti3Al with H2 gas at pressure of 0.1 MPa and a temperature of 600°C, and

determined its structure using neutron diffraction. [1999Mae] used the same technique on a product of

interaction of Ti3Al with D2 gas at p = 9.2 bar and 200°C; they determined the composition of a higher

72


MSIT®

Al–H–Ti

hydride to be Ti3AlH8-z (z 0.8) and determined its crystal structure. The latter phase was identified with

fcc hydride of [1978Rud], but no bcc phase was detected under that conditions. A structural study of the

reaction products of Ti3Al with H2 gas at 127°C was performed by [2002Ito], who found and investigated

by X-ray diffraction and electronography both phases discovered by [1978Rud]. The metal sublattice of

“bcc” hydride called as “ H” hydride proved to be an orthorhombic superlattice to bcc. Two modifications

of “fcc” phase of [1978Rud] called “ H1” and “ H2” were discovered; the metal sublattices of both have

bcc superlattices, close to fcc. The positions of H atoms were not determined but the H2 phase was

identified as Ti3AlH8-z [1999Mae]. The decomposition of higher hydrides gives TiH2 and some Al

enriched product. For alloys with 30 % Al and more [1999Mae] found amorphization under H2 treatment.

The results of [1999Mae] and [2002Ito] generally confirm those of [1978Rud] and refine the structural data.

[2002Ito] also suggested a possible mechanism of mutual transformation of these phases. These authors

correlated relative stability of different hydride phases with the kinetics of hydrogen desorption.

The H solubility in two-phase samples (Ti3Al+TiAl) from 450 to 570°C was measured by [1995Tak].

[1976Gri] investigated the solubility and thermodynamics of solution of H in liquid Al-Ti alloys up to

8.7 mass% Al between 1700 and 2100°C. The starting materials were Al(A999) and Ti sponge with main

impurities of 0.04 mass% Fe, 0.01% Mn, 0.002% Si, 0.004% C, 0.04% O and 0.01% N. The specimens

obtained were analyzed yielding 0.03 to 0.4 mass% O and 0.01 mass% N. To prevent contamination, the H

saturation was conducted by electromagnetic levitation with subsequent quenching. H content was

measured by vacuum extraction.

The solubility of H in Ti3Al is theoretically analyzed using a geometrical model [1985Mro]. [1994Bel]

performed investigation of H influence on ordering in the Ti3Al phase using GBW model.

Ab initio calculation of electronic structure, chemical bonding and hydrogen site preferences in two

modifications of Ti3Al and Ti3AlH phase was performed by [2000Sor1].

Binary Systems

For the Al-H and Al-Ti binary systems the updated versions [2002Per, 2003Sch] are accepted. The H-Ti

edge is believed to be correct as described by [Mas2].

Solid Phases

The ternary hydride phases are stable under hydrogen pressure. For instance, at 127°C, Ti3AlH8-z is stable

above 10 kPa [2002Ito], and their appearance strongly depends on the conditions of preparation.

The crystallographic data for all reported phases, including metastable ternary hydrides, are given in

Table 1. For the H phase only the structure of the metal sublattice is known. The H1 and H2 phases are

claimed to differ only by H content [2002Ito]; no direct structural data for the former seem to exist. The H

composition in H1 is not reported; it was estimated by the present author from its position in hydridation

sequence after H and claims of [2002Ito] that it contains less hydrogen than H2. The identity of the latter

with fcc hydrides of [1978Rud] and [1999Mae] are accepted, though the real structure may be more

complex than determined by [1999Mae].

Isothermal Sections

Only [1981Buk] and [1974Sch1, 1974Sch2] tried to present true phase equilibria. The latter data are

preferred, mainly because in the former work the H2 pressure in declared three-phase field was not constant.

Figures 1 and 2 display the sections at 800 and 900°C after [1974Sch1, 1974Sch2]; in addition these authors

give an isothermal section at 850°C which is very similar to that at 800°C and not reproduced here.

Thermodynamics

A selection of isoactivity lines of H [1974Sch2] at 800 and 900°C is presented in Figs. 1 and 2.

A Wagner expansion of the activity coefficient [1976Gri] fits the experimental data within their scatter

between 1700 and 2100°C up to 10% Al:

log10 ((%H)/p(H2)1/2) = 2323/T - 2.043 - (92.2/T - 0.03) (%Al)

73


MSIT®

Al–H–Ti

where (%H), (%Al) are in mass%, p(H2) in bar, T in K. The interaction parameters of Al in Ti and Ti

have been calculated by [1974Sch2]:

in Ti: H(Al) = dln H / dx(Al) = +5.51

in Ti: H(Al) = dln H / dx(Al) = +6.04.

The positive values of show that Al dissolved in Ti decreases the hydrogen solubility. The result was

experimentally confirmed by [1975Buk] with solid Ti, then by [1976Gri] with liquid Ti.

The solubility of H in Ti3Al has been investigated at various temperatures and pressures, up to 200°C and

up to 10 MPa [1978Rud]. The 150°C isotherm presents a plateau from H-Ti3AlH2 to H1-Ti3AlH3 under

1 MPa.

The hydrogen uptake goes up to Ti3AlH4 under 10 MPa H2. The same plateau is estimated under 0.1 MPa

at 100°C and under 0.01 MPa at 50°C. The hydrogen pressure at equilibrium H- H1 is given by:

RTln(pH2/bar) = -47280 + 127.2T

This relation agrees with measurements made later by [2002Ito] which propose a plateau at 127°C and 0.2

MPa.

The solubility of H under 1 bar in TiAl has been experimentally determined between 450 and 570°C

[1995Tak]. It is given by the following expressions:

for Ti50Al50 c/ppm = 1.12 104exp(-4380/T),

for Ti55Al40 c/ppm = 1.53 106exp(-7010/T).

These alloys show endothermic uptake of hydrogen. Only the Ti47Al53 alloy takes up hydrogen

exothermically.


The use of Ti3Al for hydrogen storage is discussed from technical point of view in [1981Ive], [1995Tak].

At room temperature under 1 MPa H2 Ti3Al may absorb hydrogen up to the composition Ti3AlH5.6

(H/Me=1.4), under 5 MPa the hydride obtained is Ti3AlH6 (H/Me = 1.5). The hydrogen capacity decreases

with off-stoichiometry. For instance, under 5 MPa H2, Ti0.7Al0.3 alloy absorbs hydrogen up to the

composition Ti0.7Al0.3H. The desorption of hydrogen reaches 50% by heating at 600°C; it reaches 100%

by heating at 800°C [2001Has]. Careful investigation has been carried out by [2002Ito] at higher

temperature (127°C). The pressure composition curve of Fig. 3 shows a plateau with hysteresis. On the

absorbing edge a plateau is observed at 0.2 MPa for the transition Ti3AlH2 (H/Me=0.5) to Ti3AlH4

(H/Me=1). On the desorbing edge the plateau (narrower and less well defined) is observed at 8 kPa. The

position of the plateaus does not change significantly with the preparation of the samples (single crystalline,

homogenized, pulverized and as arc-melted).

Miscellaneous

[1954Ram] suggested as a preparative method to obtain Ti hydride with low O and N content the saturation

of an Al-10Ti (mass%) alloy with H2 at 1000°C. This suggestion, however, seems to contradict all other

data, especially [1972Gab] who did not find any decomposition of TiAl3 with H2 up to 0.4 to 0.6 kbar of

the latter. It may be correlated to some degree with [1972Sch] and [1978Rud], although an observation of

[1972Sch] of the decomposition of alloys with only 1 to 6 mass% Al to give free Al (!) seems to be quite

surprising. Nevertheless the possible formation of pure Al during decomposition of H2 phase was

discussed by both [1999Mae] and [2002Ito], although none of authors could detect it. The suggestion of

[1999Mae] on the formation of TiAl (possibly in nanocrystalline or amorphous state) in addition to TiH2

under that conditions seems to be more realistic.

The substitution of Ti by Zr or Hf decreases slightly the hydrogen storage capacity of Ti3Al; the substitution

of one atom of Ti in Ti3Al by one atom of Mn, Ni, Cu, V or Co decreases hydrogen storage capacity by a

factor of 3. The alloys Ti2CrAl or Ti2FeAl had no hydrogen storage capacity at all [2001Ish].

74


MSIT®

Al–H–Ti

References

[1954Ram] Ramamurthi, S., “Formation of Titanium Hydride in Aluminium- Titanium Alloys”, J. Sci.

Ind. Research (India), 13B, 306-307 (1954) (Experimental, 3)

[1958Ber] Berger, L.W., Williams, D.H., Jaffe, R.J., “Hydrogen in Titanium-Aluminium Alloys”,

Trans. Met. Soc. AIME, 212, 509-513 (1958) (Experimental, 7)

[1971Pat] Paton, N.E., Hickman, B.S., Leslie, D.H., “Behavior of Hydrogen in a Phase Ti-Al Alloys”,

Metall. Trans., 2, 2791-2796 (1971) (Experimental, *, 16)

[1972Gab] Gabidullin, R.M., Shevtsov, I.N., Kolachev, B.A., Archakov, Yu.I., “Solubility of H in Al

Intermetallics with Mg, Cu, Mn, Ti and Zr” (in Russian), Stroenie Svoistva i Primenenie

Metall., (Publ. 1974), 188-190 (1972) (Experimental, 2)

[1972Sch] Schekhotsov, M.G., Kolomytsky, F.M., Rubtsov, A.N., “Investigation of Thermal Stability

of Titanium Hydride and Hydridated Titanium Based Alloys” (in Russian), Stroenie

Svoistva i Primenenie Metall., (Publ. 1974), 185-188 (1972) (Experimental, 4)

[1974Sch1] Schuermann, E., Kootz, T., Preisendranz, H., Schueller, P., Kauder, G., “On the Hydrogen

Solubility in the Ti-Al-H, Ti-V-H and Ti-Al-V-H in the Temperature Range 800 to 1000°C

at H2 Pressures 0.1 to 250 mbar. Part 1: Theoretical Basis and Experimental Data” (in

German), Z. Metallkd., 65, 167-172 (1974) (Experimental, Thermodyn., *, 32)

[1974Sch2] Schuermann, E., Kootz, T., Preisendranz, H., Schueller, P., Kauder, G., “On the Hydrogen

Solubility in the Ti-Al-H, Ti-V-H and Ti-Al-V-H in the Temperature Range 800 to 1000°C

at H2 Pressures 0.1 to 250 mbar. Part 2: Thermodynamic Evaluation” (in German),

Z. Metallkd., 65, 249-255 (1974) (Equi. Diagram, Thermodyn., #, *, 3)

[1975Buk] Bukhanova, A.A., Kolachev, B.A., Nazimov, O.Z., Seregina, E.V., “On the Influence of Al

to H Solubility in Ti” (in Russian), Tekhnol. Legk. Splavov, (8), 48-53 (1975)

(Experimental, 7)

[1976Che] Chernetsov, V.I., Tseiger, E.N., “On the Solubility of H in Aluminium-Bearing Titanium

Alloys”, Sov. J. Non-Ferrous Met., (5), 69 (1976), translated from Tsvetn. Met., (5), 67


[1976Gri] Grigorenko, G.M., Lakomskii, V.I., Korzhov, M.P., Tetyukhin, V.V., Konstantoniv, V.S.,

Kalinyuk, N.M., Gontchar, V.Ya., Solomentsev, A.N., “The Influence of Al to H Activity

in Molten Ti” (in Russian), Probl. Spets. Elektrometall., (5), 88-93 (1976) (Thermodyn.,

Experimental, 12)

[1977Rud] Rudman, P.S., Reilly, J.J., Wiswall, R.H., “Hydrogen Absorption in Ti3Al”, Ber.

Bunsen-Ges. Phys. Chem., 31, 71-80 (1977) (Experimental, 10)

[1978Rud] Rudman, P.S., Reilly, J.J., Wiswall, R.H., “The Formation of Metastable Hydrides

Ti0.75Al0.25Hx with x < 1.5”, J. Less-Common Met., 58, 231-240 (1978) (Experimental,

Crys. Structure, 10)

[1981Buk] Bukhanova, A.A., Kolachev, “On the Phase Diagram of the Ti-Al-H System between 500

to 800°C” (in Russian), Fazovje Ravnovesija v Metallicheskych Splavach, Publ. Nauka,

Moscow, 127-131 (1981) (Equi. Diagram, 3)

[1981Ive] Ivey, D.G., Northwood, D.O., “Metal Hydrides for Energy Storage”, Can. Metall. Quart.,

20, 397-405 (1981) (Review, 40)

[1981Kol] Kolachev, B.A., Gontchar, V.Ya., Liskovitsch, V.A., “Phase Composition of the

Hydrogenation Products of Titanium Alloys”, Inorg. Mater., 17, 1527-1530 (1982),

translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 17, 2048-2052 (1981)

(Experimental, 10)

[1985Mro] Mrowietz, M., Weiss, A., “Solubility of Hydrogen in Titanium Alloys. II. Blocking Models

and Hole Size Considerations”, Ber. Bunsen-Ges. Phys. Chem., 89, 362-371 (1985)

(Thermodyn., Theory, 82)

[1987Ere] Eremenko, V.N., Tretyachenko, L.A., “Physico-Chemical Properties of Titanium”, in

“Ternary Systems of Titanium with Transition Metals of IV-VI Groups” (in Russian),

Naukova Dumka, Kiev, 5-6 (1987) (Equi. Diagram, Crys. Structure, Review, 14)

75


MSIT®

Al–H–Ti

[1989Ili] Il'in, A.A., Mamonov, A.M., Mikhailov, Yu.V., “The Phase Diagrams of H Alloyed Ti

Alloys” (in Russian), Abstr. 5th All-Union Conf. Phase Diagrams of Metallic Systems, 162

(1989) (Equi. Diagram, Abstract, 0)

[1990Sch] Schuster, J.C., Ipser, H., “Phases and Phase Relations in the Partial System TiAl3-TiAl”,

Z. Metallkd., 81, 389-396 (1990) (Crys. Structure, Equi. Diagram, Experimental,

Review, #, 33)

[1991Sch] Schwartz, D.S., Yelon, W.B., Berliner R.B., Lederich, R.J., Sastry, S.M., “A Novel Hydride

Phase in Hydrogen Charged Ti3Al”, Acta Met. Mater., 39, 2799-2803 (1991) (Crys.

Structure, Experimental, *, 8)

[1992Kat] Kattner, U.R., Lin, J.-C., Chang, Y.A., “Thermodynamic Assessment and Calculation of the

Ti-Al System”, Metall. Trans. A, 23(8), 2081-2090 (1992) (Assessment, Calculation, Equi.

Diagram, Thermodyn., #, *, 51)

[1994Bel] Belov, S.P., Il'in, A.A., Mamonov, A.M., Aleksandrova, A.V., “Theoretical Analysis of

Ordering in Ti3Al-Base Alloys. II. Effect of Hydrogen on Stability of Ti3Al Intermetallic

Compound”, Russ. Metall., (2), 52-55 (1994), translated from Izv. Ross. Akad. Nauk. Met.,

(2), 76-78 (1994) (Crys. Structure, Theory, 13)

[1995Tak] Takasaki, A., Furuya, Y., Ojima, K., Taneda, Y., “Hydrogen Solubility of Two-Phase

(Ti3Al+TiAl) Titanium Aluminides”, Scr. Metall. Mater., 32, 1759-1764 (1995) (Phys.


[1999Mae] Maeland, A.J., Hauback, B., Fjellvag, H., Sorby, M., “The Structure of Hydride Phases in

the Ti3Al/H System”, Int. J. Hydrogen Energy, 24, 163-168 (1999) (Crys. Structure,

Experimental, *, 12)

[2000Sor1] Sornadurai, D., Panigrahi, B., Ramani, “Electronic Structure, Hydrogen Site Occupation

and Phase Stability of Ti3Al upon Hydrogenation”, J. Alloys Compd., 305, 35-42 (2000)

(Crys. Structure, Theory, 22)

[2000Sor2] Sornadurai, D., Panigrahi, B.K., Shashikala, K., Raj, P., Sastry, V.S., Ramani, “X-Ray

Diffraction and Differential Scanning Calorimetry Investigations on High-Pressure

Hydrogen Gas Charged Ti3Al”, J. Alloys Compd., 312, 251-256 (2000) (Crys. Structure,

Kinetics, *, 10)

[2001Bra] Braun, J., Ellner, M., “Phase Equilibria Investigations on the Aluminium-Rich Part of the

Binary System Ti-Al”, Metall. Mater. Trans. A, 32A, 1037-1048 (2001) (Crys. Structure,

Equi. Diagram, Experimental, #, *, 34)

[2001Has] Hashi, K., Ishikawa, K., Aoki, K., “Hydrogen Absorption and Desorption in Ti-Al Alloys”,

Met. Mater. Int., 7(2), 175-179 (2001) (Equi. Diagram, Experimental, 8)

[2001Ish] Ishikawa, K., Hashi, K., Suzuki, K., Aoki, K., “Effect of Substitutional Elements on the

Hydrogen Absorption-Desorption Properties of Ti3Al Compounds”, J. Alloys Compd., 314,

257-261 (2001) (Crys. Structure, Kinetics, Experimental, *, 8)

[2002Ito] Ito, K., Okabe, Y., Zhang, L.T., Yamaguchi, M., “Reversible Hydrogen

Absorption/Desorbtion and Related Phase Transformations in a Ti3Al Alloy with

Stoichiometry Composition”, Acta Mater., 50, 4901-4912 (2002) (Equi. Diagram,


[2002Has] Hashi, K., Ishikawa, K., Syzuki, K., Aoki, K., “Hydrogen Absorption and Desorption in the

Binary Ti-Al System”, J. Alloys Compd., 330/332, 547-550 (2002) (Equi. Diagram,

Experimental, 11)




Assessment, 21)

[2003Sch] Schmid-Fetzer, R., “Al-Ti (Aluminium-Titanium)”, MSIT Binary Evaluation Program, in

MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services,

GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 86)

76


MSIT®

Al–H–Ti


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Al)

< 660.452

cF4

Fm3m

Cu

a = 404.88 pure Al [V-C]

( Ti)

1670 - 882

cI2

Im3m

W

a = 330.65

a = 328.4

pure Ti at 900°C [V-C]

at room temperature, extr.

from solid solution [1987Ere]

( Ti)

< 882

hP2

P63/mmc

Mg

a = 295.2

c = 498.9

[V-C]

TiAl3< 1387

tI8

I4/mmm

TiAl3

a = 384.88

c = 859.82

[1990Sch]

“Ti2Al5”

1416 - 990

tetragonal

superstructure of

AuCu-type

[2001Bra]

tP28

P4/mmm

“Ti2Al5”

a* = 395.3

c* = 410.4

a* = 391.8

c* = 415.4

a = 390.53

c = 2919.63

chosen stoichiometry [1992Kat]

summarizing several phases:

Ti5Al11 stable range 1416- 995°C

[2001Bra]

66 to 71 at.% Al at 1300°C [2001Bra]

(including the stoichiometry Ti2Al5!);

[1990Sch] claimed: 68.5 to 70.9 at.% Al

and range 1416 -1206°C;

at 66 at.% Al [2001Bra]

* AuCu subcell only

at 71 at.% Al [2001Bra]

* AuCu subcell only

“Ti2Al5”

~1215 - 985°C [1990Sch];

included in homogeneity region of

Ti5Al11 [2001Bra]

Ti5Al111416 - 1206

tI16

I4/mmm

ZrAl3

a = 392.30 to

393.81

c = 1653.49 to

1649.69

29.1 to 31.5 at.% Ti [1990Sch]

TiAl2(h)

1433 - 1214

oC12

Cmmm

ZrGa2

a = 1208.84

b = 394.61

c = 402.95

33 to 34 at.% Ti [1990Sch]

TiAl2(r)

< 1216

tI24

I41/amd

HfGa2

a = 396.7

c = 2429.68

[1990Sch]

Ti1-xAl1+x

~1445 - 1424

oP4 a = 402.62

b = 396.17

c = 402.62

at x = 0.28

[1990Sch]

TiAl

< 1460

tP4

P4/mmm

AuCu(I)

a = 398.69

c = 405.39

at 38.5 to 52 at.% Ti [1990Sch]

at 38.5 at.% Ti, 1000°C

77


MSIT®

Al–H–Ti

a) Only metal atoms are counted for Pearson symbol

Ti3Al

< 1180

hP8

P63/mmc

Ni3Sn

a = 580.6

c = 465.5

a = 574.6

c = 462.4

at 78 at.% Ti [L-B]

at 62 at.% Ti [L-B]

TiHx> 315

cF12

Fm3m

CaF2

a = 445.4 x = 1.05 to 2.0 [Mas, V-C]

TiHx< 315

tI6

I4/mmm

ThH2

a = 320.2

c = 427.9

x = 1.72 to 2.0 [Mas, V-C]

* Ti3AlH cP5

Pm3m

CaTiO3

a = 408.79 [1991Sch]

The parameter is given by [1999Mae]

for Ti3AlD

* Ti0.75Al0.25Hx hP?

a = 289

c = 466

metastable, x < 0.2 [1978Rud] decomp.

at 200°C, form. at 50 to 150°C

at x 0

* Ti0.75Al0.25Hx cI?

a = 328

metastable, 0.4 < x < 0.5 (in-reactor

state) [1978Rud] decomp. at 200°C,

form. at 50 to 150°C at x = 0.35

(estimated)

* Ti0.75Al0.25Hx cF?

a = 435

metastable, x > 1.5 [1978Rud] decomp.

at 200°C, form. at 50 to 150°C

at x = 1.6 (two-phase sample, estimated

comp.)

* H, Ti3AlH2 oP8 a)

a 310

[2002Ito]; two-phase sample with

H/Me = 0.55; 2×2×1 superstructure to

bcc; 2c/a ratio close to 1, varies

depending on sample probably identical

to cI? phase of [1978Rud] approx.

Value for bcc sublattice

* H1, Ti3AlHx t?? a = 390

c = 313

[2002Ito]; two-phase sample with

H/Me = 0.55 (same as previous);

x probably between 2 and 7.8

* H2, Ti3AlH8-z tP12 a = 439.77 [1999Mae]; sample of gross

composition Ti3AlH5.9 contained also

6-7% of TiAl and Ti3AlH;

z 0.8; composition of metal sublattice

is Ti3(Al0.25Ti0.75)

after [2002Ito] may be bct with c/a ratio

close to fcc

* “TiAlyHx” - - [1981Kol], the composition of both Al

and H is not known

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

78


MSIT®

Al–H–Ti

80

90

10 20

10

20

Ti Ti 70.00Al 30.00H 0.00

Ti 70.00Al 0.00H 30.00 Data / Grid: at.%

Axes: at.%

0.018

0.073

0.127

0.1

81

0.2

36

0.327

0.508

(αTi)

(βTi)

(αTi)+(βTi)

Fig. 1: Al-H-Ti.

Partial isothermal section

with superimposed

(dashed) isoactivity lines

(aH) at 800°C.

The numbers given are

aH=(pH2(bar)/0.981)1/2)

80

90

10 20

10

20

Ti Ti 70.00Al 30.00H 0.00

Ti 70.00Al 0.00H 30.00 Data / Grid: at.%

Axes: at.%

(βTi)

(αTi)

0.726

0.508

0.327

0.2180.163

0.109

0.0180.073

Fig. 2: Al-H-Ti.

Partial isothermal section

with superimposed

(dashed) isoactivity lines

(aH) at 900°C.

The numbers given are

aH=(pH2(bar)/0.981)1/2)

79


MSIT®

Al–H–Ti

Hydrogen content (mass%)

Pre

ssure

(MP

a)

10

1

0.1

0.01

0.003

Hydrogen content (H/Me)

0.6 0.8 1.0 1.2

1.5 2.0 2.5 3.01.0

(A)

(B)

Fig. 3: Al-H-Ti.

Absorbtion (A) and

desorption (B)

isotherms at 127°C

for Ti3Al

80


MSIT®

Al–Hf–Ni

Aluminium – Hafnium – Nickel

Gautam Ghosh

Literature Data

[1969Mar1] was the first to report the isothermal section of the entire system at 800°C. They prepared about

100 ternary alloys in an arc furnace under Ar atmosphere using the elemental metals Al (99.998 mass%),

iodide Hf (99.95 mass%) and electrolytic Ni (99.9 mass%). The alloys were annealed at 800°C for 830 h in

evacuated silica tubes followed by quenching into cold water. Phase analysis was performed by

microstructural observation and X-ray diffraction techniques. [1981Nas] reported the partial isothermal

sections of the Ni corner at 1200 and 1000°C. They prepared 12 ternary alloys containing up to 35 at.% Al

and 23 at.% Hf. The alloys were prepared from 99.99 mass% Al, Hf containing about 3 at.% Zr and other

impurities of about 0.38 mass%, and 99.99 mass% Ni. The alloy buttons were prepared in an arc furnace

under Ar atmosphere. They were placed in alumina crucibles, sealed in silica tubes partially filled with Ar

and were homogenized at 1200 and 1000°C for 168 h followed by quenching into water. Phase analysis was

carried out by optical microscopy, X-ray diffraction and electron probe microanalysis. [1981Bal]

investigated microstructure of two Ni rich ternary alloys, both as-cast and annealed conditions. These

results were reviewed by [1991Lee, 1993Gho]. Brief reviews of phase equilibria were presented by

[1977Abr, 1990Kum].

Recently, Miura et al. [1999Miu] investigated the solid-liquid phase equilibria of Ni-rich ternary alloys

using DTA, XRD and SEM-WDS analysis. They prepared ternary alloys using 99.99 mass% Al, 99.95

mass% Ni, and 95 mass% Hf. [1991Mis] determined the solvus boundary of (Ni) using DTA and SEM-

EDX analysis. Other recent investigations of the ternary system involve rapid solidification [2002Lou], and

very limited thermodynamic measurements [1992Alb].

Binary Systems

The Al-Ni binary phase diagram is accepted from [2003Sal], and the Al-Hf binary phase diagram is

accepted from [2003Sch].

Recently, Miura et al. [1999Miu, 2001Miu] have determined the liquidus of Ni rich alloys containing up to

13 at.% Al. Unlike Hilpert et al. [1987Hil], Miura et al. [2001Miu] observed a maximum (1466°C) in the

liquidus at about 2 at.% Al. Except for [2001Miu], this feature has not been considered in the CALPHAD

modeling of the Al-Ni phase diagram [2003Sal].

Unlike [1998Mur], Schuster [2003Sch] did not consider Hf2Al phase in the assessment of Al-Hf

equilibrium diagram. This phase was first reported by [1961Now] but subsequent investigations failed to

confirm. It is believed that Hf2Al and HfAl3(TiAl3) might have been stabilized by silicon, and they are not

a Al-Hf equilibrium phase [1962Poe1, 1962Poe2, 1964Rie].

The Hf-Ni binary phase diagram is accepted from [1983Nas].

Solid Phases

The data of [1981Nas] suggest that the lattice parameter of (Ni) increases more rapidly in the ternary regime

than in the binary solid solutions [1985Mis]. In the Hf-Ni system, the rate of increase in the lattice

parameter, da/dc, is reported to be 1.0 pm/at.% Hf [1984Och2, 1985Mis]. Figure 1 shows the solubility

isotherms of (Ni) [1991Mis].

Ni3Al is reported to dissolve about 8.5 at.% Hf at 1200°C [1981Nas, 1985Mis] and 7 at.% Hf at 1000°C

[1983Och], 8 at.% at 1000°C [1981Nas]. On the other hand, [2002Lou] reported a maximum solubility of

11 at.% Hf in Ni3Al in a rapidly solidified Ni74Al15Hf11 alloy with lattice parameter of a = 364.0 pm, even

though the alloy contained another metastable cubic phase. Substitution of Al by Hf causes a linear increase

in the lattice parameter with increasing Hf content [1984Och1, 1984Och2, 1985Mis]. The rate of increase

in the lattice parameter of Ni3Al, da/dc, is reported to be 0.73 pm/at.% Hf [1985Mis].

81


MSIT®

Al–Hf–Ni

The maximum solid solubility of Hf in NiAl is reported to be about 5 at.% at 1350°C [1990Tak]. Lattice

parameters of Ni3Al and NiAl as a function of alloy composition and heat treatment were reported by

[1981Nas].

[1981Nas] reported that Hf2Ni7 can dissolve up to about 14 at.% Al with a width of 1.5 at.% Hf at 1200 and

1000°C. On the other hand, [1969Mar1] found that none of the Hf-Ni binary compounds, including Hf2Ni7,

can dissolve more than 1 at.% Al at 800°C. Lattice parameters of HfNi3, HfNi5 and Hf2Ni7 phases as a

function of alloy composition and heat treatment were also determined by [1981Nas].

At least ten ternary phases have been reported in this system, of which nine were first reported by Markiv

and co-workers [1964Mar, 1966Mar, 1969Mar1, 1969Mar2]. The Hf6Ni8Al15 phase was first reported by

[1966Gan1, 1966Gan2] and subsequently confirmed by [1969Mar1]. The ternary phase 2 (Hf10Ni19Al)

was reported to be stable above 1000°C, but was not observed by [1981Nas] in the 1200 and 1000°C

isothermal sections. This phase was suggested to be an extension of HfNi2 into the ternary region [1972Pet],

but it has been disproved [1979Bse]. Also, there is experimental evidence [1979Bse, 1981Nas] suggesting

that such a structure is not an equilibrium phase, but most probably stabilized by silica.

Incidentally, in Markiv's [1969Mar1] experiment the specimens were in direct contact with silica tubes,

whereas [1981Nas] kept their specimens in alumina crucibles during annealing treatments. The structures

of the Hf5Ni4Al phase [1969Mar2] and the Hf4Ni16Al5 ( 3 or L phase) were not determined [1969Mar1].

The latter phase was reported to be present in the isothermal section at 800°C [1969Mar1], but was not

observed in the isothermal sections at 1200 and 1000°C [1981Nas].

Accordingly, it has been suggested that the 3 phase forms by a solid state reaction between 1000 and 800°C

[1981Nas]. The ternary phase HfNi2Al has been predicted to form by an invariant transition type reaction

[1981Nas]. According to [1968Dwi], the structure of the HfNiAl phase can be better described by

introducing a slight variation in stacking sequence and by doubling the c-parameter.

The details of the crystal structures and lattice parameters of all the solid phases are listed in Table 1.


The section NiAl-HfNi2Al is established as quasibinary using X-ray analysis, metallography and by

determining the melting temperatures, but only part of this section up to 30 at.% Hf has been reported

[1990Tak]. As shown in Fig. 2, a pseudobinary eutectic reaction L NiAl+HfNi2Al takes place at 1350°C

and 15 at.% Hf. Further experiments are necessary to confirm this phase diagram.

[1981Bal] observed a eutectic microstructure of NiAl and (Ni,Al)7Hf2 phases embedded in Ni3Al matrix in

an as-cast alloy of Ni-20Al-7.5Hf (at.%). This result suggests the possibility of the existence of a

pseudobinary eutectic between NiAl and Ni7Hf2. To corroborate this interpretation further experiments are

needed.


Two transition invariant reactions have been reported [1981Nas] to take place during solidification of the

Ni rich alloys: L+Ni3Al (Ni)+Hf2Ni7 (U1) and L+Hf2Ni7 (Ni)+HfNi5 (U2). However, the temperatures

of occurrence of the above invariant reactions were not reported, but estimated to be between 1275 and

1200°C. Based on the observation of equilibria in the region (Ni)-Ni3Al-Hf2Ni7 between 1200 and 1000°C,

[1981Nas] predicted the presence of an invariant U type reaction (Ni)+Hf2Ni7 Ni3Al+HfNi5 at some

temperature between 1200 and 1000°C. In the same temperature range, [1981Nas] also predicted the

possibility of another invariant U type reaction Hf2Ni7+Hf3Ni7 HfNi3(r)+HfNi2Al. [1991Lee] reported a

speculative invariant reaction scheme for the Ni rich portion of the ternary system. In addition to above

mentioned four U type invariant reactions, their speculative reaction scheme includes eight more U type

invariant reactions involving the liquid phase.

Liquidus Surface

Addition of Al in Ni rich Hf-Ni alloys or addition of Hf in Ni rich Al-Ni alloys decrease the liquidus

temperature [1999Miu]. Figure 3 shows the probable liquidus surface of the Ni corner [1981Nas]. It is based

82


MSIT®

Al–Hf–Ni

on the observation of as-cast microstructures of Ni-(2.5 to 35)Al-(5 to 25)Hf (at.%) alloys. This is in

substantial disagreement with the calculated liquidus surface by Kaufman et al. [1974Kau, 1975Kau]. Also,

in their calculation [1974Kau, 1975Kau] assumed that the 3 phase (Hf4Ni16Al5) melts congruently which

is not supported by the results of [1981Nas]. Additionally, the calculated liquidus temperatures [1974Kau,

1975Kau] were substantially lower than the measured solidus temperatures [1981Nas]. In other words, the

thermodynamic parameters derived by [1974Kau, 1975Kau] certainly overestimate the stability of the

liquid phase. Nonetheless, combining the calculated liquidus of [1975Kau] and limited experimental data

of [1981Nas], Lee and Nash [1991Lee] proposed a tentative liquidus surface up to 40 at.% Al and 50

at.% Hf.

Isothermal Sections

Figures 4 and 5 show the partial isothermal sections of the Ni corner at 1200°C [1981Nas, 1985Nas] and

1000°C [1981Nas], respectively. It should be mentioned that the homogeneity ranges of binary Ni3Al at

1200 and 1000°C as reported by [1981Nas, 1985Nas] were considerably higher than those given by the

presently accepted binary phase diagram [Mas, 1987Hil, 1988Bre]. Figure 6 shows the isothermal section

at 800°C [1969Mar1]. The three-phase fields (Ni)+Ni3Al+HfNi5 and Ni3Al+HfNi5+Hf2Ni7, as reported in

the 800°C isothermal section [1969Mar1], were also found to be present in the 1000°C isothermal section

[1981Nas]. These three-phase fields result from an invariant transition type reaction

(Ni)+Hf2Ni7 Ni3Al+HfNi5 [1981Nas]. However, the calculated isothermal section at 800°C [1974Kau,

1975Kau] showed the presence of a (Ni)+HfNi5+Hf2Ni7 three-phase field, and thus does not take into

account the above transition type reaction [1981Nas]. In Figs. 4 to 6, minor adjustments have been made in

order to comply with the accepted binary phase diagrams. Since Hf2Al phase is not considered to be an

equilibrium phase, previously reported three-phase fields (Hf)+Hf2Al+ 5 and Hf2Al+Hf3Al2+ 5 in the

isothermal section at 800°C [1969Mar1] have been replaced by (Hf)+Hf3Al2+ 5. Computer calculated

isothermal sections, in the range of Ni-50 at.% (Hf+Al), at 1423 and 1323°C [1974Kau, 1975Kau], at 1223,

1123, 1023°C [1974Kau, 1975Kau, 1976Kau] and at 800°C [1974Kau, 1975Kau] have also been reported.

Thermodynamics

Experimental thermodynamic data of ternary alloys is very limited. [1992Alb] determined the activity of Hf

and Al in (Ni3Al)1-xHfx and Ni0.75Al0.25-xHfx alloys in the temperature range of 1088 and 1407°C. Their

data indicate the substitution of Hf for Al in Ni3Al. In fact, thermal conductivity measurement of Ni3(Al,Hf)

also corroborate this behavior [2001Ter].

[1999Dar] measured the low-temperature (3.2 to 10.3 K) specific heat of HfNi2Al ( 2) using an adiabatic

calorimeter, and analyzed the specific heat data in terms of electronic, Debye lattice and Einstein models.

The analysis of experimental data yields the Debye temperature D=15°C. They also calculated the

electronic structure by tight-binding linearized muffin-tin orbital (TB-LMTO) method. Their results

underscore the importance of electron-phonon coupling on the phase stability.

Kaufman and Nesor [1974Kau, 1975Kau, 1976Kau] have performed CALPHAD modeling of the ternary

system, and calculated several isothermal sections.


The constitutional equilibria of this ternary system is very important for developing creep resistant high-

temperature alloys. [1991Miu] studied the creep behavior of Ni-23.5Al-2Hf (at.%) alloy single crystals

oriented close to [001] direction at 850, 900 and 950°C under compressive loads. They observed power-law

creep behavior with an exponent of 3.89 and an activation energy of 360 kJ mol.

Hf is a good solid solution strengthener of Ni3Al. Theoretical calculations show that the strengthening effect

is related to both the site occupancy and local segregation of Hf at antiphase boundaries [1991Wu].

The microstructure and mechanical properties of melt-spun and bulk Hf1Co9Ni61Al29 specimens were

compared to Al-Co-Ni samples [1990Pan].

83


MSIT®

Al–Hf–Ni

Miscellaneous

The solidus temperatures [1981Nas] of some ternary alloys are listed in Table 2.

[1997Nag] found that in the presence of boron, the solubility of Al in Ni3Al is increased while the solubility

of Ni in NiAl is decreased by about 1.25 at.% at 1130°C. These results suggest that it is easier for Al and

Hf to occupy the Ni sites, and it was rationalized in terms of occupancy of interstitial sites by boron atoms.

[1991Sas] studied the microstructure of arc-melted (NiAl)99.5Hf0.5 alloy, and did not find any evidence of

grain refining effect. Even though they observed the presence of precipitates, the absence of grain refining

effect was attributed to the solid-state precipitation.

[2002Lou] carried out rapid solidification of Ni78Al12.5Hf9.5, Ni74Al15Hf11 and Hf20Ni66Al14 alloys. In the

former two alloys they observed a hitherto unknown body-centered cubic phase with lattice parameter

a = 220 pm, while the latter alloy has an amorphous structure. It is uncertain if this cubic phase is indeed 4

which also has similar lattice parameter but with face-centered symmetry. This point was not discussed by

[2002Lou]. Calorimetric study shows that the crystallization temperature of the amorphous alloy is about

577°C at a heating rate of 1.33 °C/s. Other aspects of crystallization behavior of the rapidly solidified alloys

have been discussed by [2002Lou].

References

[1961Now] Nowotny, H., Schob, O., Benesovsky, F. “The Crystal Structure of Zr2Al and Hf2Al” (in


[1962Poe1] Poetschke, M., Schubert, K., “On the Constitution of Some Systems Homologous and

Quasihomologous to T4 - B3. Part I” (in German), Z. Metallkd., 53, 474-488 (1962)

(Experimental, *, 18)

[1962Poe2] Poetschke, M., Schubert, K., On the Constitution of Some Systems Homologous and

Quasihomologous to T4 - B3. Part II” (in German), Z. Metallkd., 53, 548-561 (1962) (Crys.

Structure, Experimental, Equi. Diagram, *, 45)

[1964Mar] Markiv, V.Ya., Voroshilov, Yu.V., Kripyakevich, P.I., Cherkashin, E.E., “New Compounds

of MnCu2Al and MgZn2 Types Containing Al and Ga”, Sov. Phys.-Crystallogr., 9, 619-620

(1965), transl. from Kristallografiya, 9, 737-738 (1964) (Crys. Structure, Experimental, 4)

[1964Rie] Rieger, W., Nowotny, H., Benesovsky, F. “Investigations in Systems Transition Metal (T)-

Boron-Aluminium” (in German), Monatsh. Chem., 95, 1417-1423 (1964) (Crys. Structure,

Experimental, 11)

[1966Gan1] Canglberger, E., Nowotny, H., Benesovsky, F., “On Some New G-Phases” (in German),

Monatsh. Chem., 97, 219-220 (1966) (Crys. Structure, Experimental, 3)

[1966Gan2] Ganglberger, E., Nowotny, H., Benesovsky, F., “New G-Phases” (in German), Monatsh.

Chem., 97, 829-832 (1966) (Crys. Structure, Experimental, 4)

[1966Mar] Markiv, V.Ya., Kripyakevich, P.I., “Compounds of the Type R(X', X'')2 in Systems with

R = Ti, Zr, Hf; X' = Fe, Co, Ni, Cu; and X'' = Al or Ga and Their Crystal Structure”, Sov.

Phys.-Crystallogr., 11, 733-738 (1967), translated from Kristallografiya, 11, 859-865


[1967Kri] Kripyakevich, P.I., Markiv, V.Ya., Mel´nik, Ya.V., “Crystal Structure of Zr-Ni-Al, Zr-Cu-

Ga and Analogous Compounds” (in Ukrainian), Dopov. Akad. Nauk Ukr. RSR, (A8), 750-

753 (1967) (Crys. Structure, Experimental, 9)

[1968Dwi] Dwight, A.E., Mueller, M.H., Conner, R.A., Downey, J.W., Knott, H., “Ternary

Compounds with the Fe2P-Type Structure”, Trans. Metall. Soc. AIME, 242, 2075-2080


[1969Mar1] Markiv, V.Ya., Burnashova, V.V., “The Hf-Ni-Al System”, Russ. Metall. (Engl. Transl.),

(6), 113-115 (1969), translated from Izv. Akad. Nauk SSSR, Met., (6), 181-182 (1969) (Equi.

Diagram, Experimental, #, *, 17)

[1969Mar2] Markiv, V.Ya., Burnashova, V.V., “New Ternary Compounds in the (Sc, Ti, Zr, Hf)-(V, Cr,

Mn, Fe, Co, Ni, Cu)-(Al, Ga) Systems” (in Ukrainian), Dopov. Akad. Nauk Ukr. RSR, (A5),


84


MSIT®

Al–Hf–Ni

[1969Tes] Teslyuk, M.Yu., Intermetallic Compounds with Structure of Laves Phases (in Russian),

Moscow, Nauka, 1969, 1-138 (1969) (Crys. Structure, Equi. Diagram, Review)

[1972Pet] Pet´kov, V.V., Markiv, V.Ya. Gorsky, V.V., “Compound with the MgCu2-Type of

Structure in Alloys of Ni, Zr and Hf” (in Russian), Izv. Akad. Nauk SSSR, Met., (2), 188-192


[1974Fer] Ferro, R., Marazza, R., Rambaldi, G., “Equi-Atomic Ternary Phases in the Alloys of the

Rare Earths with In and Ni and Pd”, Z. Metallkd., 65, 37-39 (1974) (Crys. Structure,

Experimental, 2)

[1974Kau] Kaufman, L., Nesor, H., “Computer Calculated Phase Diagrams for the Ni-W-Al, Ni-Al-Hf,

Ni-Cr-Hf and Co(Cr,Ni)-Ta-C Systems”, NASA Contract No NAS3-17304, National

Aeronautics and Space Administration, Washington, D.C. 20546, 1-58 (1974) (Equi.

Diagram, Thermodyn., Theory, 28)

[1975Kau] Kaufman, L., Nesor, H., “Calculation of the Ni-W-Al, Ni-Al-Hf, Ni-Cr-Hf Systems”, Can.

Metall. Quart., 14, 221-232 (1975) (Equi. Diagram, Thermodyn., Theory, 22)

[1976Kau] Kaufman, L., Nesor, H., “Application of Computer Techniques of Prediction of Metastable

Transitions in Metallic Systems”, Mater. Sci. Eng., 23, 119-123 (1976) (Equi. Diagram,

Theory, 13)

[1977Abr] Abrikosov, N.Kh., “Phase Diagrams of Al and Mg Alloy Systems” in “Phase Diagrams of

Al and Mg Alloy Systems”, Nauka, Moscow, 22-25 (1977) (Crys. Structure, Review, 5)

[1979Bse] Bsenko, L., “The Hf-Ni and Zr-Ni Systems in the Region 65-80 at.% Ni”, J. Less-Common

Met., 63, 171-179 (1979) (Equi. Diagram, Experimental, 13)

[1981Bal] Baldan, A. and West, D.R.F., “Structural Features of Certain Ni-Al-Ta and Ni-Al-Hf Alloys

Containing the ´ and -Phases”, J. Mater. Sci., 16, 24-34 (1981) (Crys. Structure,

Experimental, 28)

[1981Fer] Ferro, R., Marazza, R., “Crystal Structure and Density Data” in “Hafnium: Physicochemical

Properties of its Compounds and Alloys”, Atomic Energy Review, Special Issue No.8., K.L.

Komarek, Ed., IAEA, Vienna, (8), 121-250 (1981) (Crys. Structure, Review, 645)

[1981Nas] Nash, P., West, D.R.F., “Phase Equilibria in Ni-Rich Region of the Ni-Al-Hf System”, Met.

Sci., 15, 347-352 (1981) (Equi. Diagram, Experimental, #, *, 20)

[1983Nas] Nash, P., Nash, A., “The Hf-Ni (Hafnium-Nickel) System”, Bull. Alloy Phase Diagrams, 4,

250-253 (1983) (Equi. Diagram, Review, #, *, 23)

[1983Och] Ochiai, S., Oya, Y., Suzuki, T., “Solubility Data of Ni3Al with Ternary Additions”, Bull. P.

M. E., (52), 1-17 (1983) (Equi. Diagram, Experimental, Review, 39)

[1984Och1] Ochiai, S., Oya, Y., Suzuki, T., “Alloying Behaviour of Ni3Al, Ni3Ga, Ni3Si and Ni3Ge”,

Acta Metall., 32, 289-298 (1984) (Equi. Diagram, Experimental, 90)

[1984Och2] Ochiai, S., Mishima, Y., Suzuki, T., “Lattice Parameter Data of Ni( ), Ni3Al( ') and

Ni3Ga( ') Solid Solutions”, Bull. P. M. E., (53), 15-28 (1984) (Crys. Structure,

Experimental, 66)

[1985Mis] Mishima, Y., Ochiai, S., Suzuki, T., “Lattice Parameters of Ni( ), Ni3Al( ') and Ni3Ga( ')

Solid Solutions with Additions of Transition and B-Subgroup Elements”, Acta Metall., 33,


[1985Nas] Nash, P., “Ni-Base Intermetallics for High-Temperature Alloy Design” in “High-

Temperature Ordered Intermetallic Alloys”, Koch, C.C., Liu, C.T., Stoloff, N.S., (Eds.),

Mat. Res. Soc., Pittsburgh, PA, 423-427 (1985) (Equi. Diagram, Review, #, *, 15)

[1987Hil] Hilpert, K., Kobertz, D., Venugopal, V., Miller, M., Gerads, H., Bremer, F.J., Nickel, H.,

“Phase Diagram Studies on the Al-Ni System”, Z. Naturforsch., 42a, 1327-1332 (1987)

(Equi. Diagram, Experimental, *, 17)

[1988Bre] Bremar, F.J., Beyss, M., Karthaus, E., Hellwig, A., Schober, T., Welter, J.-M., Wenzl, H.,

“Experimental Analysis of the Ni-Al Phase Diagram”, J. Cryst. Growth, 87, 185-192 (1988)


[1990Kum] Kumar, K.S., “Ternary Intermetallics in Aluminium-Refractory Metal-X Systems (X=V,

Cr,Mn,Fe,Co,Ni,Cu,Zn)”, Int. Mat. Rev., 35, 293-327 (1990) (Equi. Diagram, Review, 158)

85


MSIT®

Al–Hf–Ni

[1990Pan] Pank, D.R., Nathal, M.V., Koss, D.A., “Microstructure and Mechanical Properties of

Multiphase NiAl-Based Alloys”, J. Mater. Res., 5, 942-949 (1990) (Experimental, Mechan.

Prop., 18)

[1990Tak] Takeyama, M., Liu, C.T., “Microstructure and Mechanical Properties of NiAl-Ni2AlHf

Alloys”, J. Mater. Res., 5, 1189-1196 (1990) (Equi. Diagram, Experimental, #, *, 22)

[1991Lee] Lee, K.J., Nash, P., “The Al-Hf-Ni System”, J. Phase Equilib., 12, 94-104 (1991) (Equi.

Diagram, Crys. Structure, Review, #, 16)

[1991Sas] Sasaki, K., Morinaga, M., Yukawa, N., “Alloying Effect on thr Solidified Structure of

NiAl”, Proc. Conf. Intermetal. Comp. - Struct. Mechan. Prop., 877-881 (1991) (Abstract,

Equi. Diagram, Experimental, Mechan. Prop., 10)

[1991Mis] Mishima, Y., Hong, Y.M., Suzuki, T., “Determination of the Solvus Surface in Ni-Al-X

Ternary Systems”, Mater. Sci. Eng. A, 146, 123-130 (1991) (Assessment, Equi. Diagram,

Experimental, 5)

[1991Miu] Miura, S., Hayashi, T., Takekawa, M., Mishima, Y., Suzuki, T., “The Compression Creep

Behavior of Ni3Al-X Single Crystals”, Mater. Res. Soc. Symp. Proc.: High-Temp. Ordered

Intermetallic Alloys IV, 213, 623-628 (1991) (Experimental, Phys. Prop., 9)

[1991Wu] Wu, Y.P., Sanchez, J.M., Tien, J.K., “Effect of APB Microsegregation on the Strength of

Ni3Al with Ternary Additions”, Mater. Res. Soc. Symp. Proc.: High-Temp. Ordered

Intermetallic Alloys IV, 213, 87-94 (1991) (Calculation, 22)

[1992Alb] Albers, M., Baba, M.S., Kath, D., Miller, M., Hilper, K., “Chemical Activities in the Solid

Solution of Hf in Ni3Al”, Ber. Bunsen-Ges. Phys. Chem., 96(11), 1663-1668 (1992) (Equi.

Diagram, Experimental, Thermodyn., 25)

[1993Gho] Ghosh, G., “Aluminium-Hafnium-Nickel”, in MSIT Ternary Evaluation Program, in MSIT



Assessment, 30)

[1997Nag] Nagarajan, R.R., Jena, A.K., Ray, R.K., “Phase Equilibria in the ´-Rich Region of the Ni-

Al-Hf System”, Z. Metallkd., 88(1), 87-90 (1997) (Equi. Diagram, Experimental, 16)

[1998Mur] Murray J.L., McAlister A.J., Kahan D.J., “The Al-Hf (Aluminium-Hafnium) System”,

J. Phase Equilib., 19, 376-379 (1998) (Assessment, Crys. Structure, Equi. Diagram, *,14)

[1999Dar] Da Rocha, F.S., Fraga, G.L.F., Brandao, D.E., Da Silva, C.M., Gomes, A.A., “Specific Heat

and Electronic Structure of Heusler Compounds Ni2TAl (T = Ti, Zr, Hf, V, Nb, Ta)”,

Physica B (Amsterdam), 269, 154-162 (1999) (Crys. Structure, Experimental)

[1999Miu] Miura, S., Hong, Y.-M., Suzuki, T., Mishima, Y., “Liquidus and Solidus Temperatures of

Ni-Solid Solution in Ni-Al-X (X: Ti, Zr, and Hf) Ternary Systems”, J. Phase Equilib.,

20(3), 193-198 (1999) (Equi. Diagram, Experimental, 11)

[2001Miu] Miura, S., Unno, H., Yamazaki, T., Takizawa, S., Mohri. T., “Reinvestigation of Ni-Solid

Solution/Liquid Equilibria in Ni-Al Binary and Ni-Al-Zr Ternary Systems”, J. Phase

Equilib., 22, 457-462 (2001) (Equi. Diagram, Experimental, #, *, 9)

[2001Ter] Terada, Y., Ohkubo, K., Mohri, T., Suzuki, T., “Site Preference Determination in

Intermetallic Compounds by Thermal Conductivity Measurement”, J. Mater. Res., 16(8),

2314-2320 (2001) (Calculation, Crys. Structure, Experimental, Thermal Conduct., 63)

[2002Lou] Louzguine, D.V., Inoue, A., “Structure and Transformation Behaviour of Rapidly Solidified

Ni-Al-Hf Alloys”, J. Alloys Compd., 340, 151-156 (2002) (Crys. Structure, Equi. Diagram,

Experimental, 9)

[2003Sal] Saltykov, P., Cornish, L., Cacciamani, G., “Al-Ni (Aluminum-Nickel)”, MSIT Binary


International Services, GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi.


[2003Sch] Schuster, J.C, “Al-Hf (Aluminium-Hafnium)”, MSIT Binary Evaluation Program, in MSIT

Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH,


86


MSIT®

Al–Hf–Ni


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Al)

660.452

cF4

Fm3m

Cu

a = 404.88

a = 404.96

[V-C], pure Al at 24°C

[Mas2], Al at 25°C

( Hf)

2231 - 1743

cI2

Im3m

W

a = 361.5

a = 361.0

[V-C], [Mass2]

[2003Sch] dissolves up to 34 at.% Al at

1450°C

( Hf)

1743

hP2

P63/mmc

Mg

a = 319.8

c = 506.1

[V-C], pure Hf at 25°C [Mas2]

[2003Sch] dissolves up to 30 at.% Al at

1450°C

(Ni)

1455

cF4

Fm3m

Cu

a = 352.32

a = 353.55

a = 353.88

a = 352.4

[V-C], pure Ni at 20°C

at 0.95 at.% Hf [1985Mis]

at 8.0 at.% Al [1985Mis]

[Mas2] dissolves 21.3 at.%Al at 1372°C

[2003Sal]

Hf2Al

< 1160

tI12

I4/mcm

CuAl2

a = 677.6 to 677.9

c = 537.2 to 543.3

[1981Fer], in Hf rich two-phase

alloys

Si stabilized [20003Sch]

Hf3Al2 1590 25

tP20

P42/mnm

Zr3Al2

a = 753.5 to 754.9

c = 690.6 to 691.1

[1981Fer]

Hf4Al3 1200

hP7

P6/mmm

Zr4Al3

a = 513.43 to

533.10

c = 542.2 to 541.4

[1981Fer]

[20003Sch]

HfAl

1800

oC8

Cmcm

CrB

a = 325.3

b = 1083.1

c = 428.2

[1981Fer]

[20003Sch]

Hf2Al3 1640 25

oF40

Fdd2

Zr2Al3

a = 952.1

b = 1376.3

c = 552.2

[1981Fer]

HfAl2 1650 25

hP12

P63/mmc

MgZn2

a = 523.0 to 529.0

c = 865.0 to 874.0

[1981Fer]

HfAl3(h)

1590 - 700

tI8

I4/mmm

TiAl3

a = 389.0 to 393.0

c = 893.0 to 889.0

[1981Fer]


HfAl3(r)

700

tI16

I4/mmm

ZrAl3

a = 398.0 to 401

c = 1714.0 to

1713.0

[1981Fer]

NiAl3 854

oP16

Pnma

NiAl3

a = 661.3

b = 736.7

c = 481.1

[2003Sal]

for 37 at.% Al

Ni2Al3 1138

hP5

P3m1

Ni2Al3

a = 402.8

c = 489.1

[2003Sal]

59.5 to 63.2 at.% Al

87


MSIT®

Al–Hf–Ni

NiAl

1651

cP2

Pm3m

CsCl

a = 286.00 to

288.72

[2003Sal] solid solubility ranges


dissolves < 5 at.% Hf at 1350°C

[1990Tak]

Ni5Al3 700

oC16

Cmmm

Pt5Ga3

a = 753.0

b = 661.0

c = 376.0

[2003Sal], for 37 at.% Al

solid solubility ranges from 32.0 to 37.0

at.% Al

Ni3Al

1372

cP4

Pm3m

AuCu3

a = 356.77 to

358.90

[2003Sal], solid solubility ranges


dissolves 8 at.% Hf at 1000°C

[1981Nas], 11 at.% Hf by rapid

solidification

Ni3Al4< 702

cI112

Ia3d

Ni3Ga4

a = 1140.8 0.1 [2003Sal]

Hf2Ni

1200

tI12

I4/mcm

CuAl2

a = 674.3

c = 558.0

[1981Fer]

[1983Nas]

HfNi

1530

oC8

Cmcm

CrB

a = 322.0

b = 982.0

c = 412.0

[1981Fer]

[1983Nas]

Hf9Ni11

< 1340

tI40

I4/mcm

(or I4/m)

Zr9Pt11

a = 979.0

c = 653.0

[1981Fer]

HfNi2 1200

cF24

Fd3m

Cu2Mg

a = 690.6 [1981Fer]


Hf3Ni71016 - 1250

aP20

P1

Hf3Ni7

a = 651.38

b = 658.9

c = 762.71

= 104.87°

= 104.60°

= 112.71°

[1981Fer]

[1983Nas]

Hf7Ni10

1290

oC68

Aba2

Zr7Ni10

C2ca

a = 912.6

b = 907.8

c = 1227.5

a = 1227.5

b = 907.8

c = 912.6

[1981Fer]

[1983Nas]

Hf8Ni21

1300 - 1175

aP29

P1

Hf8Ni21

a = 642.75

b = 800.07

c = 855.4

= 75.18°

= 68.14°

= 75.61°

[1981Fer]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

88


MSIT®

Al–Hf–Ni

HfNi3(h)

1350 - 1200

hR12

R3m

BaPb3

a = 527.87

c = 1923.24

a = 525.5

c = 1926.0

[1981Fer]

[1981Nas]

HfNi3(r)

1200

hP40

P63/mmc

TaRh3

a = 527.10 to 528.6

c = 2130.0

to 2139.16

[1981Fer, 1981Nas, 1983Nas]

Hf2Ni7 1480

mC36

C2/m

Zr2Ni7

a = 462.0 to 468.0

b = 819.1 to 831.7

c = 1210.2 to

1224.0

= 94.7 to 95.905°

[1981Fer, 1981Nas, 1983Nas]

dissolves up to 11 at.% Al at 1000°C and

14 at.% at 1400°C

HfNi5 1240

cF24

F43m

AuBe5

a = 668.3 to 669.7 [1981Fer, 1981Nas]

* 1, HfNiAl hP9

P62m

Fe2P

a = 686.0

c = 342.0

a = 684.7

c = 345.9

a = 688.5

c = 683.8

a = 687.3

c = 343.7

[1966Mar], annealed at 900°C for 480 h

[1967Kri]

[1968Dwi], annealed between 700

and 900°C

[1974Fer], annealed at 600°C (> 168 h)

* 2, HfNi2Al

1450

cF16

Fm3m

MnCu2Al

a = 608.1

a = 601.8

a = 607.3

a = 606.5

a = 608.2

a = 601.1

a = 607.4

a = 608.1

[1964Mar], at 50 at.% Ni,

25 at.% Al and 25 at.% Hf,

annealed at 800°C for 480 h

[1981Nas], in an alloy of 60 at.% Ni,



[1981Nas], in the same alloy as

above but annealed at 1000°C for 168 h







[1981Nas], in an alloy of

61.3 at.% Ni, 20.3 at.% Al and

18.4 at.% Hf, annealed at 1200°C

for 168 h

[1981Nas], in the same alloy as

above but annealed at 1000°C for 168 h

[1999Dar]

* 3, Hf3Ni6Al16 tI16

I4/mmm

ZrNi2Al5

a = 401.0

c = 1412.0

[1969Mar1, 1969Mar2]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

89


MSIT®

Al–Hf–Ni

Table 2: Solidus Temperatures as a Function of Alloy Composition [1981Nas]

* 4, Hf6Ni8Al15 cF16

Fm3m

Th6Mn23

a = 1200.0 [1966Gan1, 1966Gan2, 1969Mar1]

* 5, Hf6NiAl2 hP9

P62m

Hf6CoAl2

a = 783.0

c = 329.0

[1969Mar1, 1969Mar2]

* 6, Hf5Ni4Al ? - [1969Mar2]

* 1, Hf5Ni3Al7 hP12

P63/mmc

MgZn2

a = 518.0

c = 841.0

[1969Mar1]

* 2, Hf3NiAl5 cF24

Fd3m

Cu2Mg

a = 734.7 [1966Mar, 1969Tes]

* 2, Hf10Ni19Al

1000

cF24

Fd3m

Cu2Mg

a = 690.5 [1969Mar1], possibly stabilized by silica

* 3, Hf4Ni16Al5 - - Denoted as L phase by [1969Mar1]

Alloy Composition (at.%) SolidusTemperature [°C 12°C]

Al Hf Ni

9

25

13

15

5

20

5

20.3

2.5

8

15

13

23

20

16

5

18.4

22.5

83

60

74

62

75

64

70

61.3

75

1237

1233

1262

1227

1262

1237

1233

1233

1233

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

90


MSIT®

Al–Hf–Ni

10

90

10

Hf 20.00Ni 80.00Al 0.00

Ni

Hf 0.00Ni 80.00Al 20.00 Data / Grid: at.%

Axes: at.%

1127°C

1027°C

927°C

827°C

(Ni)

Fig. 1: Al-Hf-Ni.

Solubility isotherms

of (Ni)

10 20

1000

1250

1500

1750

Hf 0.00Ni 50.00Al 50.00

Hf 30.00Ni 50.00Al 20.00Hf, at.%

Te

mp

era

ture

, °C

NiAl

NiAl+τ2

τ2

L+NiAl

L

1350°C

15% Hf

1651°C

L+τ2

Fig. 2: Al-Hf-Ni.

Pseudobinary system

NiAl-HfNi2Al

91


MSIT®

Al–Hf–Ni

10

20

80 90

10

20

Hf 30.00Ni 70.00Al 0.00

Ni


Axes: at.%

(Ni)

Ni3Al

Hf2Ni

7

HfNi5

U1

U2

p1

p2

e1

10

20

30

40

60 70 80 90

10

20

30

40

Hf 50.00Ni 50.00Al 0.00

Ni


Axes: at.%

(Ni)

(Ni)+Ni3Al

L+(Ni)

(Ni)+HfNi5

LHfNi5Hf

2Ni

7

HfNi3(r)

Hf3Ni

7

τ2

NiAl

NiAl+Ni

3Al

Ni3Al

τ 2+

NiA

l

τ2 +Hf

2 Ni7

Hf8Ni

21

Hf 2

Ni 7

+N

i 3A

l

Hf 2

Ni 7

+τ 2

+N

iAl

Fig. 3: Al-Hf-Ni.

Liquidus surface of

the Ni corner

Fig. 4: Al-Hf-Ni.

Partial isothermal

section at 1200°C.

The dashed lines

represent interpolated

boundaries

92


MSIT®

Al–Hf–Ni

20

40

60

80

20 40 60 80

20

40

60

80

Hf Ni


Axes: at.%

(Ni)

Ni3Al

NiAl

Ni2Al

3

NiAl3

L

HfAl3(r)

HfAl2

Hf2Al

3

HfAl

Hf4Al

3

Hf3Al

2

(αHf)

Hf2Ni HfNi Hf

9Ni

11 Hf7Ni

10

HfNi3

Hf2Ni

7

HfNi5

λ2

λ1

τ1

τ2

τ6

τ5

τ3

τ4

λ3

10

20

30

40

60 70 80 90

10

20

30

40

Hf 50.00Ni 50.00Al 0.00

Ni


Axes: at.%

(Ni)

(Ni)+HfNi5

HfNi5

βHfNi3

τ2

τ2+βHfNi

3+Hf

2Ni

7

NiAl+τ2

+Hf2Ni

7

NiAl

NiA

l+N

i 3A

l

Ni3Al

(Ni)+Ni3Al

(Ni)+HfNi5

+Ni3Al

NiA

l+N

i 3A

l+H

f 2N

i 7

Hf2Ni

7

Fig. 6: Al-Hf-Ni.


800°C

Fig. 5: Al-Hf-Ni.

Partial isothermal

section at 1000°C.

The dashed lines


boundaries

93


MSIT®

Al–Li–Mg

Aluminium – Lithium – Magnesium

Gautam Ghosh

Literature Data

The ternary system contains many technologically important alloys for light weight, high-strength and

corrosion resistant applications. Therefore, the phase equilibria of the system are of experimental and

theoretical interest. Extensive studies have been carried out on the aging behavior and the structure-property

relationship of Al rich alloys. Comprehensive reviews of the phase equilibria have been published by

[1990Goe, 1993Gho]. [1948Sha2] was the first to report the entire liquidus surface. Later, [1977Dri,

1979Vos, 1981Sch3] reinvestigated the liquidus surface. Isothermal sections have been investigated several

times [1948Sha2, 1954Wei, 1956Lev, 1955Row, 1956Row, 1973Dri1, 1973Dri2, 1976Pad, 1977Dri,

1980Sch4]. Until recent studies on the constitutional equilibria of the entire system by Schürmann and

co-workers [1979Vos, 1979Gei, 1980Sch4, 1981Sch3], earlier results were subjected to considerable doubt

and inaccuracy in view of high reactivity and volatility of Li and Mg. To overcome this problem,

Schürmann et al. [1980Sch1, 1981Sch1] designed a special experimental apparatus to prepare the binary

and ternary alloys, and the phases were analyzed by X-ray diffraction, optical metallography and electron

probe microanalysis. Accordingly, much of their results are reproduced in this assessment with some

amendments.

Binary Systems

The Al-Li binary phase diagram is taken from the assessment of McAlister [1982McA, Mas]. In his

assessment, the experimental results of [1979Vos, 1980Sch2] were not reviewed. Nevertheless, the liquidus

data and the invariant reaction temperatures involving liquid, (Al), LiAl, and Li3Al2 phases of [1979Vos,

1980Sch2] agree very well with those of [1982McA]. Also, all these results are in good agreement with the

recent thermodynamic assessment of Saunders [1989Sau]. According to [1982McA], the peritectic reaction

L+Li3Al2 Li9Al4 occurs at 335°C. On the other hand, [1979Vos, 1981Sch2] proposed the peritectic

reaction to be L+Li3Al2 (Li) at 329°C. But, this was found to be incompatible with the thermodynamic

modeling by Saunders [1989Sau]. Voss [1979Vos] reported a eutectoid reaction Li3Al2+(Li) (~Li21Al4)

at 242°C, and this feature was also absent in the assessments of [1982McA, 1989Sau]. Also, [1979Vos]

reported an unusually high solid solubility (about 13.0 at.%) of Al in (Li). Once again, this feature was found

to be incompatible with the thermodynamic modeling of the Al-Li system [1989Sau].

In an earlier assessment [1993Gho] of the Al-Li-Mg system, the Al-Mg binary phase diagram was accepted

from the experimental work of [1979Vos, 1980Sch3] which was somewhat different from Murray's

assessment [1982Mur, Mas]. The major discrepancy lied in the composition range of 40 to 50 at.% Mg.

Schürmann et al. [1979Vos, 1980Sch3] reported two intermediate phases (Mg10Al11) and (Mg9Al11),

which were absent in the assessed phase diagram of [1982Mur]. Also, [1980Sch3] did not observe the

R-phase which was reported to exist between 320 to 370°C and at 42 at.% Mg [1982Mur].

Thermodynamically assessed [1990Sau] Al-Mg phase diagram was in excellent agreement with the

experimental phase diagram of [1979Vos, 1980Sch3]. Recently, the high-temperature phase equilibria

between (Mg2Al3) and (Mg17Al12) phases has been reinvestigated in detail [1997Su, 1998Don,

1998Lia]. Therefore, the Al-Mg phase diagram is accepted from the experimental and thermodynamic

calculation of [1998Lia], which was also accepted in the recent evaluation by [2003Luk]. In the composition

range of 50 to 60 at.% Al, the phase diagram of [1998Lia] is substantially different from that of [1979Vos,

1980Sch3] and similar to the assessed diagram of [1982Mur]. [1998Lia] found that between (Mg2Al3) and

(Mg17Al12) phases, there is only one intermediate phase (Mg23Al30). Moreover, phase forms by a

peritectoid reaction at 410°C and decomposes by a eutectoid reaction at 210°C. Recent experimental

investigations by [1997Su, 1998Don] have shown that the phase reported by Schürmann et al. [1980Sch3,

1981Sch2] in the temperature range of 410 to 452°C does not exist. To account for the additional peaks

observed in the X-ray diffraction, [1997Su] assumed the presence of a hypothetical phase having

94


MSIT®

Al–Li–Mg

composition between 57 to 58 at.% Al. Donnadieu et al. [1998Don] carried out electron diffraction

experiments of several alloys containing 47 to 59 at.% Al which were annealed between 425 to 445°C. They

observed modulated microstructure of the phase. The wave vector characterizing the commensurate

modulation is temperature and composition dependent. Therefore, the additional peaks observed by

[1997Su] in X-ray diffraction could be explained by the commensurate modulation.

The Li-Mg binary phase diagram is taken from the recent review and thermodynamic assessment of

Nayeb-Hashemi et al. [1984Nay].

Solid Phases

Depending on temperature, (Al) can dissolve up to 16.5 at.% Mg [1979Vos, 1980Sch4] and 15.8 at.% Li

[1982McA, 1980Sch2]. Solid solubility of (Al) in the ternary regime is shown in Fig. 1, as a function of

temperature [1965Fri, 1973Dri2, 1980Sch4]. The results of [1973Dri2, 1980Sch4] agree fairly well, but the

results of [1965Fri] indicate that at a given temperature and Li content the solid solubility of Mg in (Al) was

less than those reported by [1973Dri2, 1980Sch4].

The solid solubility of Al and Li in (Mg) has been reported by several investigators [1948Sha2, 1952Jon,

1976Pad, 1979Gei, 1980Sch4]. The compositions of (Mg), as a function of temperature, in the

(Mg)+(Li)+LiAl( ) and (Mg)+LiAl( )+Mg17Al12( ) three phase fields are listed in Table 1. In general,

there is systematic disagreement between the results of [1980Sch4] and those of the others. Figures 2 and 3

show (Mg)/(Mg)+(Li) phase boundaries in vertical sections at 1.0 and 2.0 mass% Al respectively [1952Jon,

1954Wei, 1955Row, 1979Gei, 1980Sch4]. Figures 4, 5 and 6 show the (Mg)/(Mg)+Mg17Al12( ) phase

boundaries in vertical sections at 1.0, 2.0 and 4.0 mass% Li respectively [1952Jon, 1954Wei, 1955Row,

1976Pad, 1979Vos, 1980Sch4]. Along these sections, there is significant disagreement between the results

of Voss [1979Vos, 1980Sch4] and those due to [1952Jon, 1954Wie, 1955Row, 1976Pad]. In drawing the

phase boundaries in Figs. 4 to 6, weightage is given to the results of Voss [1979Vos, 1980Sch4].

In the ternary regime, Mg17Al12( ) dissolves up to about 20 at.% Li, Mg2Al3( ) dissolves up to about

7 at.% Li, Mg23Al30( ) dissolves about 0.8 at.% Li, and LiAl ( ) dissolves up to about 17 at.% Mg

[1979Vos, 1980Sch4]. The lattice parameter of Mg17Al12( ) decreases with the addition of Li [1956Lev].

Two ternary phases, 1 and 2, have been reported in this system. The 1 phase was first reported by

Shamray [1948Sha1, 1948Sha2], and has been confirmed by subsequent investigators [1954Wei, 1955Lev,

1956Lev, 1973Tho, 1976Pad, 1979Vos, 1980Sch4]. Originally, the stoichiometry of the 1 phase was

designated as LiMgAl2 [1948Sha1, 1948Sha2, 1954Wei, 1955Lev, 1956Lev]. However, recent results of

[1976Pad, 1979Vos] indicate that 1 phase contains 32.0 to 34.2 at.% Li and 13.5 to 14.0 at.% Mg, and this

composition is accepted here in drawing the isothermal sections. The Li and Mg contents of 1 phase

reported by [1948Sha2, 1954Wei, 1955Lev, 1956Lev] differ significantly as compared to those of

[1976Pad, 1979Vos]. The 2 phase, having stoichiometry Li2MgAl and NaTl type of structure, was reported

by earlier investigators [1952Jon, 1955Row], but could not be confirmed in subsequent investigations

[1954Wei, 1968Pau, 1979Vos]. Rather, it has been reported that 2 is a nonequilibrium transitional phase

[1954Wei, 1985Nik]. Accordingly, this phase is not considered in drawing the isothermal sections.

The details of the crystal structures and lattice parameters of the equilibrium solid phases are listed in

Table 2.


Figure 7 shows the reaction scheme associated with the solidification of Al-Li-Mg alloys after [1981Sch3].

However, several modifications are made for consistency with the accepted Al-Mg binary phase diagram.

Three pseudo-binary reactions p1, e3 and e4, all of which give rise to a maximum on the liquidus surface,

have been reported [1981Sch3]. From the vertical sections reported by Voss [1979Vos, 1981Sch3], the

temperatures of the three maxima are estimated to be 545, 485 and 480 10°C respectively. The pseudo

binary reactions p1 [1948Sha1, 1981Sch3] and e3 [1981Sch3] give rise to the formation of the ternary phase

1, but the latter reaction was originally reported to be occurring at 477°C and peritectic type i.e.,

L+ 1=LiAl( ) [1948Sha1]. [1981Sch3] reported that three U type reactions U6, U7 and U8 occur at 458,

451 and 449°C, respectively. In this assessment, the U6 invariant reaction of [1981Sch3] is rewritten as a

95


MSIT®

Al–Li–Mg

ternary peritectic reaction P1. Since the (Mg10Al11) phase of [1981Sch3] does not exist in the accepted

Al-Mg phase diagram and the (Mg23Al30) phase forms by a solid-state reaction, the U7 and U8 reactions

of [1981Sch3] are not accepted. Also, it is doubtful if three invariant reactions, as proposed by Schürmann

et al. [1981Sch3], occurring within a temperature interval of 9°C could be firmly established. A new ternary

peritectic reaction P2, so far undetected, has been introduced at the Li corner and it is expected to be

occurring at 350 20°C. All these amendments are consistent with the experimentally observed isothermal

sections, vertical sections and the accepted binary phase diagrams. The compositions of the phases

participating in the invariant equilibria [1979Vos, 1981Sch3] are listed in Table 3.

Liquidus, Solidus and Solvus Surfaces

Figure 8 shows the liquidus surface and the melting grooves separating eleven areas of primary

crystallization [1979Vos, 1981Sch3]. Approximate isotherms at 25°C interval are also shown in Fig. 8.

There is considerable discrepancy between the liquidus surface reported by Voss [1979Vos, 1981Sch3] and

those due to Shamray [1948Sha2] and Drits et al. [1977Dri]. Also the binary phase diagrams accepted by

Shamray is quite different from the presently accepted ones. Accordingly, the liquidus surfaces reported by

[1948Sha2, 1977Dri] were not considered here. [1986Dub, 1987Dub] employed CALPHAD technique to

calculate the liquidus surface of the Al corner. According to their calculation, the temperatures of invariant

reactions U1 and U2 agree very well with those of experimental ones. But, [1987Dub] predicted a ternary

eutectic reaction L=(Al)+Mg17Al12+Mg2Al3 at 447°C since they assumed no Li-solubility in the Mg2Al3( ) phase.

Figure 9 shows the solidus surface of the entire ternary system, after [1981Sch3]. The diagram is still

incomplete in the Li corner. The solidus temperatures as given in the vertical sections reported by Drits et

al. [1973Dri1, 1977Dri] are in reasonably good agreement with [1981Sch3].

Isothermal Sections

Partial isothermal sections have been reported several times [1948Sha2, 1952Jon, 1954Wei, 1956Lev,

1956Row, 1956Lev, 1973Dri1, 1973Dri2, 1976Pad, 1977Dri, 1979Gei, 1980Sch4]. Among these, the

results of Schürmann et al. [1980Sch4] are considered to be the most accurate. They prepared about 178

ternary alloys in a specially designed vacuum induction furnace [1980Sch1]. The alloys were annealed at

400, 300 and 200°C for 260 h and subsequently quenched in water or oil. The phase analysis was carried

out by metallography and electron probe microanalysis. The isothermal sections at 400, 300 and 200°C are

shown in Figs. 10, 11 and 12, respectively. In this composition range, the essential feature of the phase fields

remain same down to room temperature [1948Sha2, 1954Wei, 1956Row]. Even though several three-phase

fields are shown dashed in Figs. 10 to 12, they are consistent with the reaction scheme shown in Fig. 7.

Minor adjustments have been made in Figs. 8 to 12 along the binary edges. The partial isothermal sections

of [1948Sha2, 1954Wei, 1956Lev, 1956Row, 1973Dri1, 1973Dri2, 1976Pad] agree qualitatively with those

of [1980Sch4]. The discrepancies between the results of [1980Sch4] and those of others are primarily due

to the fact that the solid solubilities of the binary intermediate phases in the ternary regime were not

determined accurately.

[1977Sab, 1978Sab] reported the computer calculated isothermal sections in the temperature range of 375

to 500°C. The calculations were done based on the binary solution-phase interaction parameters and

compound parameters. Also, binary intermediate phases were assumed to be stoichiometric, and no ternary

interaction parameter and ternary phase were taken into account. Accordingly, substantial disagreement

between the calculated and the experimental isothermal sections was noticed. However, the isothermal

sections of the Al corner calculated by Dubost et al. [1987Dub], agree reasonably well with those

experimentally observed.


[1948Sha1] reported vertical sections at 5, 10, 15, 20, 30, 50 and 60 at.% Li and also along Mg17Al12-LiAl

and LiMg2-Al. Among these, the vertical section at 50 at.% Li was reported to be pseudobinary type.

96


MSIT®

Al–Li–Mg

However, most of his results are incompatible with the accepted binary phase diagram. [1954Wei] reported

two vertical sections along Mg-Li5Mg4 and LiMg7-Al. [1973Dri1] determined two isopleths at 30 and 32

mass% Mg. [1977Dri] reported an isopleth at 60 mass% Mg.

Thermodynamics

[1991Mos] determined the enthalpies of mixing of liquid Al-Li-Mg alloys in the temperature range of 596

to 758°C using an isothermal high temperature mixing calorimeter. Their data indicate the presence of

ternary interactions.

Miscellaneous

Because of structural applications, the decomposition behavior of supersaturated Al-Li-Mg alloys have

been studied several times [1950Fro, 1971Fri1, 1971Fri2, 1982Cha, 1983Fri, 1985Nik, 1986Kru, 1987Flo,

1994Kra, 1997Kim, 1998Cho]. The mechanical properties associated with such decomposition process

have also been studied a number of occasions [1950Bus, 1956Row, 1965Fri, 1971Fri1, 1982Cha, 1983Fri,

1984Gil, 1985Nik, 1986Kru, 1994Kra, 1997Hwa]. Decomposition of supersaturated Al-(1.5 to

2.0)Li-(4 to 6)Mg (mass%) alloys take place through the formation of a metastable phase [1973Tho]. The

structure of this metastable phase has been reported [1980Shc] to be face-centered monoclinic having lattice

parameters a = c = 2000.4 pm, b = 1979.7 pm and = 88.83°. [1993Nii, 1994Tsa] reported the formation

of a face-centered icosahedral phase in rapidly solidified Li25Mg25Al50 and Li10Mg40Al50 alloy,

respectively. The electronic origin of such a quasicrystalline phase has been discussed by [1997Del]. It has

been predicted [1994Hos] that Mg will occupy the Al sublattice in the metastable phase LiAl3 having L12

structure.

References

[1948Sha1] Shamray, F.I., Kurnakov, N.S., “The Ternary System Aluminium-Magnesium-Lithium. II.

State Diagrams of Auxilliary Sections” (in Russian), Bull. Acad. Sci. URSS, Classe Sci.

Chim., (1), 83-94 (1948) (Experimental, Equi. Diagram, 0)

[1948Sha2] Shamray, F.I., “Ternary System: Aluminium-Magnesium-Lithium. III. Description of the

Ternary System Aluminium-Magnesium-Lithium. Projection of the Liquidus Surface,

Isotherms at 400°C and 20°C, and the Process of Crystallisation” (in Russian), Izv. Akad.

Nauk SSSR, Otdel Khim. Nauk, (3), 290-301 (1948) (Experimental, Equi. Diagram, *, 0)

[1950Bus] Busk, R.S., Leman, D.L., Casey, J.J., “The Properties of Some Magnesium-Lithium Alloys

Containing Aluminium and Zinc”, Trans. AIME, J. Met., 188, 945-951 (1950)

(Experimental, 6)

[1950Fro] Frost, P.D., Kura, J.G., Eastwood, L.W., “Aging Characteristics of Magnesium-Lithium

Base Alloys”, Trans. AIME, J. Met., 188, 1277-1282 (1950) (Experimental, 3)

[1952Jon] Jones, A., Lennon, J.H., Nash, R.R., W.H. Chang, E.G. Macpeek, “Magnesium Alloy

Research Studies”, U. S. At. Energy Comm. Publ., (AF-TR-52-169), 1-130 (1952)

(Experimental, Equi. Diagram, #, 16)

[1954Wei] Weinberg, A.F., Levison, D.W., McPherson, D.J., Rostoker, W., Wolfe, C.P.,

Humphreys, A., Dvorak, J., Manasevit, H., DuPraw, W., “Phase Relationships in

Magnesium - Lithium - Aluminum and Magnesium - Lithium - Zinc Alloys”, Armour Res.

Found. Rep., (AD-16567), 1-94 (1954) (Experimental, Equi. Diagram, #, *)

[1955Row] Rowland, J.A., Armantrout, Jr.,C.E., Walsh, D.F., “Magnesium-Rich Corner of the

Magnesium-Lithium-Aluminum System”, Trans. AIME, J. Met., 203, 355-359 (1955)

(Experimental, Equi. Diagram, #, *, 11)

[1955Lev] Levison, D.W., “Discussion on Magnesium-Rich Corner of the Magnesium - Lithium -

Aluminum System by Rowland, J.A.,Jr., Armantrout, C.E., Walsh, D.F.”, Trans. AIME,

J. Met., 203, 1267 (1955) (Experimental, 1)

97


MSIT®

Al–Li–Mg

[1956Lev] Levison, D.W., McPherson, D.J., “Phase Relations in Magnesium-Lithium-Aluminum

Alloys”, Trans. Am. Soc. Met., 48, 689-697 (1956) (Experimental, Equi. Diagram, #, *, 9)

[1956Row] Rowland, J.A., Armantrout, C.E., Walsh, D.F., “Experimental Magnesium Alloys

Containing Nickel, Manganese, Lithium and Aluminum”, U. S. Bur. Mines, Rep. Invest.,

5250, 1-21 (1956) (Experimental, Equi. Diagram, #, 11)

[1965Fri] Fridlyander, I.N., Shamray, V.F., Shiryaeva, N.V., “Phase Composition and Mechanical

Properties of Alloys of Aluminum with Magnesium and Lithium” (in Russian), Izv. Akad.

Nauk SSSR, Met., (2), 153-158 (1965) (Experimental, Equi. Diagram, #, 9)

[1968Pau] Pauly, H., Weiss, A., Witte, H., “FCC Alloys of Composition Li2MgX with Body-Centred

Substructure” (in German), Z. Metallkd., 59, 414-418 (1968) (Crys. Structure,


[1971Fri1] Fridlyander, I.N., Sandler, V.S., Nikol'skaya, T.I., “Change in the Phase Composition of

Aluminum-Magnesium-Lithium Alloy 01420 During Aging” (in Russian), Metall. i Term.

Obra. Metallov., (5), 2-5 (1971) (Crys. Structure, Experimental, 7)

[1971Fri2] Fridlyander, I.N., Sandler, V.S., Nikol'skaya, T.I., “Investigation of the Aging of

Aluminum-Magnesium-Lithium Alloys”(in Russian), Fiz. Met. Metalloved., 32, 767-774


[1972Sam] Samson, S., “Structural Relationships Among Complex Intermetallic Compounds”

(Abstract Only), IXth International Congress of Crystallography, Kyoto, Japan, VII-7, 96


[1973Dri1] Drits, M.E., Padezhnova, E.M., Guzei, L.S., “On the Question of the Mg-Li-Al System” in

“Certain Regularities in the Structure of Phase Diagrams of Metallic Systems”, Baikov Inst.

Met., Nauka, Moscow, 147-153 (1973) (Experimental, Equi. Diagram, #, *, 5)

[1973Dri2] Drits, M.E., Kadaner, E.S., Turkina, N.I., Kuz'mina, V.I., “Study of Phase Equilibria in the

Solid State in the Al-Corner of the Al-Mg-Li System” (in Russian), Izv. Akad. Nauk SSSR,

Met., (2), 225-229 (1973) (Experimental, Equi. Diagram, #, *, 5)

[1973Tho] Thompson, G.E., Noble, B., “Precipitation Characteristics of Al-Li Alloys Containing Mg”,

J. Inst. Met., 101, 111-115 (1973) (Crys. Structure, Experimental, 6)

[1976Pad] Padezhnova, E.M., Melmik, E.V., Guzei, L.S., Guseva, L.N., “Phase Equilibria in the

Magnesium-Lithium-Aluminum System at 300°C” (in Russian), Izv. Akad. Nauk SSSR,

Met., (4), 222-226 (1976) (Experimental, Equi. Diagram, #, *, 8)

[1977Dri] Drits, M.E., Padezhnova, E.M., Guzei, L.S., “Magnesium - Lithium - Aluminum Phase

Diagram” (in Russian), Izv. Akad. Nauk SSSR, Met., (2), 205-209 (1977) (Experimental,

Equi. Diagram, #, *, 5)

[1977Sab] Saboungi, M.L., Hsu, C.C., “Computation of Isothermal Sections of the Al-Li-Mg System”,

Calphad, 1, 237-251 (1977) (Equi. Diagram, Theory, Thermodyn., 29)

[1978Sab] Saboungi, M.L., Hsu, C.C., “Estimmation of Isothermal Sections of Ternary Phase

Diagrams of Lithium Containing Systems: The Al-Li-Mg System” in “Applications of

Phase Diagrams in Metallurgy and Ceramics”, Vol. 2, NBS Special Publ. No 496,

Washington, DC, 1109-1138 (1977) (Equi. Diagram, Theory, Thermodyn., 29)

[1979Gei] Geissler, I., “Phase Equilibria of Al-Li-Mg Alloys at 200, 300 and 400°C and their Hardness

in the as Cast State” (in German), Ph. D. Thesis, TU Clausthal (1979) (Experimental, Equi.

Diagram, #, *, 44)

[1979Vos] Voss, H.-J., “Development of an Apparatus for Melting Lithium-Containing

Magnesium-Aluminium Alloys and its use for Thermal Analysis” (in German), Ph. D.

Thesis, TU Clausthal, 82 pp., (1979) (Experimental, Equi. Diagram, #, *, 14)

[1980Sch1] Schuermann, E., Geissler, I.K., “Phase Equilibria in the Solid Condition of the Aluminum

rsp. the Magnesium-rich Corner of the Ternary System of Aluminum-Lithium-Magnesium.

Part I. Testing Methods and Design of a Proper Melting Aggregate for

Aluminum-Lithium-Magnesium Alloys” (in German), Giessereiforschung, 32(2), 163-164

(1980) (Experimental, Equi. Diagram, #, *, 4)

98


MSIT®

Al–Li–Mg


rsp. the Magnesium-rich Corner of the Ternary System of Aluminum -Lithium - Magnesium

Part II. Phase Equilibria in the Solid Condition of the Aluminium rsp Magnesium Rich

Zones of the Binary Systems Aluminium-Lithium and Magnesium-Lithium” (in German),

Giessereiforschung, 32(2), 165-167 (1980) (Experimental, Equi. Diagram, #, *, 17)


rsp. the Magnesium-rich Corner of the Ternary System of Aluminum-Lithium-Magnesium.

Part III. Phase Equilibria in the Solid Condition of the Binary System Aluminium-

Magnesium” (in German), Giessereiforschung, 32(2), 167-170 (1980) (Experimental, Equi.

Diagram, #, *, 15)


rsp. the Magnesium-Rich Corner of the Ternary System of Aluminum-Lithium-Magnesium.

Part IV. Phase Equilibria in the Solid Condition of the Ternary System of Aluminum-

Lithium-Magnesium” (in German), Giessereiforschung, 32(2), 170-174 (1980)


[1980Shc] Shchegoleva, T.V., Rybalko, O.F., “The Structure of the Metastable S'-Phase in an

Al-Mg-Li Alloy” (in Russian), Fiz. Met. Metalloved, 50(1), 86-90 (1980) (Crys. Structure,

Experimental, 7)

[1981Sch1] Schuermann, E., Voss, H.-J., “Investigation of the Melting Equilibria of the Magnesium-

Lithium-Aluminum Alloys. Part I. Description of the Melting Equipment and Realization

of the Research” (in German), Giessereiforschung, 33(1), 33-35 (1981) (Experimental,

Equi. Diagram, *, 5)

[1981Sch2] Schuermann, E., Voss, H.-J., “Investigation of the Melting Equilibria of the

Magnesium-Lithium-Aluminum Alloys. Part IV. Melting Equilibria of the Binary System

Magnesium - Lithium” (in German), Giessereiforschung, 33(2), 43-46 (1981)


[1981Sch3] Schuermann, E., Voss, H.-J., “Investigation of the Melting Equilibria of the

Magnesium-Lithium-Aluminum Alloys. Part V. Melting Equilibria of the Ternary System

of Magnesium - Lithium-Aluminum” (in German), Giessereiforschung, 33(2), 47-53 (1981)


[1982Cha] Chanani, G., Narayanan, G. H., Telesman, I.J., “Heat Treatment, Microsrtucture and

Mechanical Property Correlations in Al-Li-Cu and Al-Li-Mg P/M Alloys”, “High-Strongth

Powder Metallurgy Aluminum Alloys”, Proc. Conf., Dallas, TX, 1982, TMS-AIME,

Warrandale, PA, 341-368 (1968) (Crys. Structure, Experimental, 14)

[1982McA] McAlister, A.J., “The Al-Li (Aluminum-Lithium) System”, Bull. Alloy Phase Diagrams,

3(2), 177-183 (1982) (Assessment, Equi. Diagram, Thermodyn., #, *, 31)

[1982Mur] Murray, J.L., “The Al-Mg (Aluminum-Magnesium) System”, Bull. Alloy Phase Diagrams,

3(1), 60-74 (1982) (Equi. Diagram, Review, Thermodyn., #, *, 112)

[1983Fri] Fridlyader, I.N., Sandler, V.S., Nikol'skaya, T.I., “Characteristics of the Structure and

Properties of 1420 Aluminum Alloy” (in Russian), Metall. Term. Obra. Metallov, (7), 20-22


[1984Gil] Gilman, P.S., “The Physical Metallurgy of Mechanically Alloyed, Dispersion-Strengthened

Al-Li-Mg and Al-Li-Cu Alloys” in “Aluminum-Lithium Alloys II”, Proc. Conf., Monterey,

1984, TMS-AIME, Warrandale, PA, 485-506 (1984) (Crys. Structure, Experimental, 11)

[1984Nay] Nayeb-Hashemi, A.A., Clark, J.B., Pelton, A.D., “The Li-Mg (Lithium-Magnesium)

System”, Bull. Alloy Phase Diagrams, 5(4), 365-374 (1984) (Equi. Diagram, Review,

Thermodyn., #, *, 37)

[1985Nik] Nikulin, L.V., Shevrikuko, S.B., Belozerova, E.V., “Properties and Structure of Cast

Mg-Li-Al -Alloys” (in Russian), Tsvetn. Met., (12), 56-59 (1985) (Experimental, 5)

[1986Dub] Dubost, B., Bompard, P., Ansara I., “Contribution to the Establishment of the Equilibrium

Diagram of Phases of the Al-Li-Mg System” (in French), Mem. Etud. Sci. Rev. Metall., 83,

437 (1986) (Experimental, Equi. Diagram, Theory, #, 6)

99


MSIT®

Al–Li–Mg

[1986Kru] Kruglov, B.F., Khristoferov, C.M., Sheikman, A.I., “Effect of Natural Aging in an Al-2.2

wt.% Li-5.6 wt.% Mg Alloy” (in Russian), Fiz. Met. Metalloved, 61(1), 190-191 (1986)

(Experimental, 11)

[1987Dub] Dubost, B., Bompard, P., Ansara, I., “Experimental Study and Thermodynamic Calculation

of the Al-Li-Mg Equilibrium Phase Diagram”, J. Phys.(France), C3, 473-479 (1987)

(Experimental, Equi. Diagram, Theory, Thermodyn., #, 15)

[1987Flo] Flower, H.M., Gregson, P.J., “Solid State Phase Transformations in Aluminum Alloys

Containing Lithium”, Mater. Sci. Technol., 3, 81-90 (1987) (Crys. Structure, Review, 116)

[1989Sau] Saunders, N., “Calculated Stable and Metastable Phase Equilibria in Al-Li-Zr Alloys”,

Z. Metallkd., 80, 894-903 (1989) (Equi. Diagram, Theory, Thermodyn., #, 78)

[1990Goe] Goel, N.C., Cahoon, J.R., “The Al-Li-Mg System (Aluminum-Lithium-Magnesium)”, Bull.

Alloy Phase Diagrams, 11, 528-546 (1990) (Equi. Diagram, Review, #, *, 25)

[1990Sau] Saunders, N., “A Review of Thermodynamic Assessment of the Al-Mg and Mg-Li

Systems”, Calphald, 14, 61-70 (1990) (Equi. Diagram, Theory, Thermodyn., #, 78)

[1991Mos] Moser, Z., Agarwal, R., Sommer, F., Predel, B., “Calorimetric Studies of Liquid Al-Li-Mg

Alloys”, Z. Metallkd., 82, 317-321 (1991) (Experimental, Thermodyn., 9)

[1993Gho] Ghosh, G., “Aluminium-Lithium-Magnesium”, MSIT Ternary Evaluation Program, in



Assessment, 48)

[1993Nii] Niikura, A., Tsai, A.P., Inoue, A., Masumoto, T., Yamamoto, A., “Novel Face-Centered

Icosahedral Phase in Al-Mg-Li System”, Jpn. J. Appl. Phys., 32, L1160-L1163 (1993)


[1994Hos] Hosoda, H., Sato, T., Tezuka, H., Mishima, Y., Kamio, A., “A Substitution Behavior of

Additional Elements in the L12-Type Al3Li Metastable Phase in Al-Li Alloys” (in

Japanese), J. Jpn. Inst. Met., 58(8), 865-871 (1994) (Theory, 26)

[1994Kra] Kramer, L.S., Langan, T.J., Pickens, J.R., “Development of Al-Mg-Li Alloys for Marine

Applications”, J. Mater. Sci., 29, 5826-5832 (1994) (Experimental, Equi. Diagram, 23)

[1994Tsa] Tsai, A.P., Yamamoto, A., Niikura, A., Inoue, A., Masumoto, T., “Structural Model of a

Face-Centered Icosahedral Phase in Al-Mg-Li Alloys”, Philos. Mag. Lett., 69, 343-349


[1997Del] Dell'Acqua, G., Krajci, M., Hafner, J., “Face-Centered Al-Mg-Li Alloys: a Free-Electron

Quasicrystal”, J. Phys.: Condensed Matter, 9, 10725-10738 (1997) (Crys. Structure,

Theory, 46)

[1997Hwa] Hwang, Y.H., Han, C.H., Kim, Y.W., Cho, B.J., Kim, D.H., Hong, C.P., “Effects of Heat

Treatment on the Mechanical Properties in Squeeze Cast Mg-Li-Al Alloys” (in Korean),

J. Korean Inst. Met. Mater., 35(12), 1653-1659 (1997) (Experimental, 15)

[1997Kim] Kim, Y.W., Hwang, Y.H., Park, T.W., Kim, D.H., Hong, C.P., “Precipitation Behavior of

and During Heat Treatment in Squeeze Cast Mg-Li-Al Alloys” (in Korean), J. Korean

Inst. Met. Mater., 35(12), 1609-1615 (1997) (Experimental, 9)

[1997Su] Su, H.-L., Harmelin, M., Donnadieu, P., Baetzner, C., Seifert, H.J., Lukas, H.L., Effenberg,

G., Aldinger, F., “Experimental Investigation of the Mg-Al Phase Diagram from 47 to 63

at.% Al”, J. Alloys Compd., 247, 57-65 (1997) (Crys. Structure, Experimental, Equi.

Diagram, #, *, 20)

[1998Cho] Cho, B.J., Kim, D.H., Hong, C.P., “Formation and Growth of Widmanstaetten HCP phase

in Mg-Li-Al Alloy” (in Korean), J. Korean Inst. Met. Mater., 36(5), 647-654 (1998)

(Experimental, 11)

[1998Don] Donnadieu, P., Harmelin, M., Seifert, H.J., Aldinger, F., “Commensurately Modulated

Stable States Related to the -Phase in Mg-Al Alloys”, Philos. Mag. A, 78, 893-905 (1998)

(Crys. Structure, Experimental, *, 21)

[1998Lia] Liang, P., Sung, H.-L., Donnadieu, P., Harmelin, M.G., Quivy, A., Ochin, P., Effenberg, G.,


100


MSIT®

Al–Li–Mg

Calculation of the Central Part of the Mg-Al Phase Diagram”, Z. Metallkd., 89, 536-540

(1998) (Equi. Diagram, Experimental, Thermodyn., #, *, 33)




Assessment, 49)

Table 1: Temperature Dependence of Solid Solubility of (Mg) in Three-Phase Fields


Three-Phase Field Temperature [°C] Composition (at.%) References

Al Li

(Mg) + (Li) + 400 6.0

5.2

20.8

19.2

[1980Sch4]

[1977Dri]

300 2.7

1.25

18.7

17.0

[1980Sch4]

[1976Pad]

200 1.3

0.63

0.87

18.0

16.9

16.9

[1980Sch4]

[1977Dri]

[1955Row]

100 0.28 16.9 [1955Row]

(Mg) + + 400 11.0

7.8

11.7

11.3

[1980Sch4]

[1977Dri]

300 4.7

3.35

10.0

8.7

[1980Sch4]

[1976Pad]

200 3.3

1.45

2.26

8.7

8.1

5.7

[1980Sch4]

[1977Dri]

[1955Row]

100 1.78 3.59 [1955Row]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Al)

660.45

cF4

Fm3m

Cu

a = 404.88 pure Al at 24°C [V-C]

(Li)

180.6

cI2

Im3m

W

a = 351.0 pure Li at 25°C [V-C]

(Mg)

650

hP2

P63/mmc

Mg

a = 320.89

c = 521.01

pure Mg [V-C]

, LiAl

700

cF16

Fd3m

NaTl

a = 637.0 [V-C, 1982McA], at 50 at.% Li

45 to 55 at.% Li

Li3Al2 520

hR15

R3m

Bi2Te3

a = 450.8

c = 1426.0

[V-C, 1982McA]

60 to 61 at.% Li

101


MSIT®

Al–Li–Mg

Table 3: Invariant Equilibria

, Li9Al4, (h)

330 - 275

mC26

C2/m

Li9Al4

a = 1915.51

b = 542.88

c = 449.88

= 107.67°

[V-C, 1982McA]

', Li9Al4, (r)

275

- - [1982McA]

, Mg2Al3 452

cF1168

Fd3m

Mg2Al3

a = 2816 to 2824 60-62 at.% Al [2003Luk]


[2003Luk]

, Mg17Al12

< 458

cI58

I43m

Mn

a = 1054.38 at 41.4 at.% Al [V-C2]

39.5 to 51.5 at.% Al [2003Luk]

, Mg23Al30

410 - 250

hR159

R3

Mn44Si9

a = 1282.54

c = 2174.78

54.5 to 56.5 at.% Al [2003Luk]

1, LiMgAl2 c*456 a = 2031.0 [1972Sam, V-C]

Reaction T [°C] Type Phase Composition (at.%)

Al Li Mg

L + (Al) + 1 536 U1 L

(Al)

1

66.0

53.5

81. 5

54.2

19.4

40.7

12.0

34.5

14.6

5.8

6.5

11.3

L + 1 (Al) + 483 U2 L

1

(Al)

61.5

54.5

79.3

48.3

10.8

31.0

8.4

16.4

27.7

14.5

12.3

35.3

L + 1 + 464 U3 L

1

39.8

51.2

45.5

42.5

20.1

34.4

40.8

18.6

40.1

14.4

13.7

38.7

L + (Al) + 458 P1 L

(Al)

60.5

80.7

51.9

60.5

6.0

2.9

10.7

7.2

33.5

16.4

37.4

32.3

L + (Li) (Mg) + 436 U4 L

(Li)

(Mg)

23.9

0.2

7.9

39.5

29.3

37.5

20.2

44.5

46.8

62.3

71.9

16.0

L (Mg) + + 418 E1 L

(Mg)

19.0

10.2

37.7

41.6

20.6

12.6

17.7

42.5

50.4

76.2

44.6

15.9

L + Li3Al2 + (Li) 411 U5 L

Li3Al2(Li)

12.6

39.4

30.5

0.2

61.0

50.5

67.2

63.0

26.4

10.1

2.3

36.8

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

102


MSIT®

Al–Li–Mg

0 10

Al2

Li, mass%

10

20

(Al)

300°C

Mg,m

ass

%

430°C400°C

200°C

3

(Al)+ +� �1

(Al)+ +� �1

(Al)+ +� �

Fig. 1: Al-Li-Mg.

The solid solubility of

(Al) at different

temperatures

70 80 90

0

100

200

300

400

Li 38.20Mg 61.10Al 0.70

Li 0.00Mg 99.10Al 0.90Mg, at.%

Te

mp

era

ture

, °C

(Mg)(Mg)+(Li)

70 80 90

0

100

200

300

400

Li 38.20Mg 61.10Al 0.70

Li 0.00Mg 99.10Al 0.90Mg, at.%

Te

mp

era

ture

, °C

(Mg)(Mg)+(Li)

Fig. 2: Al-Li-Mg.

The (Mg)/(Mg)+(Li)

phase boundary at a

constant Al content of

1.0 mass%

103


MSIT®

Al–Li–Mg

70 80 90

0

100

200

300

400

Li 38.30Mg 60.40Al 1.30

Li 0.00Mg 98.20Al 1.80Mg, at.%

Te

mp

era

ture

, °C

(Mg)(Mg)+(Li)

10

0

100

200

300

400

Li 3.40Mg 96.60Al 0.00

Li 3.40Mg 83.20Al 13.40Al, at.%

Te

mp

era

ture

, °C

(Mg)

(Mg)+Mg17Al12(γ)

Fig. 3: Al-Li-Mg.

The (Mg)/(Mg)+(Li)

phase boundary at a

constant Al content of

2.0 mass%

Fig. 4: Al-Li-Mg.

The (Mg)/(Mg)+(Li)

phase boundary at a

constant Li content of

1.0 mass%

104


MSIT®

Al–Li–Mg

10

0

100

200

300

400

Li 12.70Mg 87.30Al 0.00

Li 13.00Mg 70.30Al 16.70Al, at.%

Te

mp

era

ture

, °C

(Mg)

(Mg)+Mg17Al12(γ)

10

0

100

200

300

400

Li 6.70Mg 93.30Al 0.00

Li 6.80Mg 80.20Al 13.00Al, at.%

Te

mp

era

ture

, °C

(Mg)

(Mg)+Mg17Al12(γ)

Fig. 6: Al-Li-Mg.

The (Mg)/(Mg) +

Mg17Al12 ( ) phase

boundary system at a


4.0 mass%

Fig. 5: Al-Li-Mg.

The (Mg)/(Mg) +

Mg17Al12 ( ) phase

boundary at a


2.0 mass%

105


MSIT®

Al–Li–Mg

Fig

. 7

:

Al-

Li-

Mg

. R

eact

ion s

chem

e fo

r th

e so

lidif

icat

ion o

f A

l-L

i-M

g a

lloys

l (

Al)

+ η

600

e 1

L +

ητ 1

ca.

545

p1

l +

η L

i 3A

l 2

520

p2

Lτ 1 +

γ4

85

e 3

L +

η (

Al)

+ τ

15

36

U1

L +

τ1

(A

l) +

γ4

83

U2

L +

(A

l) +

τ1 L

+ (

Al)

+ γ

L (L

i) +

η4

80

e 4

l (

Li)

+ (

Mg)

588

e 2

L +

τ1

η +

γ4

64

U3 L +

(A

l) +

γβ

458

P1

L+

η +

γ

L+

β +

γl

(A

l) +

β4

50.5

e 5

lβ

+ γ

44

9.5

e 6

l γ

+ (

Mg)

436

e 7

β +

γε

410

p3

εβ

+ γ

250

e 8

l (

Li)

+ δ

167

e 10

l +

Li 3

Al 2

δ'3

35

p5

L +

Li 3

Al 2

+ (

Li)

δ,

δ'

ca.

350

P2

L +

η L

i 3A

l 2 +

(L

i)4

11

U5

L +

Li 3

Al 2

+ (

Li)

L (

Mg)

+ γ

+ η

418

E1

L +

(L

i)

(M

g)

+ η

436

U4

L +

(M

g)

+ η

(Al)

+ γ

+ τ 1

η +

γ +

τ 1

(Al)

+ γ

+ β

(Al)

+ η

+ τ 1

(Mg

) +

η +

γ

(Li)

+ (

Mg)

+ η

(Li)

+ L

i 3A

l 2 +

(δ,

δ')

η +

Li 3

Al 2

+ (

Li)

L +

(A

l) +

β

Al-

Li

Li-

Mg

Al-

Li-

Mg

Al-

Mg

106


MSIT®

Al–Li–Mg

20

40

60

80

20 40 60 80

20

40

60

80

Li Mg


Axes: at.%

600

550

500

650

600

550

500

450

600550

500

450350 400

U1

U e5

P1

e6

e7

e2

e10

p5

p2

e1

p1

e3

U3

E1

U4

e4

U5

P2

τ1

δ

Li3Al

2

LiAl(η)

(Al)

Mg2Al

3(β)

Mg17

Al12

(γ)

(Mg)

(Li)

2

20

40

60

80

20 40 60 80

20

40

60

80

Li Mg


Axes: at.%

(Al)

β

γ

(Mg)(Li)

ητ

1

450.5596

450.5458,

483, U2

436

436

588 588

418, E1

436, U4

411, U5

520

596

536,

U1

464, U3

449.5

449.5P

1

Fig. 8: Al-Li-Mg.

Liquidus surface

Fig. 9: Al-Li-Mg.

Solidus surface

107


MSIT®

Al–Li–Mg

20

40

60

80

20 40 60 80

20

40

60

80

Li Mg


Axes: at.%

(Al)

β

γ

(Mg)(Li)

η

τ1

η+γ+(Mg)

(Li)+(Mg)+η

(Al)+τ1+η

η+τ1+γ

(Al)+τ1+γ

(Al)+γ+β

L

Li3Al

2

L+

(Li)+

Li

3 Al

2

Li3Al

2+η+(Li)

εβ+γ+ε

20

40

60

80

20 40 60 80

20

40

60

80

Li Mg


Axes: at.%(Al)

β

γ

(Mg)(Li)

η

τ1

η+γ+(Mg)

(Li)+(Mg)+η

(Al)+τ1+η

η+τ1+γ

(Al)+γ+β

(Al)+τ1+γ

β+γ+εε

L

Li3Al

2

δ'

Li3Al

2+δ'

+(Li)

(Li)

+L

+δ'

Li3Al

2+η+(Li)

Fig. 10: Al-Li-Mg.


400°C

Fig. 11: Al-Li-Mg.


300°C

108


MSIT®

Al–Li–Mg

20

40

60

80

20 40 60 80

20

40

60

80

Li Mg


Axes: at.%(Al)

β

γ

(Mg)

(Li)

η

τ1

η+γ+(Mg)

η+(Li)+(Mg)

η+τ1+γ

(Al)+τ1+η

(Al)+γ+β

(Al)+τ1+γ

εβ+γ+ε

Li3Al

2

δ

L

Li3Al

2+η+(Li)

Li3Al

2+δ+(Li)

δ+(Li)+L

Fig. 12: Al-Li-Mg.


200°C

109


MSIT®

Al–Li–Si

Aluminium – Lithium – Silicon

Oksana Bodak

Literature Data

The first studies on Al-Li-Si were published in 1926 and the first reviews were made by [1991Goe] and[1995Pav]. Thermal analysis and metallographic techniques have been used to construct a partial liquidusprojection for Al rich alloys, Fig. 1, using the data of [1977Dri, 1984Han]. Although topologically similarin the sense that both groups reported the presence of a pseudobinary eutectic reactions L (Al)+ 1, and twoternary eutectic reactions L (Al)+(Si)+ 1 and L (Al)+LiAl+ 1, the results of both groups differ largely inlocating the invariant points and in the liquidus isotherms for the primary 1 region. There is alsoconsiderable uncertainty with regard to the composition of the ternary compound 1 and to the extension ofits homogeneity range.Historically [1926Ass] was the first to study Al-Li-Si alloys with a view to improve their mechanicalproperties, by ageing between 25 and 525°C. He deduced that the section Al-Li3Si (Li13Si4?) was apseudobinary section, which is understandable as he was not aware that there is an additional compound, 1.The first report of a ternary compound [1949Boo1] merely stated that the addition of sufficient Li to Al-Sialloys revealed a new phase LixAlySiz. Much more details were revealed by [1949Boo2]. Alloys from 1 to20 mass% Si were thermally analyzed at cooling rates of 8 K·min-1, remelted under a 50 KCl, 50 LiCl fluxwith the addition of 1 mass% Li and the thermal analysis repeated. A ternary eutectic reaction was locatedat 569°C. For hypereutectic Al-Si alloys additions of >1 mass% Li gave a ternary compound as the primaryphase. The most significant finding concerned the composition of the ternary compound. An alloy with7.4Li-11.9Si (at.%) was shown by metallography to contain primary ternary compound. This phase wasextracted with hot HCl, the extract was dried and chemically analyzed as 44.1Li-29.6Si (at.%). Thiscomposition is close to the formula Li3Al2Si2 for 1. In later work [1976Kad] showed a pseudobinaryeutectic e7 L (Al)+Li3Al2Si2, Fig. 1. Using electron probe microanalysis combined with the nuclearmicroprobe [1987Deg] showed that the primary phase in as cast alloy containing 16.1Li-6.6Si (at.%) wasLi3Al2Si2. The crystal structure was not established. However, [1960Now, 1976Sch, 1984Han] refer to theternary compound as LiAlSi, the lattice parameter of which are very close to the LiAlSi after [1960Now].The designation of the ternary compound 1 as LiAlSi stems from [1960Now] who prepared about 30 alloysfrom the elements by heating them in sealed (welded) Fe crucibles at 900-1000°C for 2 h. Practically noattack was observed on the crucible. Examination of the alloys, presumably in their cast state, was solely byX-ray powder diffraction analysis. A cubic phase with a = 594 pm was found at the composition “LiAlSi”.The new phase with a lattice parameter a = 613 pm was detected at the composition "Li2Al2Si". With lowerSi contents, on the section “LiAlSi” - LiAl, at a composition of 43.5Li-13Si (at.%), the X-ray examinationproved that the alloy was heterogeneous. At the composition “Li2AlSi” the cubic phase had a latticeparameter a = 612 pm. Equilibria in the solid state were studied in alloys containing less then 8.0 at.% of Liand less then 12.0 at.% Si. Aluminum (99.99 mass%), lithium (99.8 mass%), and silicon of semiconductorpurity were used as initial materials. [1976Kad] who used thermal analysis and metallographic techniquesto study the equilibria in Al-rich alloys showed a wide two-phase region in which (Al)+ 1 coexist andtherefore a wide homogeneity region for 1.In [1995Pav] it is accepted that the ternary compound 1 is based on the formula Li3Al2Si2, as shownindependently by [1949Boo2, 1987Deg], with a homogeneity region that includes the composition “LiAlSi”and “Li2Al2Si”. There is disagreement on the composition “Li2AlSi”; [1960Now] reports it as a cubic phasewithin the homogeneity region of 1, whereas [1978Ble] regards it as a single phase with cubic structure,different from “LiAlSi”, with a = 606.1 pm and a density of 1.92 g cm-3. These data were measured fromsamples prepared under optimum conditions, reacting elements (99.98 Li, 99.999 Al and Si mass%) for 5 dat 500-600°C followed by slow cooling to room temperature. [1978Ble] indicate that a phase with thestoichiometry Li2AlSi did not form. [1992Pav1] studied the system at 200°C and did not detect the Li2AlSicompound and interpreted the ternary compound 1 as LiAlSi with no homogeneity range. All the

110


MSIT®

Al–Li–Si

conflicting data reported by [1960Now, 1974Boc, 1976Sch, 1978Ble, 1992Pav1] rely exclusively on X-raydiffraction analysis of cubic phases with lattice parameters varying from a = 593 pm to a = 612 pm.Experimental difficulties with Li rich alloys have precluded the use of thermal analysis and metallographictechniques.Until results from new techniques are available for these alloys it is concluded that the only ternary phasein alloys containing 50 at.% Li is the compound 1.In the studies of alloys containing >50 at.% Li [1978Ble] reported the presence of the ternary compoundLi5.3Al0.7Si2 with 1 formula unit in the elementary cell. This compound showed superlattice reflections,which were ascribed to the presence of a phase with the same composition containing 3 formula units in theelementary cell and having an “a” axis enlarged by . Due to the reactivity of the alloys it was not possibleto use high temperature X-ray diffraction analysis to determine whether Li16Al2Si6, with 3 formula units,was a low temperature polymorph to Li5.3Al0.7Si2. [1992Pav1] prepared ternary alloys from 98.2 Li,99.9998 Al and 99.999 mass% Si by arc-melting in purified Ar atmosphere under 1.01·105 Pa pressure. Thealloys were annealed for 240h at 200°C in Ta containers and examined by X-ray diffraction analysis. Theternary compound Li5.3Al0.7Si2 [1978Ble] was confirmed. A ternary compound Li12Al3Si4 was alsoobserved. This compound probably corresponds to a phase called W in [1978Ble]. It has a lattice parametera of 612 to 615 pm which is about a the lattice parameter of Li12Al3Si4, Table 1.Further studies on the phase relations and crystal structures of the compounds were made by [2000Kev,2001Kev, 2001Gro].To clarify the relations among the ternary phases [2001Kev] prepared three series of samples made fromaluminum powder (99.8 mass%, Alfa), lithium bulk material (99.9 mass%, Chemetall, Frankfurt), andsilicon chips (99,9998 mass%, Wacker) as starting materials. The first samples were prepared byarc-melting in purified argon atmosphere. Due to high weight losses (5-10 mass%) by arc-melting,levitation melting under purified argon was performed for most of the alloys. Samples were packed into Tacontainers and sealed in silica ampoules. The annealing was carried out at 250°C for up to 1 month. Theresults for alloys of 15 compositions in the range of 30 to 60 at.% Li, 20 to 50 at.% Si, and 10 to 50 at.% Alare reported by [2001Kev]. Alloys were powdered and investigated using an X-ray powder diffractometerSiemens D-5000 with CoK radiation. The mechanically extracted single crystals of the new ternary phaseswere also investigated using electron microscope Leitz-AMR 1600T with EDX-detector for thedetermination of composition. The 1 and 2 phases are confirmed and a new phase of the Al3Li8Si5composition designated as 3 is found. The other ternary phases reported earlier are assumed to bemetastable. The isothermal section at 250°C is presented.[2001Gro] investigated the ternary Al-Li-Si alloys by differential thermal analysis. Melting temperatureswere established for the three ternary compounds LiAlSi ( 1), Li5.3Al0.7Si2 ( 2), and Li8Al3Si5 ( 3).Additionally selected ternary alloys were also studied by DTA. These results were combined with the phaserelations examined in [2001Kev]. Using these data together with some of the available information from theliterature the ternary phase diagram was calculated applying the Calphad method. The thermodynamicmodel of the ternary system was built by extrapolating the thermodynamic data of the binary subsystemsinto the ternary. The liquid phase and (Al) were modeled by a simple regular solution model without anyternary interaction parameter. The three experimentally found ternary phases were modeled asstoichiometric phases although there is a homogeneity range confirmed for 1. The phase transformationtemperatures found by [2001Kev] were used to fit Gibbs energy functions for the ternary phases. Theresulting calculation reproduced the measured DTA data quite well, the model parameters, however werenot cited [2001Gro]. This work also presents a number of isothermal sections calculated at 250, 590, 597,605, 700, 800°C, the liquidus surface and a set of invariant equilibria, but does not give the compositionsof the phases. The latest results, published by [2003Spi], puts newly questions on the composition of the compounds inthe Al-Li-Si system. The authors synthesized the Li12Al3Si4 compound which according to [1992Pav1,1992Pav2, 1992Pav3, 1996Dmy] does exist, and which categorically is denied to exist by [2001Kev,2001Gro]. The alloy was prepared in a tantalum tube weld-sealed under an argon atmosphere. This tube wasprotected from air by a silica jacket sealed under vacuum. The mixture was heated for 10 h at 950°C in avertical furnace and shaken several times for homogenization. It was then cooled down at a rate of 6 K h-1

3

3

111


MSIT®

Al–Li–Si

for crystal growth. The product of the reaction appeared to be not quite homogeneous, but containedpredominantly black and well-crystallized material. A few black crystals were selected and analyzed byatomic absorption flame spectrometry to identify the composition. This analysis led to an Li/Al/Si ratio of1:0.223(2):0.41(1), corresponding to a mean formula of Li14.63Al3.26Si6. The compound could then bere-prepared following this stoichiometry and obtained in practically 100% yield, as confirmed by X-raypowder pattern (m. p. 824°C).The structure of the Li15Al3Si6 compound, determined by X-ray powder and single crystal analysis, agreeswell with data found earlier by [1978Ble], for the compound Li16Al2Si6.Al-Li-Si alloys were investigated in [1994Hos] with the purpose to study the influence of the thirdcomponent, in this case Si, on the precipitation of metastable phase LiAl3 ( `).

Binary Systems

The binary system Li-Si from [Mas2] is accepted. The binary Al-Si system from [2003Luk] and Al-Li from[2003Gro] are accepted.

Solid Phases

The 1, 2 and 3 phases of constant composition are the only stable phases in the system according to[2000Kev, 2001Kev, 2001Gro], who worked with high purity initial materials and under well controlledconditions of the experimental environment. The essential differences in composition of compound withcubic structure reported in early works [1949Boo2, 1960Now, 1976Dri, 1976Kad, 1977Dri, 1984Han,1992Pav1] become understandable after results of [2001Kev]. In this work, in addition to 1 phase, 3phase, also cubic, but with larger cell parameter and closely-related crystal structure, has been found (Table1). The resemblance of the crystal structures of 1 and 3 phases and limitations of the film-method used forthe determination of crystal structure in early works can be the reason of noticed inaccuracies. All threeternary phases are proposed to melt congruently: 1 at 811°C, 2 at 793°C and 3 at 833°C. The existence of a hexagonal ternary 4 phase, found by [1978Ble, 2003Spi] and the cubic 5 found by[1978Ble, 1992Pav2] need to be confirmed. These phases possibly are stabilized by impurity of othercomponents, contained in the initial metals.


Seven pseudobinary eutectics exist in the system according to computation [2001Gro]. Unfortunately theircompositions are not given. The existence of a pseudobinary section extending between the (Al) solidsolution and 1 is well established experimentally by [1976Dri, 1976Kad, 1977Dri, 1984Han]. Theinvariant curve for the liquid phase undergoes a maximum at 635°C for an invariant eutectic reactionaccording to [1976Kad, 1977Dri], at ~630°C according to [1984Han] and at 657°C according to [2001Gro].There is disagreement on the reported composition of the eutectic maximum, Fig. 1. [1976Kad, 1977Dri]give a vertical section from the Al corner to 17.5 mass% Li3Al2Si2 with the eutectic composition at 9 mass%Li3Al2Si2 (5.35Li-3.57Si (at.%). [1984Han] quotes a eutectic composition that does not lie on theirmonovariant curve E2´E3´, Fig. 1 the scaled composition converts to 14Li-4Si (at.%) for e7´. As shown inFig. 1 the eutectic maximum found by [1984Han] at ~630°C lies very near to the 670°C isotherm given by[1977Dri]. The discrepancies between [1976Kad, 1977Dri, 1984Han] can only be resolved by furtherinvestigation.


From the eutectic maximum, e7 or e7´ in Fig. 1, monovariant curves descend to a ternary eutectic E2´ or E2´´and to a ternary eutectic E3´ or E3´´, respectively. According to [1976Kad, 1977Dri] E2 has the composition28.3Li-1.5Si (at.%) whereas [1984Han] place E´2 at 31.6Li-0.8Si (at.%). The temperature of the reactionwas given as 595°C [1976Kad, 1977Dri] and 592°C [1984Han]. The liquidus and the eutectic compositionfor binary Al-Li alloys given by [1976Kad, 1977Dri] agree more closely with [1989Che] than do the valuesfound by [1984Han]. For example the binary Al-Li eutectic composition is quoted as 25.8 at.% Li

112


MSIT®

Al–Li–Si

[1989Che], 27.1 at.% Li [1976Kad, 1977Dri] and 30.2 at.% Li [1984Han]. On this basis the ternary eutecticpoint E2´´ is preferred above E2´ reported by [1984Han]. The composition of the liquid phase in the secondternary eutectic reaction shows similar differences. [1976Kad, 1977Dri] and [1984Han] place E3 atdifferent compositions; respectively at 0.2Li-11.1Si (at.%) (shown as E3´´ in Fig. 1) and at 5.3Li-12.8Si(at.%) (shown as E3´ in Fig. 1). The ternary eutectic temperature is quoted by [1976Kad, 1977Dri] to be565°C. [1984Han] locates it at 575°C and according to [2001Gro] the reaction happens at 577°C. The dataof [1984Han] for the binary Al-Si eutectic agree well with that given by [2003Luk] whereas that of[1976Kad, 1977Dri] do not. The data of [1984Han] and the composition/temperature of E3´ are preferredabove those of [1976Kad, 1977Dri]. On the basis of the assessed experimental data a partial reaction schemeis given in Fig. 2. The calculated invariant equilibria after [2001Gro] are listed in Table 2.The optimization of [2001Gro] revealed a contradiction between the higher melting temperature of 3(compared to 1) and the eutectic E1: L (Al)+LiAl+ 1 reported by [1976Kad, 1976Dri, 1984Han]. A highermelting phase 3 will always result in a tie line between 3 and (Al) at higher temperature. Therefore aeutectic between (Al), LiAl, and 1 will not occur. On the other hand, the four phases, 1, 3, LiAl, and (Al),found in as-cast (not equilibrated) alloys near to 1 and 3 give a hint for an invariant reaction which maychange the tie line of 3+(Al) to LiAl+ 1. In fact, in the calculation an invariant reaction 3+(Al) LiAl+ 1at 591°C emerges by fitting the parameter for 3 and 1 to the measured melting temperatures and to theexperimentally observed phase equilibria at 250°C. This final version of the thermodynamic data setreproduced all experimental results of [2001Kev]. However, the calculated liquidus surface of 1 extendsmuch closer to the Al corner than reported by [1984Han] and somewhat closer than given by [1976Kad].Figure 3 illustrates the discrepancies between the different reports shown with dashed lines [1976Kad] anddotted lines [1984Han] and the calculation [2001Gro] of the partial liquidus surface shown with solid lines.As discussed above, a eutectic E2´´´: L LiAl+(Al)+ 1 does not take place in this calculation. However, atnearly the same temperature as given for E2´´ by [1976Kad] an invariant reaction, E2: L LiAl+(Al), 3 ispresent in the calculation. It was concluded that the ternary eutectic with 3 instead of 1 describes thecorrect equilibrium.

Liquidus Surface

As follows from Fig. 3 there is discrepancy between experimental data of different authors. The calculatedliquidus surface given in Fig. 4 after [2001Gro] differs from both experimental series shown in Fig. 3. Asfollows from three previous chapters additional investigations for liquidus surface are necessary.

Isothermal Sections

An isothermal section at 550°C was published by [1976Dri] and one at 500°C by [1977Dri]. [1976Kad] hasdrawn four-phase eutectic planes at 595°C and 565°C and four vertical sections, along 5 mass% Li, along92 mass% Al, along 2 mass% Si, and one section along Al- 1, up to 17.5 mass% Li3Al2Si2. [1976Dri] usedan extended annealing schedule, involving 30 h homogenization at 400°C of the cast ingots followed bydeformation of 70 % with different annealing, 200 h at 550°C or 200 h at 550°C, plus subsequent 400 h at500°C; or 200 h at 550°C plus subsequent 1000 h at 200°C. All annealing procedures terminated with waterquenching of the samples. Thermodynamically the resulting data for the combined solubility of Li and Siin Al at 550, 500 and 200°C are not feasible. These data were used by [1977Dri] to produce a 500°Cisothermal section confined to Al contents above 88 mass%. Plotting in a common scale data from[1976Dri] at 550°C, [1977Dri] at 500°C and [1976Kad] at 595 and 565°C gives an impression of the widthof the (Al)+ 1 phase region represented by this group of workers. Figure 5 summarizes the data andextrapolates the boundary (Al)- 1 tie lines after [1976Kad] to the 1 “composition line” between “LiAlSi”and “Li2AlSi”. See also the discussion in “Introduction”. [1976Kad] did not determine any phaseboundaries below the two ternary eutectic temperatures. No check can be made to compare their verticalsections with the isothermal sections at 550 and 500°C. Comparison of the delineation of the 595 and 565°Cternary eutectic planes [1976Kad] with their published vertical sections shows reasonable agreement for the(Al)- 1 tie line at 565°C but substantial disagreement for the (Al)- 1 tie line at 595°C. In Fig. 5 the tie linesgiven by [1976Kad] at 595 and 565°C have been preferred to those derived from vertical sections.

113


MSIT®

Al–Li–Si

As can be seen from Fig. 5 there is some difference between the data reported by [1976Dri, 1977Dri,1976Kad]. The isothermal section at 600°C assessed by [1991Goe] is based on the above mentionedexperimental data. The systematic investigations of isothermal sections given in [1992Pav1] and [2001Kev](Fig. 6) indicate the punctual composition for the compound 1 (LiAlSi). The difference in compositiongiven in works [1976Dri, 1976Kad, 1977Dri] is connected with the existence of 1 and 3 phases withclosely-related crystal structures. This is clearly described in [2001Kev].The equilibria at different constant temperatures, between 590 and 800°C are given in Figs. 7 to 11 ascalculated by [2001Gro]. At 800°C (Fig. 11) only two ternary phases, 3 and 1, are present. One hundreddegrees lower (Fig. 10) the third ternary phase, 2 appears together with the binary phases Li13Si4 andLi7Si3. The liquid phase extends along the Al-Li edge up to the binary Li-Si eutectic, with little extensioninto the ternary. At 605°C (Fig. 9) the (Al) solid solution is in equilibrium with 1. The need to reconcilethe high melting point of 3 with the solid state LiAl+ 1 equilibrium were resolved by [2001Gro] in a seriesof three nonvariant equilibria: U5: L+ 1 (Al)+ 3 at 600°CE2: L LiAl+(Al)+ 3 at 596°CU6: 3+(Al) LiAl+ 1 at 591°C.The high melting point of 3 gives a tie line between 3 and an Al-rich liquid, which is also in equilibriumwith 1 at 605°C (Fig. 9). The reaction U5 transforms this tie line, L+ 1, into a tie line 3+(Al) shown inFig. 8. The heat evolution of U5 is suspected to be slow, because a substantial amount of 1 would have tobe consumed in this cross-reaction. At 596°C the liquid decomposes by the eutectic reaction E2 to formLiAl+(Al)+ 3. At 591°C the (Al)+ 3 tie-line transforms into the 1+LiAl equilibrium which is stable downto room temperature, Figs. 7 and 6. As a result from the reaction U6 (Übergangsreaktion) the triangle

1+LiAl+(Al) appears in Fig. 7, describing a three phase field which is well supported by literature data[1976Kad, 1976Dri, 1984Han]. These results, however, would be different if the experimentally foundhomogeneity range for 1 phase is taken into account.


Lithium is an important alloying element for weight saving in conventional aluminium alloys. Lithiumadditions to Al-Si increases the strength and elasticity of alloy, with silicon increasing in particular theirhardness [1976Kad]. The improvement of the physical properties by adding silicon to aluminium-lithiumalloys is attributed to the formation of lithium silicides. For compositions close to 1 a microhardness of946 kg·mm-2 has been measured. On quenching, the alloys are in an unstable state, supersaturated with silicides which later, during ageing,appear in a highly dispersed form. Although Al-Li-Si alloys are heat treatable, the improvement inproperties is small. The main effect of lithium in Al-Si alloys is the improvement of hardness by thecombined effect of Li and Si [1926Ass]. During microprobe study of structure of alloys with compositionnear E3 the epitaxy between silicon and silicide was observed, leading to the formation of fine siliconestructure [1963Boo]. Additions of aluminium to silicides of lithium increase their stability during hydrolyzein dilute H2SO4 under argon [1974Boc].

References

[1926Ass] Assmann, P., “The Importance of Si for the Mechanical Improvement of Al by Li or Mg”(in German), Z. Metallkd., 18, 256-260 (1926) (Experimental, 7)

[1949Boo1] Boom, E.A., “A New Phase in the Al-Li-Si System” (in Russian), Dokl. Akad. Nauk SSSR,66, 645-646 (1949) (Experimental, 3)

[1949Boo2] Boom, E.A., “Physico-Chemical Investigation of Al-Li-Si Alloys” (in Russian), Dokl.

Akad. Nauk SSSR, 67, 871-874 (1949) (Experimental, 5) [1960Now] Nowotny, H., Holub, F., “Investigations of Metallic Systems with Fluorspar Phases” (in

German), Monatsh. Chem., 91, 877-887 (1960) (Crys. Structure, Experimental, 15)[1963Boo] Boom, E.A., “On the Mechanism of the Modification of Silumin” (in Russian), Dokl. Akad.

Nauk SSSR, 151, 96-97 (1963) (Experimental, 5)

114


MSIT®

Al–Li–Si

[1974Boc] Bockelmann, W., Schuster, H.U., “Crystallographic Aspects of Ternary Phases of Li with3B and 4B Elements in Ionic and Non-Ionic Compounds” (in German), Anorg. Allg. Chem.,410, 741-750 (1974) (Crys. Structure, Experimental, 5)

[1976Dri] Drits, M.E., Kadaner, E.S., Kuz’mina, V.I., Turkina, N.I., “Phase Composition of Al-RichAl-Li-Si Alloys”, Russ. Metall., (5), 177-178 (1976), translated from Izv. Akad. Nauk SSSR,

Met., (5), 206-208 (1976) (Equi. Diagram, Experimental, 4)[1976Kad] Kadaner, E.S., Turkina, N.I., Kuz’mina, V.I., “Phase Diagram of Al-Li-Si System in the

Al-Rich Region”, Russ. Metall., (1), 150-153 (1976), translated from: Izv. Akad. Nauk SSSR,

(1), 181-184 (1976) (Equi. Diagram, Experimental, 14)[1976Sch] Shuster, H.U., Hinterhauser, H.W., Schäfer, W., Will, G., “Neutron Diffraction

Investigations of the Phases LiAlSi and LiAlGe” (in German), Z. Naturforsch., 31B,1540-1541 (1976) (Crys. Structure, Experimental, 3)

[1977Dri] Drits, M.E., Bochvar, N.R., Kadaner, E.S., Padezhnova, E.M., Rokhlin, L.l.,Sviderskaya, E.A., Turkina, N.I., Phase Diagrams of al and Mg Systems (in Russian),Abrikosov, N.Kh., (Ed.), Nauka, Moscow, 57-58 (1977) (Equi. Diagram, Review, 4)

[1978Ble] Blessing, J., “Synthesis and Study of Ternary Phases of li with Elements of the 3 and a subGroups”(in German), Ph. D. Thesis, Univ. Cologne,167 pp. (1978) (Experimental, Crys.Structure, 87)

[1984Han] Hanna, M.D., Hellawell, A., “The Liquidus Surface for the Al-Li-Si System from 0 to 20wt.% Li and Si”, Metall. Trans. A, 15A, 595-597 (1984) (Equi. Diagram, Experimental, 6)

[1987Deg] Degreve, F., Dubost, B., Dubus, A., Thorne, N. A., Bodart, F., Demortier, G., “QuantitativeAnalysis of Intermetallic Phases in Al-Li Alloys by Electron, Ion and NuclearMicroprobes”, J. Phys. Colloq., 48, (Suppl. C3), 505-511 (1987) (Experimental, 13)

[1989Che] Chen, S.-W., Jan, C.-H., Lin, J.-C., Austin Change, Y., “Phase Equilibria of the Al-LiBinary System”, Metall. Trans. A, 20A, 2247-2258 (1989) (Equi. Diagram, Thermodyn.,Experimental, 59)

[1991Goe] Goel, N.C., Cahoon, J.R., “Tha Al-Li-Si (Aluminium-Lithium-Silicon)”, J. Phase Equilib.,12(2), 225-230 (1991) (Equi. Diagram, Review, 9)

[1992Pav1] Pavlyuk, V.V., Bodak, O.I., Dmytriv, G.S., “Interaction of Components in Li-(Mg, Al)-SiSystems” (in Russian), Ukr. Khim. Zh. (Russ. Ed.), 58, 735-737 (1992) (Equi. Diagram,Experimental, #,6)

[1992Pav2] Pavlyuk, V.V., Bodak, O.I., “The Crystal Structure of Li12Mg3Si4 and Li12Al3Si4Compounds” (in Russian), Neorgan. Mater., 28(5), 988-990 (1992) (Crys. Structure,Experimental, 3)

[1992Pav3] Pavlyuk, V.V., Dmytriv, G.S., Starodub, P.K., “Crystal Structure of the Compounds of theLi-M-X (M = Mg, Al; X = Si, Ge, Sn) Systems” (in Russian), VI Conf. Cryst. Chem. Inorg.

Coord. Compounds, L’viv (Abstact), 210 (1992) (Crys. Structure, Experimental)[1994Hos] Hosoda, H., Sato, T., Tezuka, H., Mishima, Y., Kamio, A., “Substitution Behavior of

Additional Elements in the L1(2)-Type Al3Li Metastable Phase in Al-Li Alloys” (inJapanese), J. Jpn. Inst. Met., 58(8), 865-871 (1994) (Crys. Structure, Thermodyn.,Theory, 26)

[1995Pav] Pavlyuk, V., Bodak, O., “Aluminium-Lithium-Silicon”, MSIT Ternary EvaluationProgram, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science InternationalServices GmbH, Stuttgart; Document ID: 10.16694.1.20, (1995) (Crys. Structure, Equi.Diagram, Assessment, 15)

[1996Dmy] Dmytriv, G.S., “Phase Equilibria and Crystal Structure of Compounds in Mg-Li-Si,Ca-Li-{Si, Ge}, Al-Li-{Si, Ge, Sn}, Zn-Li-{Al, Sn}”, Summary of the thesis for kandidate

science degree, 1-23 (1996) (Crys. Structure, Equi. Diagram, Experimental, 10)[2000Kev] Kevorkov, D., Gröbner, J., Schmid-Fetzer, R., “Experimental Investigations and

Thermodynamic Calculation of the Ternary Al-Li-Si Phase Diagram”, Proc. Disc. Meet.

Thermodyn. Alloys, 27 (2000) (Thermodyn., Abstract)

115


MSIT®

Al–Li–Si

[2001Gro] Groebner, J., Kevorkov, D., Schmid-Fetzer, R., “The Al-Li-Si System. 2. ExperimentalStudy and Thermodynamic Calculation of the Polythermal Equilibria”, J. Solid State Chem.,156, 506-511 (2001) (Equi. Diagram, Thermodyn., Experimental, Calculation, 12)

[2001Kev] Kevorkov, D., Groebner, J., Schmid-Fetzer, R., “The Al-Li-Si System. 1. A New StructureType Li8Al3Si5 and the Ternary Solid State Phase Equilibria”, J. Solid State Chem., 156,500-505 (2001) (Crys. Structure, Equi. Diagram, Experimental, 16)

[2003Spi] Spina, L., Tillard, M., Belin, C., “Li15Al3Si6(Li14.6Al3.4Si6), a Compound Displaying aHeterographite-Like Anionic Framework”, Acta Crystallogr., Sect. C: Cryst. Struct.

Commun., C59(2), i9-i10 (2003) (Crys. Structure, Experimental, 9)[2003Gro] Groebner, J., “Al-Li (Aluminium-Lithium)”, MSIT Binary Evaluation Program, in MSIT

Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH,Stuttgart; Document ID: 20.13517.1.20, (2003) (Crys. Structure, Equi. Diagram,Assessment, 21)

[2003Luk] Lukas, H.L., “Al-Si (Aluminium-Silicon)”, MSIT Binary Evaluation Program, in MSIT

Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH,Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 29)


Phase/Temperature Range [°C]

Pearson Symbol/Space Group/Prototype

Lattice Parameters [pm]

Comments/References

(Li)< 180.6

cI2Im3m

W

a = 351.0 pure Li at 25°C[V-C2]

(Al)< 660.45

cF4Fm3m

Cu

a = 404.96 pure Al at 25°C [Mas2]dissolves up to 15 at.% Li and up to 1.5 at.% Si

Li9Al4 ( )< 347 - 275

mC26C2/mLi9Al4

a = 1915.51b = 542.88c = 449.88

= 107.671°

[2003Gro]

Li9Al4 ( ´)< 275

? ? [Mas2]

Li3Al2 ( )< 520

hR15R3m

Li3Al2

a = 450.8c = 1426

[2003Gro]60 to 61 at.% Li[Mas2]

LiAl ( )< 700

cF16Fd3m

NaTl

a = 637 at 50 at.% Li [2003Gro]45 to 55 at.% Li [Mas2]

LiAl3 ( ´)< 190 - ~120

cP4Pm3m

Cu3Au


Li2Si mC12C2/mGe2Os

a = 770b = 441c = 656

= 113.4°

Metastable?[V-C2]

116


MSIT®

Al–Li–Si

Li7Si2 oP36Pbam

Li7Si2

a = 799b = 1521c = 443

Metastable?[V-C2]

Li7Si3< 752

hR7R3m

Li7Si3

a = 443.5c = 1813.4

[Mas2, V-C2]

Li12Si7< 648

oP152Pnma

Li12Si7

a = 861.0b = 1973.8c = 1434.1

[Mas2, V-C2]

Li13Si4< 722

oP34Pbam

Li13Si4

a = 799b = 1521c = 443

[Mas2, V-C2]

Li22Si5< 628

cF432F23

Li22Pb5

a = 1875 [Mas2, V-C2]

Li41Si11 cF416F43m

Cu41Sn5

a = 1871 Metastable? [V-C2]

* 1< 811

cF12F43m

LiAlSi

a = 594

a = 593

a = 593a = 593a = 592.82

at Li0.33Al0.33Si0.33 (LiAlSi) [1960Now]

at Li0.33Al0.33Si0.33

m = 1.95 g cm-3

x = 1.97 g cm-3 [1976Sch]at Li0.33Al0.33Si0.33 [1984Han]at Li0.33Al0.33Si0.33 [1992Pav1]at Li0.33Al0.33Si0.33 [2001Kev]

* 2< 793

hP8P63/mmc

Li5.3Al0.7Si2

a = 435.9c = 813.6

a = 434.10c = 810.52

at Li0.66Al0.09Si0.25 (Li5.3Al0.7Si2)

m = 1.35 g cm-3

x = 1.38 g cm-3 [1978Ble]at Li0.66Al0.09Si0.25 [2001Kev]

* 3< 833

cP16P43m

Li8Al3Si5

a = 611.46

a = 613

a = 612

at Li0.50Al0.19Si031 (Li8Al3Si5)[2001Kev]Li0.42Al0.29Si0.29 (Li3Al2Si2)[1949Boo2, 1976Kad]at Li0.40Al0.40Si0.20 (Li2Al2Si)[1960Now]at Li0.50Al0.25Si0.25 (Li2AlSi)[1960Now]Li0.42Al0.29Si0.29 (Li3Al2Si2) [1987Deg]




Comments/References

117


MSIT®

Al–Li–Si


* 4 hP24P63/mLi15Al3Si6

a = 754.9c = 809.7

a = 755.0c = 813.6

at Li0.63Al0.13Si0.24 (Li15Al3Si6)[2003Spi], not shown on the diagram, stability is not confirmedat Li0.67Al0.08Si0.25 (Li16Al2Si6)[1978Ble], not shown on the diagram, stability is not confirmed

* 5 cI76 a = 1062.0 at Li0.63Al0.16Si0.21 (Li12Mg3Si4)[1992Pav2], not shown on the diagram, stability is not confirmed

Reaction T [°C] Type

L 3 832 congruent

L 1 + 3 809 e1 (max)

L 1 810 congruent

L (Si) + 1 802 e2 (max)

L 2 800 congruent

L 3 + 2 798 e3 (max)

L + 1 3 + (Si) 788 U1

L 2 + Li7Si3 746 e4 (max)

L 2 + Li13Si4 730 e5 (max)

L Li13Si4 + Li7Si3 + 2 727 E1

L + 2 3 + Li7Si3 718 U2

L 3 + LiAl 686 e6 (max)

L + 3 2 + LiAl 679 U3

L (Al) + 1 657 e7 (max)

L + Li7Si3 Li12Si7 + 3 630 D1

L + Li13Si4 Li22Si5 + 2 616 U4

L (Si) + Li12Si7 + 3 604 D2

L + 1 3 + (Al) 600 U5

L LiAl + (Al), 3 596 E2

3 + (Al) LiAl + 1 591 U6

L (Al) + (Si) + 1 577 E3

L + LiAl Li3Al2 + 2 518 U7

L + Li3Al2 Li9Al4 + 2 334 U8

L (Li) + Li22Si5, 2 180 D3

L (Li) + Li9Al4, 2 175 D4




Comments/References

118


MSIT®

Al–Li–Si

10

20

30

70 80 90

10

20

30

Li 35.00Al 65.00Si 0.00

Al

Li 0.00Al 65.00Si 35.00 Data / Grid: at.%

Axes: at.%

620

E3´´

E3´

e7´´

τ1 700

680

660

640 e7´

630620610

LiAl

E2´

E2´´

640650

660

670

(Al)

(Si)

[1977Dri, 1976Kad]

[1984Han]

Fig. 1: Al-Li-Si.

Partial liquidus projection showing the data of [1977Dri] and [1984Han]; numbering of invariant reactions is adapted to [2001Gro]

Fig. 2: Al-Li-Si. Partial reaction scheme from assessed experimental data

Al-Li A-B-C

l (Al) + LiAl

600 e

Al-Li-Si

L (Al)+LiAl+τ1

595 E2

Al-Si

l (Al) + (Si)

577 e

L (Al) + τ1

635 e7(max)

L (Al)+(Si)+τ1

575 E3

L+LiAl+τ1

L+(Si)+τ1

?

(Al)+LiAl+τ1

(Al)+(Si)+τ1

?

119


MSIT®

Al–Li–Si

10

20

30

70 80 90

10

20

30

Li 40.00Al 60.00Si 0.00

Al

Li 0.00Al 60.00Si 40.00 Data / Grid: at.%

Axes: at.%

e7,657

(Al)

e7,635

e7,632

E3,565

E3,577

U5,600E

2,596

E2,592 E

2,595

E3,575

(Si)

τ1

τ3

LiAl

[1984Han][1976Cad]

[2001Gro]

20

40

60

80

20 40 60 80

20

40

60

80

Li Al

Si Data / Grid: at.%

Axes: at.%

(Si)

τ1

τ3

τ2

1300

1200

1100

1000

900

800

775

750

700

600500

Li3Al

2

U7,518

LiAl U3,679

e6

E2,596 U

5,600

e7

E3,577

e2,802

e1,809

e3,798

Li12

Si7

U2,718

Li7Si

3

E1,727

Li13

Si4

Li22

Si5

U4

825

U1,788

(Al)

Fig. 3: Al-Li-Si.

Partial liquidus projection. Comparison between calculation [2001Gro] (solid) and estimations after [1976Kad] (dashed) and [1984Han] (dotted)

Fig. 4: Al-Li-Si.

Calculated liquidus surface

120


MSIT®

Al–Li–Si

20

40

60

80

20 40 60 80

20

40

60

80

Li Al


Axes: at.%

"LiAlSi"

τ1

"Li2AlSi"

(Al)+(Si)+τ1

550500565

(Al)+τ1

595

500

550(Al)+LiAl+τ

1

500 [1977Dri]

550 [1976Dri]

565 [1976Kad]

595 [1976Kad]

Fig. 5: Al-Li-Si.

The (Al)+ 1 phase region after experimental data

20

40

60

80

20 40 60 80

20

40

60

80

Li Al


Axes: at.%

τ1

τ3

τ2

Li12

Si7

Li7Si

3

Li13

Si4

Li22

Si5

Li9Al

4Li

3Al

2LiAl (Al)

(Si)

(Li)

Fig. 6: Al-Li-Si.

Experimental isothermal section at 250°C after [2001Kev]

121


MSIT®

Al–Li–Si

20

40

60

80

20 40 60 80

20

40

60

80

Li Al


Axes: at.%

τ1

τ3

τ2

L

(Al)LiAl

Li12

Si7

Li7Si

3

Li13

Si4

Li22

Si5

L

(Si)Fig. 7: Al-Li-Si.

Calculated isothermal section at 590°C

20

40

60

80

20 40 60 80

20

40

60

80

Li Al


Axes: at.%

τ1

τ3

τ2

(Si)

Li12

Si7

Li7Si

3

Li13

Si4

Li22

Si5

L

L

(Al)LLiAl

Fig. 8: Al-Li-Si.


122


MSIT®

Al–Li–Si

20

40

60

80

20 40 60 80

20

40

60

80

Li Al


Axes: at.%(Si)

τ1

τ3

τ2

L

Li12

Si7

Li7Si

3

Li13

Si4

Li22

Si5

L

LiAl L (Al)

L

Fig. 9: Al-Li-Si.


20

40

60

80

20 40 60 80

20

40

60

80

Li Al


Axes: at.%(Si)

τ1τ

3

τ2

L

L

Li7Si

3

Li13

Si4

Fig. 10: Al-Li-Si.


123


MSIT®

Al–Li–Si

20

40

60

80

20 40 60 80

20

40

60

80

Li Al


Axes: at.%

τ1

τ3

(Si)

L

Fig. 11: Al-Li-Si.


124


MSIT®

Al–Li–Zn

Aluminium – Lithium – Zinc

Oksana Bodak

Literature Data

The research on this system started in 1942 when [1942Wei] established the temperature and composition

of a ternary eutectic in the Zn corner. One year later [1943Bad] investigated the triangle Al-LiAl-Zn by

thermal and microscopic analyses. They found two ternary compounds, 1 and 3, in the LiAl-Zn section

and gave eight vertical and two isothermal sections. The 1 phase has been confirmed by [1963Che] and the

region of its homogeneity has been determined more accurately; however, these authors found 1 to be in

equilibrium with (Zn) neglecting the 3 phase. [1987Dub] and [1989Aud] reported a stable phase to exist

in the vicinity of 4, Li3ZnAl5, not far from the Al rich end of the 1 domain. Metastable icosahedral

quasicrystals are formed in this system by rapid solidification [1986Cas, 1987Che] or as grain boundary

precipitates through solid - solid transformations [1987Cas] with compositions close to the stable 1 phase.

The substitution behavior of additional elements in the L12 type metastable compound of Li3Al ( ´ phase)

was reported in [1994Hos]. In 1995 a critical review was made inn the MSIT evaluation programs, covering

the literature published until 1992, [1995Pav].

Isothermal section of the system at 197°C and crystal structures of compounds were investigated and

published in [1993Pav, 1995Dmy, 1996Dmy, 1999Pav]. Alloys of the Al-Li-Zn system were prepared by

arc-melting pieces of the pure metals (lithium with a purity 98.2 mass%, zinc with a purity 99.98 mass%,

aluminium with a purity 99.99 mass%) under argon atmosphere. The alloys were annealed at 197°C for 400

hours in tantalum containers in evacuated quartz ampoules, quenched in cold water and examined by X-ray

diffraction analysis. There are measurements of the enthalpy of mixing of liquid Al-Li-Zn ternary made by

high temperature mixing calorimeter in the temperature range 456 - 682°C, [1997Kim]. They used their data

in an association model to calculate the thermodynamic mixing functions of the ternary alloys on the basis

of the enthalpy of mixing of the binary systems. Aluminium of purity 99.9%), 99.9% pure lithium and zinc

of 99.999% were used to prepare the alloy samples for these measurements, executed under pure argon gas

at atmospheric pressure.

Binary Systems

For the Al-Li system phase relations are accepted here as reported by [2003Gro]. For the descriptions of the

Al-Zn and Li-Zn phase diagrams the versions given in [Mas2] are accepted.

Solid Phases

The data for the solid phases are given in Table 1. The quasicrystalline phases are formed by rapid

solidification or as grain boundary precipitates by a solid-state reaction in the 1 phase region [1997Kim].

The 1 phase has a high solubility of zinc (16.7-43.3 at.% Zn at 32 at.% Li) and is formed through a

peritectic reaction at higher temperature than the 3 and 4 phases [1997Kim]. According to [1993Pav,

1996Dmy, 1999Pav] three ternary compounds are formed in this system: (a) the 1 phase,

Li1+xZn0.5-1.5Al1.5-0.5 with a large homogeneity range which includes the earlier reported composition 1,

Li26Al6(Zn1-xAlx)49 (b) the 3 phase, LiZn3Al with an unidentified structure and (c) the 4 phase, Li3ZnAl5.

Another compound 2 on the 50 at.% Li section is reported in the work of [1996Dmy].


The section LiAl-Zn shown in Fig. 1 is pseudobinary [1943Bad]. The solidus and the liquidus of the LiAl

phase in Fig. 1 are slightly corrected to agree with the congruent melting point of this phase in the binary

system.

125


MSIT®

Al–Li–Zn

The LiAl - Li2Zn3 section has been reported by [1943Bad] as pseudobinary with continuous solid solubility.

However this is unlikely because LiAl and Li2Zn3 have different crystal structures, and this contradicts to

the existence of the 2 phase proposed by [1996Dmy].


The invariant equilibria established within the triangle Al-LiAl-Zn and in the pseudobinary section LiAl-Zn

[1943Bad] are listed in Table 2. An additional four-phase equilibrium at 423°C has been proposed by

[1943Bad] following the existence of an intermediate phase in the Al-Zn system. However, in the presently

accepted Al-Zn binary this phase does not exist and therefore the invariant reaction at 423°C is eliminated

from the reaction scheme and liquidus surface in this evaluation. The reaction scheme is shown in Fig. 2.

The temperature and the concentration of the ternary eutectic E1 are reported with some uncertainty, 355°C

given by [1943Bad] and 364.25°C. In Fig. 2 and Table 2 the values of [1942Wei] are preferred.

Liquidus Surface

Figure 3 shows the liquidus surface of the Al-LiAl-Zn partial system the diagram given by [1943Bad]. It

had to be amended to match with the accepted binary equilibrium diagrams Al-Zn [Mas2] and Al-Li

[1989Che]. The ternary eutectic is incorporated using the data given by [1942Wei].

Isothermal Sections

Figure 4 shows the amended partial isothermal section Al-LiAl-Zn at 350°C according to [1943Bad]. The

section now is compatible with the accepted binary systems Al-Li [1989Che] and Al-Zn [Mas2] and

coherent with [1996Dmy] for which the phase extends in the ternary system along the 50 at.% Li. The

homogeneity region of the 1 phase follows [1963Che] and may be expressed by the approximate formula

LiZn0.5+xAl1.5+x (0 < x < 0.7). There is no experimental evidence for a large width of the 1 field, so the Li

content may be accepted as 34-35 at.% as given by [1943Bad]. [1963Che] found the 1 phase in equilibrium

with (Zn) neglecting 3. The isothermal section of the system at 193°C according to [1996Dmy] is shown

in Fig. 5. No significant solubilities of Al in binary Li-Zn compounds have been detected.

Thermodynamics

The values H(xC) of liquid Al-Li-Zn alloys were determined at different temperatures along four sections

keeping the concentration ratios of two components constant [1997Kim]: (a) Al0.25Zn0.75-Li, (b)

Al0.50Zn0.50-Li, (c) Al0.70Zn0.30-Li and (d) Al0.75Li0.25-Zn. They are plotted in Figs. 6 and 7. The H

values of the ternary liquid alloys can be obtained by adding the H value of the binary boundary systems:

H(xA/xB = const., xC) = (1 - xC) H(xA/xB = const.) + Hi(xC)(xA/xB = const.)

For the section Al0.25Zn0.75-Li the agreement between the measured and calculated values is within the

experimental error. For other concentration section the experimental H values exhibit more negative

values compared with the calculated ones. These deviations could be caused by a negative contribution to

the enthalpy of mixing due to the presence of additional ternary interactions or additional ternary associates

in the melt which have not been taken into account in the model calculation. The presence of additional

ternary interaction in the liquid state is supported by the existence of at least three ternary intermetallic

phases in this system [1995Pav]. The difference between measured and calculated values of H is shown

in Fig. 8 together with the position of the ternary intermetallic phases. Figure 8 shows that the deviation

amounts to - 3.5 kJ mol-1 in the concentration region where the 1 phase exists, which points to additional

ternary interaction in this concentration region. In the region of the ternary 2 and 3 phase the deviation is

small in comparison to that in the 1 phase region. This indicates that the ternary interactions in these

regions are relatively weak and the influence of the ternary 1 phase is predominant for liquid Al-Li-Zn

alloys.

δi

126


MSIT®

Al–Li–Zn


Al-Li base alloys have received considerable attention as potential lightweight replacements for

conventional Al base alloys in aerospace applications. The addition of 1.8-2.1% Li remarkably alter the

precipitation behavior of the Al-Cu-Mg-Zn alloys which are the highest strength aluminum alloys

[2000Wei].

References

[1942Wei] Weisse, E., Blumenthal, A., Hanemann, H., “Results of an Investigation of Eutectic Zinc

Alloys” (in German), Z. Metallkd., 34(9), 221 (1942) (Equi. Diagram, Experimental, 9)

[1943Bad] Badaeva, T.A., Sal’dau, P.Y., “Physico-Chemical Investigation of Alloys of Aluminium

with Zinc and Lithium” (in Russian), Zhur. Obshchey Khimii, 13(9/10), 643-660 (1943)

(Equi. Diagram, Experimental, 23)

[1963Che] Cherkashin, E.E., Kripyakevich, P.I., Oleksiv, G.I., “Crystal Structures of Ternary

Compounds in Li-Cu-Al and Li-Zn-Al Systems” (in Russian), Sov. Phys., -Crystallogr.,

8(6), 681-685 (1964), translated from Kristallografiya, 8(6), 846-851 (1963) (Crys.


[1986Cas] Cassada, W.A., Shen, Y., Poon, S.J., Shiflet, G.J., “Mg32(Zn,Al)49-Type Icosahedral

Quasicrystals Formed by Solid-State Reaction and Rapid Solidification”, Phys. Rev. B:

Solid State, 34(10), 7413-7416 (1986) (Experimental, 17)

[1987Cas] Cassada, W.A., Shiflet, G.J., Poon, S.J., “Quasicrystalline Grain Boundary Precipitates in

Al Alloys Through Solid-Solid Transformations”, J. Microsc., 146(3), 323-335 (1987)

(Experimental, 26)

[1987Che] Chen, H.S., Phillips, J.C., Villars, P., Kortan, A.R., Inoue, A., “New Quasicrystals of Alloys

Containing s, p and d Elements”, Phys. Rev. B, Cond. Matter, 35B(17), 9326-9329 (1987)


[1987Dub] Dubost, B., Audier, M., Jeanmurt, P., Lang, J.M., Sainfort, P., “Structure of Stable

Intermetallic Compounds of the AlLiCu(Mg) and AlLiZn(Cu) Systems”, J. Phys., Colloq.,

48C3(9), 497-504 (1987) (Crys. Structure, Experimental, 16)

[1989Aud] Audier, M., Janot, C., De Boissieu, M., Dubost, B., “Structural Relationships in

Intermetallic Compounds of the Al-Li-(Cu, Mg, Zn) System”, Philos. Mag. B, 60(4),


[1989Che] Chen, S.-W., Jan, C.- H., Lin, J.-C., Chang, Y. A., “Phase Equilibria of the Al-Li Binary

System”, Metall. Trans., 20A(11), 2247-2258 (1989) (Equi. Diagram, Experimental, #, 59)

[1993Pav] Pavlyuk, V.V., “Synthesis and Crystal Chemistry of Lithium Intermetallic Compounds”,

Doct. Thesis, Univ. L’viv, 1-35 (1993) (Equi. Diagram, Crys. Structure, Experimental,

Review, 49)

[1994Hos] Hosoda, H., Sato, T., Tezuka, H., Mishima, Y., Kamio, A., “Substitution Behavior of

Additional Elements in the L1(2)-Type Al3Li Metastable Phase in Al-Li Alloys” (in

Japanese), J. Jpn. Inst. Met., 58(8), 865-871 (1994) (Crys. Structure, Thermodyn.,

Theory, 26)

[1995Pav] Pavlyuk, V., Bodak, O., MSIT Ternary Evaluation Program, in MSIT Workplace,

Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart;

Document ID: 10.16727.1.20, (1995) (Crys. Structure, Equi. Diagram, Assessment, 9)

[1995Dmy] Dmytriv, G.S., “Isothermal Section of the Phase Diagram of the System Li-Zn-Al at 470 K”

(in Ukrainian), Lvivski Khimichni Chytannya Naukova-Praktychna Konferentsiya, LDU,

108 (1995) (Equi. Diagram, Experimental, 0)

[1996Dmy] Dmytriv, G.S., “Phase Equilibria and Crystal Structure of Compounds in Mg-Li-Si,

Ca-Li-{Si, Ge}, Al-Li-{Si, Ge, Sn}, Zn-Li-{Al, Sn}” (in Ukrainian), Summary of the Thesis

for Candidate Science Degree, Lviv, 1-23 (1996) (Crys. Structure, Equi. Diagram,

Experimental, 10)

127


MSIT®

Al–Li–Zn

[1997Kim] Kim, Y.B., Sommer, F., “Calorimetric Measurement of Liquid Aluminium-Lithium-Zinc

Alloys”, Thermochim. Acta, 291, 27-34 (1997) (Equi. Diagram, Thermodyn.,

Experimental, 16)

[1999Pav] Pavlyuk, V.V., Dmytriv, G.S., Bodak, O.I., Stepien-Damm, J., “New Variant of the

Structure of the Li1+xZn0.5-1.5Al1.5-0.5 Intermetallic Compound”, Materials Structure, 6(2),


[2000Wei] Wei, B.C., Chen, C.Q., Huang, Z., Zhang, Y.G., “Aging Behavior of Li Containing

Al-Zn-Mg-Cu Alloys”, Mat. Sci. Eng. A, 280(1), 161-167 (2000) (Mechan. Prop.,

Experimental, 9)

[2003Gro] Gröbner, J., “Al-Li (Aluminium-Lithium)”, MSIT Binary Evaluation Program, in MSIT



Assessment, 29)


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Li)

< 180.6

cI2

Im3m

W

a = 351.0 pure Li at 25°C [V-C2]

(Zn)

< 419.58

hP2

P63/mmc

Mg

a = 266.50

c = 494.70

at 25°C [Mas2]

(Al)

< 660.45

cF4

Fm3m

Cu

a = 404.96 pure Al at 25°C [Mas2]

Dissolves up to 15 at.% Li

, Li9Al4< 347 - 275

mC26

C2/m

Li9Al4

a = 1915.51

b = 542.88

c = 449.88

= 107.671°

[2003Gro]

´, Li9Al4< 275

? ? [Mas2]

Li3Al2< 520

hR15

R3m

Li3Al2

a = 450.8

c = 1426

[2003Gro]

60 to 61 at.% Li [Mas2]

, LiAl

< 700

cF16

Fd3m

NaTl

a = 637 at 50 at.% Li [2003Gro]

45 to 55 at.% Li [Mas2]

´, LiAl3< 190 - ~120

cP4

Pm3m

Cu3Au


LiZn4

< 245

hP2

P63/mmc

Mg

a = 278.8

c = 439.4

[V-C2], [Mas2]

LiZn4

481 - 65

hP2

P63/mmc

- [Mas2]

128


MSIT®

Al–Li–Zn


Li2Zn5

< 268

hP* a = 437.0

c = 251.5

[V-C2], [Mas2]

Li2Zn5

502 - 168

- - [Mas2]

LiZn2

< 93

- - [Mas2]

Li2Zn3

< 174

cP5 a = 427 [V-C2], [Mas2]

Li2Zn3

520 - 160

- - [Mas2]

LiZn

< 177

cF16

Fd3m

NaTl

a = 623.2 [V-C2], [Mas2]

* 1,

Li1+ Zn0.5-1. 3Al1.5-0.7

cI160

Im3

LiCuSi

a = 1401.7 0.3 to

a = 1390.4 0.3

[1999Pav]

single crystal data

* 2,

LiZn0.6-0.8Al0.4-0.2

cF16

Fd3m

NaTl

a = 625.7 to

a = 621.3

[1996Dmy]

* 3, LiZn3Al

< 490

- - [1943Bad], [1996Dmy]

not found by [1963Che], [1996Dmy]

* 4, Li3ZnAl5P42/mmc

a = 1391

c = 8205

a = 1390

c = 8245

[1987Dub]

sample composition

Li0.33 Zn0.11Al0.56

[1989Aud]


Al Li Zn

L + 1 + (Al) 452 U1 L

1

(Al)

33.2

< 41.5

< 35

< 86

17.5

39.5

35

7

49.3

19.0>

30 >

7 >

L + 1 3 + (Al) 368 U2 L

1

3

(Al)

15.1

< 33.3

< 20

< 88

9.3

33.3

20

4

75.6

33.3 >

60 >

8 >

L (Al) + (Zn) + 3 355 a) E1 L

(Al)

(Zn)

3

13.0

< 88.0

< 3.0

< 16.8

8.2

2.0

2.0

16.8

78.8

10.0 >

95.0 >

66.4 >

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

129


MSIT®

Al–Li–Zn

Note: values in brackets < > are estimated.a) Value given by [1943Bad], 364°C after [1942Wei].

L + 1 580 p1 L

1

32.2

< 41

< 35

32.2

41

35

35.6

18.0 >

30 >

L + 1 3 490 p2 L

1

3

18.6

< 33.3

< 20

18.6

33.3

20

62.8

33.3 >

60 >

L 3 + (Zn) 369 e3 L

3

(Zn)

11

< 16.8

< 2.5

11

16.8

2.5

78

66.4 >

95.0 >


Al Li Zn

10 20 30 40

300

400

500

600

700

Zn Li 50.00Zn 0.00Al 50.00Al, at.%

Te

mp

era

ture

, °C

L

β

τ1

580°C

490°C

369°C

419.58°C

700°C

(Zn)

τ3

Fig. 1: Al-Li-Zn.

The pseudobinary

system Zn - LiAl

130


MSIT®

Al–Li–Zn

Fig

. 2:

Al-

Li-

Zn

. R

eact

ion s

chem

e

Al-

LiA

lA

l-Z

nL

iAl-

Zn

Al-

LiA

l-Z

n

l (

Al)

+ (

Zn)

38

1e 2

l (

Al)

+ β

60

0e 1

L +

βτ 1

+ (

Al)

45

2U

1

L +

βτ 1

58

0p

1

(Al)

´´

(A

l)´

+ (

Zn)

27

7e 4

L +

τ1

τ 3

49

0p

2

L (

Zn

) +

τ3

36

9e 3

L +

τ1

τ 3 +

(A

l)3

68

U2

L (

Al)

+ (

Zn)

+ τ

33

55

E1

(Al)

´´

(A

l)´+

(Zn

)+τ 3

27

5E

2

L+

(Al)

+τ 1

(Al)

+(Z

n)+

τ 3

(Al)

´´+

(Zn)+

τ 3

L+

τ 1+

τ 3

β +

τ1 +

(A

l)

L+

(Al)

+τ 3

τ 1+

τ 3+

(Al)

131


MSIT®

Al–Li–Zn

20

40

60

80

20 40 60 80

20

40

60

80

Li Zn


Axes: at.%

β

(Al)´

(Zn)

τ3

τ1

(Al)´´+τ

3 +(Zn)

(Al)´+τ1+τ

3

(Al)´´

β+(Al)´

(Al)

´+τ 1

20

40

60

80

20 40 60 80

20

40

60

80

Li Zn


Axes: at.%

LiAl

(Al)

β

τ1

τ3

(Zn)

e1

p1

U1

p2

e3

U2 E

1

e2

600

550

500

470

450

430 420

380400

650

700 600

Fig. 4: Al-Li-Zn.

Partial isothermal

section at 350°C

Fig. 3: Al-Li-Zn.

Partial liquidus

surface

132


MSIT®

Al–Li–Zn

20

40

60

80

20 40 60 80

20

40

60

80

Li Zn


Axes: at.%(Al)

τ1

τ3

τ2

β

Li3Al

2

δ´

βLi2Zn

3 βLi2Zn

5 βLiZn4

(Zn)αLiZn4

αLi2Zn

5

-15

-10

-5

0

5

0 20 40 60 80 100

Al

Zn

Li

100-

0.00

x

xLi, at.%

∆ mix

H,

kJ

mo

l-1·

Li

5

=30 (610°C)x

x=75 (518°C)

x=50 (554°C)

Fig. 5: Al-Li-Zn.

The isothermal

section at 193°C

Fig. 6: Al-Li-Zn.

Experimental

enthalpy of mixing

for ternary

undercooled liquid

alloys

133


MSIT®

Al–Li–Zn

5

0

-5

-10

0 20 40 60 80 100

∆ mix

H,

kJ

mo

l-1·

Al

Zn

Li

Li, at.%75.0000.0025.00

Li

20

40

60

80

20 40 60 80

20

40

60

80

Li Zn


Axes: at.%

-3.5

9

-1.5

-1.0

τ2

τ3

τ1

experimental

calculated

Fig. 7: Al-Li-Zn.

Enthalpy of mixing of

(Al0.75Li0.25)1-xZnx

ternary liquid and

undercooled liquid

alloys at 682°C

Fig. 8: Al-Li-Zn.

Difference (in kJ mol-1)

between the

experimental and the

calculated enthalpy of

mixing of Al-Li-Zn

ternary liquid and

undercooled liquid

alloys at 682°C using

the association model

134


MSIT®

Al–Li–Zr

Aluminium – Lithium – Zirconium

Oksana Bodak

Literature Data

Most investigations on the Al-Li-Zr system concern the metastable phases ´, LiAl3 and ´, ZrAl3.

[1985Mak] studied the recrystallization behavior of an 3Li-0.12Zr-Al (mass%) alloy in comparison to that

of the binary alloys 2.5Li-Al (mass%), 3Li-Al (mass%) and 0.13Zr-Al (mass%). [1984Gay] prepared an

2.34Zr-Al (mass%) alloy by rapid solidification and observed after aging at 190°C a discontinuous

precipitation behavior: ZrAl3 was precipitated as aligned rods or as discrete spheres. The ZrAl3/(Al)

interface served as a nucleation site for ´, LiAl3. The resulting “composite” precipitate contained a core of

ZrAl3 and an envelope of ´, LiAl3.

[1986Gay1] found in the same alloy a ternary phase between LiAl3 and ZrAl3, expressed by the formula

(LixZr1-x)Al3 with 0.45 < x < 0.8, see Table 1. Physical and thermodynamic properties of this phase were

investigated by [1986Gay2]. The metastable phase (LixZr1-x)Al3 is also given in Table 1 because of its

technical importance [1986Gay1, 1986Gay2, 1986Sak].

In an alloy 3Zr-Al (mass%), [1986Sak] observed the ´ and ´ phases as distinct phases by a time-of-flight

atom-probe field-ion microscopy (ToF atom-probe FIM).

The nucleation of ´ on ´ as a substrate was studied theoretically by [1987Tos]. The precipitation of ´ in

several ternary and quaternary alloys was reviewed by [1987Flo]. By adiabatic scanning calorimetry

[1988Eun] examined precipitation and dissolution reactions.

Partial vapor pressures of Li over binary and ternary aluminium melts at 927°C were calculated using an

interaction parameter for Zr as a third element [1986Lee]. [1989Sau] calculated phase diagrams for stable

as well as for metastable phase equilibria in the Al-Li-Zr system.

The effects of mechanical alloying, a low temperature isothermal processing method, and the effect of

ternary addition of lithium on the phase stability of the ZrAl3 phase with metastable cubic L12 structure

were studied in [1991Des]. At 750°C it was found that adding lithium increases the stability of the L12

phase. The literature until 1989 was compiled and critically reviewed by [1995Pav].

The results of an investigation of the isothermal section of Al-Li-Zr at 197°C and data of the crystal

structure of the compounds are reported in [2002Zat]. The alloys were prepared by arc melting in purified

argon atmosphere under a pressure of ~1.01 105 Pa from a mixture of the pure metals (Zr of 99.98% mass

purity, Li of 99.0 mass% purity, and Al of 99.99 mass% purity). The alloy compositions were checked by

weight comparison of the initial mixtures and the alloys. The alloys were annealed at 197°C for 400 h in

tantalum containers in evacuated quartz ampoules and quenched in cold water. The X-ray powder method

was used for the phase analysis and structural investigation.

Binary Systems

The Al-Li system reported by [2003Gro] is accepted. The Al-Zr phase diagram presented by [2003Sch]

shows more likely features than the those given in the diagram by [Mas2], in which all the liquidus lines are

drawn tentatively. The Li-Zr system is accepted from [Mas2]. The extremely small solubility of Zr in liquid

Li was calculated by [1989Sau]. In the range of 7.5 at.% Li, the stable solid phases are (Al) and , LiAl.

However, a metastable LiAl3 occurs and creates order hardening in the alloys. The metastable solvus

(Al)/LiAl3 has been experimentally determined by [1998Nob]; Zr additions up to 0.05 at.% do not affect

the position of the metastable boundary.

Solid Phases

The ternary compounds ZrLi2Al has a narrow range of homogeneity and Zr5-xLix+yAl3 (x = 0.2 - 1.0,

y = 0 - 1) exhibits a relatively wide homogeneity range, see Table 1.

135


MSIT®

Al–Li–Zr

Liquidus Surface

The calculated liquidus surface in the composition range 0 to 10 at.% Zr and 0 to 50 at.% Li is shown in

Fig. 1 [1989Sau]. The ternary invariant reactions at 595.4°C with a composition of the melt of 24.65 at.%

Li and 5.1 10-5 at.% Zr cannot be reproduced in Fig. 1 because of the low Zr content in the melt.

Isothermal Sections

Partial isothermal sections at 500, 300 and 100°C were given for the composition range 0 to 25 at.% Zr and

0 to 50 at.% Li [1989Sau]. Since they are rather similar, only the section at 300°C is presented in Fig. 2.

Beyond this the ordering of the ´, (LixZr1-x)Al3 phase was thermodynamically described [1989Sau].

The isothermal section of the Al-Li-Zr phase diagram at 197°C is shown in Fig. 3. The results show good

agreement between the experimental data of [2002Zat] and the calculated part of isothermal section

[1989Sau].

The formation process of Zr5-xLix+yAl3 ternary intermetallic is realized by a partial substitution of Li atoms

by Zr in the 4(d) position and insertion of lithium atoms in holes at the 2(b): 000 position. The change of

the lattice parameters in the Zr5-xLix+yAl3 homogeneity range is presented in Fig. 4 after [2002Zat].

The Zr5Al4 binary compound (Ti5Ga4 structure type) has not been found in the Al-Li-Zr system at 197°C.

It is stable in the temperature range from 990 to 1530°C. The Zr5-xLix+yAl3 ternary compound, apparently,

is a remainder of the high temperature substitution-limited solid solution of the Zr5Al4 binary compound or

of the substitution- and insertion-limited solid solution of Zr5Al3 (Mn5Si3 structure type). The characteristic

feature of the Al-Li-Zr ternary system is the binary immiscibility region of Li-Zr extending up to ~10 at.%

of the third component. Limited solid solutions of the binary compounds of the Al-Zr system were observed

in the Al-Li-Zr system. Largest solubility of the third component is found in ZrAl3 (5 at.%), ZrAl2 (10 at.%)

and Zr2Al3 (15 at.%).

For the ZrLixAl3-x solid solution the change of the lattice parameters vs Li-concentration is presented in

Fig. 5 after [2002Zat].


In cast aluminium alloys Zr is typically added to achieve grain refinements and to inhibit the

recrystallization of wrought structures. This behavior is associated with the formation of coherent ZrAl3particles of metastable cubic form [1987Flo] which is stabilized by Li [1987Vec]. In addition, Zr is used to

impart superplasticity, or improve strength and toughness of rapidly solidified Al-Li alloys. In the ternary

system, Zr precipitates in a supersaturated solid solution via a normal nucleation and growth mechanism as

coherent spherical or filamentary particles, depending on the heat treatment as (Li,Zr)Al3, metastable,

Cu3Au type phase [1989Gay, 1994Hos].

References

[1984Gay] Gayle, F.W., Vander Sande, J.B., “’Composite’ Precipitates in an Al-Li-Zr Alloy”, Scr.

Metall., 18, 473-478 (1984) (Experimental, 13)

[1985Mak] Makin, P. L., Stobbs, W.M., “Comparison of the Recrystallization Behaviour of an Al-Li-Zr

Alloy with Related Binary Systems”, The Institute of Metals, London, Accession Number:

86(8), 72-312; 392-401 (1986) (Experimental, 11)

[1986Gay1] Gayle, F.W., Vander Sande, J.B., “Al3Li Precipitate Modification in an Al-Li-Zr Alloy”,

ASTM, Proc. Pennsylvania, 1984, 137-152 (Publ. 1986) (Crys. Structure, Experimental, 16)

[1986Gay2] Gayle, F.W., Vander Sande, J.B., “Al3(Li, Zr), or ´ Phase in Al-Li-Zr System”, The

Institute of Metals, London, accession Number, 86(8), 72-312, 376-384 (1986) (Crys.


[1986Lee] Lee, J.J., Sommer, F., “Thermodynamic Properties of Lithium in Liquid Aluminum Alloys”

(in Korean), Taehan Kumsok Hakhoechi, 24(10), 1185-1189 (1986) (Thermodyn.,

Experimental, 19)

136


MSIT®

Al–Li–Zr

[1986Sak] Sakurai, T., Kobayashi, A., Hasegawa, J., Sakai, A., Pickering, H.W., “Atomistic Study of

Metastable Phases in Al - 3 wt.% Zr Alloy”, Scr. Metall., 20, 1131-1136 (1986) (Crys.


[1987Flo] Flower, H.M., Gregson, P.J., “Solid State Phase Transformations in Aluminium Alloys

Containing Lithium”, Mater. Sci. Technol., 3(2), 81-90 (1987) (Review, 116)

[1987Tos] Tosten, M. H., Galbraith, J. M., Howell, P. R., “Nucleation of ´ (Al3Zr) in Al-Li-Zr and

Al-Li-Cu-Zr Alloys”, J. Mater. Sci. Lett., 6(2), 51-53 (1987) (Experimental, 10)

[1987Vec] Vecchio, K.S., Williams, D.B., “Convergent Beam Electron Diffraction Study of Al3Zr in

Al-Zr and Al-Li-Zr Alloys”, Acta Metall., 35(12), 2959-2970 (1987) (Crys. Structure,

Experimental, 19)

[1988Eun] Eun, I.-S., Woo, K.-D., Cho, H.K., “The Formation of Precursor Phase During Precipitation

in Al-Li-Zr Alloy” (in Korean), J. Korean Inst. Met., 26(11), 1007-1012 (1988)


[1989Che] Chen, S.W., Tan, C.-H., Lin, T.-C., Chang, Y.A., “Phase Equilibria of the Al-Li Binary

System”, Metall. Trans. A, 20A(11), 2247-2258 (1989) (Equi. Diagram, Experimental,

Thermodyn., #, 59)

[1989Gay] Gayle, F.W., Vandersande, B., “Phase Transformations in the Al-Li-Zr System”, Acta

Metall., 37(4), 1033-1046 (1989) (Crys. Structure, Experimental, Thermodyn., 28)

[1989Sau] Saunders, N., “Calculated Stable and Metastable Phase Equilibria in Al-Li-Zr Alloys”,

Z. Metallkd., 80(12), 894-903 (1989) (Equi. Diagram, Thermodyn., Theory, #, *, 78)

[1991Des] Desch, P.B., Schwarz, R.B., Nash, P., “Formation of Metastable L12 Phases in Al3Zr and

Al-12.5% X-25 % Zr (X = Li, Cr, Fe, Ni, Cu)”, J. Less-Common Met., 168, 69-80 (1991)


[1994Hos] Hosoda, H., Sato, T., Tezuka, H., Mishima, Y., Kamio, A., “Substitution Behavior of

Additional Elements in the L12-Type Al3Li Metastable Phase in Al-Li Alloys” (in

Japanese), J. Jpn. Inst. Met., 58(8), 865-871 (1994) (Crys. Structure, Theory,

Thermodyn., 26)

[1995Pav] Pavlyuk, V., Bodak, O., “Aluminium-Lithium-Zirconium”, MSIT Ternary Evaluation




[1998Nob] Noble, B., Bray, S.E., “On the (A1)/ ´(Al3Li) Metastable Solvus in Aluminium-Lithium

Alloys”, Acta Mater., 46(17), 6163-6171 (1998) (Calculation, Experimental, Phys. Prop.,

Thermodyn., 41)

[2002Zat] Zatorska, G.M., Pavlyuk, V.V., Davydov, V.M., “Phase Equilibria and Crystal Structure of

Compounds in the Zr-Li-Al System at 470 K”, J. Alloys Compd., 333, 138-142 (2002)

(Equi. Diagram, Crys. Structure, Experimental, #, 11)




Assessment, 21)


MSIT Workplace, Effenberg, G. (Ed.), Materials Science International Services, GmbH,

Stuttgart; submitted for publication (2003) (Crys. Structure, Equi. Diagram,

Assessment, 103)

137


MSIT®

Al–Li–Zr


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Li)

< 180.6

cI2

Im3m

W


[V-C2]

(Al)

< 660.45

cF4

Fm3m

Cu

a = 404.96 pure Al at 25°C

[Mas2]

dissolves up to 15 at.% Li

( Zr)

1855 - 863

cI2

Im3m

W

a = 360.90 [Mas2]

( Zr)

< 863

hP2

P63/mmc

Mg

a = 323.16

c = 514.75

[2003Sch]

dissolves up to 8.3 at.% Al at 910°C

Li9Al4347 - 275

mC26

C2/m

Li9Al4

a = 1915.51

b = 542.88

c = 449.88

= 107.671°

[2003Gro]

Li9Al4 ( ´)

< 275

? ? [Mas2]

Li3Al2 ( )

< 520

hR15

R3m

Li3Al2

a = 450.8

c = 1426

[2003Gro]

60 to 61 at.% Li

[Mas2]

LiAl ( )

< 700

cF16

Fd3m

NaTl

a = 637 at 50 at.% Li [2003Gro]

45 to 55 at.% Li [Mas2]

LiAl3 ( ´)

400

cP4

Pm3m

Cu3Au

a = 403.8 Metastable [1989Che, 2003Gro]

Zr3Al

< 1019

cP4

Pm3m

Cu3Au

a = 439.17 [V-C2, Mas2]

Zr2Al

< 1215

hP6

P63/mmc

Ni2In

a = 489.39

c = 592.83

[2003Sch]

Zr5Al3 (r)

1000

hP16

P63/mcm

Mn5Si3

a = 818.4

c = 570.2

[2003Sch]

Zr5Al3 (h)

1400 - 1000

tI32

I4/mcm

W5Si3

a = 1104.4

c = 539.1

[2003Sch]

Zr3Al2< 1480

tP20

P42/mnm

Zr3Al2

a = 763.0

c = 699.8

[2003Sch]

138


MSIT®

Al–Li–Zr

Zr4Al3< 1030

hP7

P6/mmm

Zr4Al3

a = 543.0

c = 539.0

[2003Sch]

Zr5Al41550 - 1000

hP18

P63/mcm

Ti5Ga4

a = 844.8

c = 580.5

[2003Sch]

ZrAl

< 1275

oC8

Cmcm

CrB

a = 335.3

b = 1086.6

c = 426.6

[2003Sch]

Zr2Al3< 1590

oF40

Fdd2

Zr2Al3

a = 960.1

b = 1390.6

c = 557.4

[2003Sch]

ZrAl2< 1660

hP12

P63/mmc

MgZn2

a = 528.24

c = 874.82

[2003Sch]

ZrLixAl3-x

ZrAl3 < 1580

tI16

I4/mmm

ZrAl3

a = 400.9

c = 1728.2

a = 401.4

c = 1727.7

x = 0.2 (Li0.2Al2.8Zr)

[2002Zat]

x = 0 (ZrAl3)

ZrAl3 cP4

Pm3m

Cu3Au

a = 408 Metastable, stabilized by Li [1987Vec,

1989Gay]

* 1, Li2ZrAl cF12

F43m

CuHg2Ti

a = 663.3 [2002Zat]

* 2, Lix+yZr5-xAl3 hP18

P63/mcm

Ti5Ga4

a = 813.36

c = 570.29

a = 817.57

c = 569.09

Li0.2Zr4.8Al3(x = 0.2, y = 0)

LiZr4Al3(x = 1, y = 0) [2002Zat]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

139


MSIT®

Al–Li–Zr

Li, at.%

Zr,

at.

%

0

0

10 20 30 40 50

2

10

4

6

8

Al

15001400

1300

12001100

900

800

ZrAl3

10

20

30

40

10 20 30 40

60

70

80

90

Li 50.00Zr 0.00Al 50.00

Li 0.00Zr 50.00Al 50.00


Axes: at.%(Al)

LiAl

ZrAl3

(Al)+LiAl+ZrAl3

Fig. 2: Al-Li-Zr.

Calculated partial

isothermal section at

300°C

Fig. 1: Al-Li-Zr.

Calculated partial

liquidus surface

140


MSIT®

Al–Li–Zr

20

40

60

80

20 40 60 80

20

40

60

80

Li Zr


Axes: at.%(Al)

τ1

τ2

L

LiAl

Li3Al

2

Li9Al

4

ZrAl3

ZrAl2

Zr2Al

3

ZrAl

Zr4Al

3

Zr3Al

2

Zr2Al

Zr3Al

(Zr)

Li, at.%

La

ttic

epa

ram

ete

r,p

m

0

813

10 20 30

569.8

326.0

V,p

m1

03

-6×

815

817

818

816

814

569.0

569.4

570.2

327.0

328.0

329.0

a

c

V

Fig. 3: Al-Li-Zr.


197°C

Fig. 4: Al-Li-Zr.

Change of lattice

parameters for

Zr5-xLix+yAl3

141


MSIT®

Al–Li–Zr

Li, at.%

La

ttic

epa

ram

ete

r,p

m

0

400.8

5 10 15

277.68

V,p

m1

03

-6×

1727.6

a

c

V

401.0

401.2

401.4

1727.8

1728.0

1728.2

277.72

277.76

277.80Fig. 5: Al-Li-Zr.

Change of lattice

parameters for the

ZrLixAl3-x solid

solution

142


MSIT®

Al–Mg–Mn

Aluminium – Magnesium – Manganese

Qingsheng Ran, updated by Joachim Gröbner

Literature Data

The system has been mainly investigated on the Al-Mg side. [1952Han, 1980Bra] reviewed information

about the system, but the scopes covered by each are quite limited.

The Al rich corner: with about 150 alloys [1938Lee] studied the system in the range 0-35.5 mass% Mg and

0-12 mass% Mn, by means of metallography and thermal analysis and in some cases supplemented by

annealing experiments and X-ray crystallography. Alloys were prepared from aluminium of 99.991%

purity, magnesium of 99.996% purity, and an aluminium-manganese hardener containing 13-25 mass% Mn

prepared directly from aluminium and dehydrated MnCl2. The results were drawn mainly from

microscopical observation of cast alloys. A short version of this work was given by [1938Han]. By X-ray

diffraction (Debye-Scherrer and rotating crystal methods) of a 16.28 mass% Mg and 4.26 mass% Mn alloy,

[1938Hof] reported the solid phases formed by a metastable eutectic in the Al corner to be Al solid solution,

Mg2Al3 and MnAl4. With about 100 alloys, prepared from Mg, 99.99% pure Al and a high purity Al-Mn

alloy, [1940Fah] investigated the joint solubilities of Mg and Mn in (Al) at 500 to 650°C by electrical

resistance measurements. To obtain more detailed information about the liquidus surface of the Al corner,

[1943But] studied 20 alloys with up to 5 mass% Mg and 2 mass% Mn. Aluminium of super-purity grade

and aluminium-manganese master alloys in the same purity degree and magnesium of 99.95% purity were

used for preparing the alloys for the determination of cooling curves. Nine alloys were studied by [1943Lit]

for determining the effect of Mg on the solubility of Mn at 500°C by microstructure observation. The

materials and experimental procedure used by [1943Lit] were the same as those of [1943But]. [1943Lit]

stated in addition that no new phases appear in the Al-5Mg-2Mn (mass%) range at 400°C. [1943Mon] drew

equilibrium diagrams for the Al corner from data by [1938Lee] and own values, but did not give any details

on the results and procedure of their own experiments. [1945But] continued the work of the constitution of

the Al corner and determined the solidus isotherms by observation of incipient melting and microstructure.

An isothermal section at 630°C was also presented. Considering the limited composition range or

nonequilibrium condition, [1948Wak] carried out microstructural observation of 45 samples for

determining the phase relationships in the region of aluminium with up to 40 mass% Mg and 25 mass% Mn

at 400°C. High purity aluminium and magnesium metals and aluminium-manganese master alloys were

melted, cast and annealed at 400°C. In some cases, slowly-cooled alloys were also examined. X-ray

diffraction was used for identifying the phases. Using 99.99% Al, 99.9% Mg and 99.9% Mn [1973Ohn1]

prepared 40 alloys. After melting, the samples were cast and then annealed at 450°C for 20 days and at

400°C for 40 days, respectively. The quenched samples were investigated by metallography and X-ray

diffraction analysis. Isothermal sections of the aluminium side with up to 15 mass% Mg and 6 mass% Mn

at 450 and 400°C were established. The structure of a ternary phase was determined. In a work primarily on

the quaternary system Al-Cr-Mg-Mn [1973Ohn2] 6 alloys were examined for studying the constitution of

the Al corner of the Al-Mg-Mn system at 550°C. Most results of the above investigations are consistent with

each other. However, the isothermal sections at 435 and 400°C, established by [1938Lee, 1948Wak,

1973Ohn1], respectively, are inconsistent. The reason might be that the equilibrium state was not achieved

by [1938Lee]. The liquidus surface from [1938Lee] is accepted, but more experiment in this region is

necessary.

The Mg rich corner: [1938Ima] investigated the Mg-35Al-6Mn (mass%) region with 17 samples. The

starting materials were 99.8% pure Mg and Al, metallic Mn, an Al-19.8%Mn master alloy and MnCl2 for

preparing ternary alloys. Thermal analysis and microscopic examination were used. [1944Bee] determined

several solubility curves of Mn and Al in Mg at different temperatures. [1948Age] studied the Mg corner

with up to 40 mass% Al and 10 mass% Mn by thermal analysis, metallography and X-ray diffraction. A

liquidus surface and some invariant equilibria are presented, but these do not agree with [1938Ima].

[1957Mir] prepared samples from metallic Mg (~99.9%), Al (~99.99%) and electrolytic Mn, from which

143


MSIT®

Al–Mg–Mn

first Mg-Mn and Al-Mn master alloys were made. The liquidus surface of the Al-Mg side with up to ~65

mass% Al was constructed from microstructural analysis on about 20 cast alloys. 75 alloys were examined

by metallography and microhardness for determining phase equilibria in the Mg corner at 200 to 400°C.

The samples were heat treated in evacuated silica ampules for 9 to 55 days and water quenched. Results

were given in two partial isothermal sections. [1986Obe, 1988Sim1, 1988Sim2] reported phases existing in

equilibrium with liquid at temperatures between 660 and 760°C in Mg alloys with up to 10 mass% Al and

1.5 mass% Mn. MnCl2 was added to the melts at 780°C to saturate the alloys with Mn. After thorough

stirring, the melts were held for 1 to 2 h at 750, 710 and 670°C, respectively. The samples were made either

by ordinary casting or rapid quench against a spinning, water-cooled wheel and examined by

microstructural, X-ray diffraction and microprobe analysis. Phases in equilibrium with liquid and the single

phase region of melt for the temperatures 750, 710 and 670°C were determined. [1992Ars] prepared

samples in the Al rich corner with constant 10 mass% Mg by rapid quenching in water. They report a

calculated metastable vertical section which is not in equilibria with the ternary phase T.

Binary Systems

The binary system Al-Mg was updated by [2003Luk]. This version is accepted. The Al-Mn system is

accepted from [2003Pis] and Mg-Mn is taken from [Mas].

Solid Phases

[1948Wak] revealed a ternary phase T by metallographic observation and X-ray diffraction. The

composition of this phase is near MnMg2Al10. The structure was determined by [1973Ohn1, 1994Fun] who

suggested the composition of the phase to be Mn2Mg3Al18. A Mn rich phase X was proposed by [1948Age]

without giving details on structure or composition. It is quite probably the same phase as X in [1957Mir]

who concluded that X should be an Al-Mn binary phase. The ternary phase T and the binary solid phases

present in the compiled phase diagrams are listed in Table 1.


Some four-phase equilibria were reported. The reactions listed in Table 2 are based on [1938Lee] (the first

three) and [1948Age, 1938Ima] (the last two). It should be noted that all these reactions are not certain.

According to [1948Wak, 1973Ohn1] the reactions given by [1938Lee] might be metastable. The region of

the primary solidification of the ternary compound T reported by [1957Mir] makes the reactions according

to [1938Lee] also doubtful. These reactions therefore need further investigation.

Liquidus Surface

A liquidus surface projection on the Al-Mg rich side is constructed using data from different investigations,

Fig. 1. Because of the different opinions on some reactions (see section Invariant Equilibria) and the

incomplete determination of other reactions, this liquidus projection has to be considered as tentative.

Isothermal Sections

Isothermal sections of the Al corner at 630°C [1945But] and 400°C [1948Wak, 1973Ohn1] are given in

Figs. 2 and 3, respectively. An isothermal section at a temperature just after the end of crystallization was

proposed by [1938Lee], but is contradictory to [1948Wak] and [1973Ohn1], who studied the topic more

carefully. The joint solubility of Mg and Mn in solid (Al) is given in Fig. 4; the data are mainly from

[1940Fah]. Isothermal sections of the Mg corner at 400°C and 200°C are plotted in Figs. 5 and 6,

respectively.

144


MSIT®

Al–Mg–Mn

References

[1938Han] Hanemann, H., Schrader, A., “On Some Ternary Systems of Aluminium, I Aluminium -

Iron - Magnesium, Aluminium-Magnesium-Manganese, Aluminium-Manganese-Silicon”

(in German), Z. Metallkd., 30, 383-386 (1938) (Equi. Diagram, Experimental, #, 11)

[1938Hof] Hofmann, W., “X-Ray Methods on Investigation of Aluminium Alloys” (in German),

Aluminium, 865-872 (1938) (Crys. Structure, Experimental, 19)

[1938Ima] Imaki, A., “On the Equilibrium Diagram of Mg-Al-Mn Alloy System” (in Japanese), Trans.

Min. Met. Alumi. Assoc., 9, 665-668 (1938) (Equi. Diagram, Experimental, 1)

[1938Lee] Leemann, W.G., “The Ternary System Aluminium-Magnesium-Manganese” (in German),

Aluminium Arch., 9, 6-17 (1938) (Equi. Diagram, Experimental, 7)

[1940Fah] Fahrenhorst, E., Hofman, W., “The Solubility of Manganese in Aluminium with up to 2 %

Mg” (in German), Metallwirtschaft, 19, 891-893 (1940) (Equi. Diagram, Experimental, 3)

[1943But] Butchers, E., Raynor, G.V., Hume-Rothery, W., “The Constitution of

Magnesium-Manganese-Zinc-Aluminium Alloys in the Range 0-5 % Magnesium, 0-2 %

Manganese, 0-8 % Zinc, I-The Liquidus”, J. Inst. Met., 69, 209-228 (1943) (Equi. Diagram,

Experimental, 9)

[1943Lit] Little, A.T., Raynor, G.V., Hume-Rothery, W., “The Constitution of Magnesium -

Manganese - Zinc - Aluminium Alloys in the Range 0-5 % Magnesium, 0-2 % Manganese

and 0-8 % Zinc, III-The 500C and 400C Isothermals”, J. Inst. Met., 69, 423-440 (1943)


[1943Mon] Mondolfo, L.F., “Metallography of Aluminium Alloys”, John Wiley and Sons, Inc., New

York, 100-101 (1943) (Equi. Diagram, Review, 1)

[1944Bee] Beerwald, A., “On the Solubility of Iron and Manganese in Magnesium and in

Magnesium-Aluminium Alloys” (in German), Metallwirtschaft, 23, 404-407 (1944) (Equi.

Diagram, Experimental, 10)

[1945But] Butchers, E., Hume-Rothery, W., “On the Constitution of Aluminium - Magnesium -

Manganese - Zinc Alloys: The Solidus”, J. Inst. Met., 71, 291-311 (1945) (Equi. Diagram,

Experimental, #, 8)

[1948Age] Ageev, N.V., Kornilov, I.I., Khlapova, A.N., “Magnesium-Rich Alloy of the System

Magnesium-Aluminium-Manganese” (in Russian), Izv. Inst. Fiz.-Khim. Anal., Inst.

Obshcheii Neorg. Khim., Akad. Nauk SSSR, 14, 130-143 (1948) (Equi. Diagram,


[1948Wak] Wakeman, D.W., Raynor, G.V., “The Constitution of Aluminium-Manganese-Magnesium

and Aluminium-Manganese-Silver Alloys, with Special Reference to Ternary Compound

Formation”, J. Inst. Met., 75, 131-150 (1948) (Equi. Diagram, Experimental, *, 27)

[1952Han] Hanemann, H., Schrader, A., “Ternary Alloys of Aluminium” (in German), Verlag

Stahleisen m.b.H., Dusseldorf, 116-120 (1952) (Equi. Diagram, Review, 3)

[1957Mir] Mirgalovskaya, M.S., Matkova, L.N., Komova, E.M., “The System Mg-Al-Mn” (in

Russian), Trudy Inst. Met. Im. A.A. Baikova, Akad. Nauk, 2, 139-148 (1957) (Equi.


[1973Ohn1] Ohnishi, T., Nakatani, Y., Shimizu, K., “Phase Diagrams and Ternary Compounds of the

Al-Mg-Cr and the Al-Mg-Mn Systems in Al-Rich Side” (in Japanese), Light Metals Tokyo,

23, 202-209 (1973) (Crys. Structure, Equi. Diagram, Experimental, *, 16)

[1973Ohn2] Ohnishi, T., Nakatani, Y., Shimizu, K., “Phase Diagram in the Al-Rich Side of the

Al-Mg-Mn-Cr Quarternary System” (in Japanese), Light Metals Tokyo, 23, 437-443 (1973)


[1980Bra] Brandes, E.A., Flint, R.F., “Manganese Phase Diagrams”, Manganese Center, 17 Ave.

Hoche, 75008 Paris, France, 82 (1980) (Equi. Diagram, Review, 2)

[1986Obe] Oberlaender, B.C., Simensen, C.J., Svalestuen, J., Thorvaldsen, A., “Phase Diagram of

Liquid Magnesium - Aluminium - Manganese Alloys”, Magnesium Technology, Pros.

Conf., London, 133-137 (1986) (Experimental, 3)

145


MSIT®

Al–Mg–Mn

[1988Sim1] Simensen, C.J., Oberländer, B.C., Svalestuen, J., Thorvaldsen, A., “Determination of the

Equilibrium Phases in Molten Mg - 4 wt.% Al-Mn Alloys”, Z. Metallkd., 79, 537-540


[1988Sim2] Simensen, C.J., Oberländer, B.C., Svalestuen, J., Thorvaldsen, A., “The Phase Diagram for

Magnesium - Aluminium - Manganese above 650°C”, Z. Metallkd., 79, 696-699 (1988)

(Experimental, 10)

[1992Ars] Arsenov, A.A., Goutan, D., Zolotarevskii, V.S., Kuznetsov, G.M., Lugin, D.V., “Study of

Decomposition of the (Al)-Solid Solution Heating for Quenching of Cast Alloys Al-10%

Mg and Al-6% Zn-15% Mg-1% Cu Containing Manganese” (in Russian), Metally, 6, 80-83


[1994Fun] Fun, H.-K., Lin, H.-C., Lee, T.-J., Yipp, B.-C., “T-Phase Al18Mg3Mn2”, Acta Crystallogr.,

C50, 661-663 (1994) (Crys. Structure, 5)




Assessment, 49)

[2003Pis] Pisch, A., “Al-Mn (Aluminium-Manganese)”, MSIT Binary Evaluation Program, in MSIT




Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Al)

< 660.5

cF4

Fm3m

Cu

a = 404.88 [V-C], pure 23°C

(Mg)

< 650

hP2

P63/mmc

Mg

a = 320.89

c = 521.01

[V-C], pure

( Mn)

< 1079

cP20

P4132

Mn

a = 631.5 pure Mn, [V-C]

( Mn)

< 710

cI58

I43m

Mn

a = 891.39 pure Mn, [V-C]

, Mg2Al3 452

cF1168

Fd3m

Mg2Al3

a = 2816 to 2824 60-62 at.% Al [2003Luk]


[2003Luk]

, Mg17Al12

< 458

cI58

I43m

Mn

a = 1054.38 at 41.4 at.% Al [V-C2]

39.5 to 51.5 at.% Al [2003Luk]

, Mg23Al30

410 - 250

hR159

R3

Mn44Si9

a = 1282.54

c = 2174.78

54.5 to 56.5 at.% Al [2003Luk]

146


MSIT®

Al–Mg–Mn


a)X is an Al-Mn binary compound [1957Mir]

MnAl6< 705

oC28

Cmcm

MnAl6

a = 754.5 0.2

b = 649.0 0.3

c = 868.1 0.2

[2003Pis]

, MnAl4< 923

hexagonal - [Mas]

Mn4Al11(r)

916

aP30

P1

Mn4Al11

a = 509.5 0.4

b = 887.9 0.8

c = 505.1 0.4

= 89.35 0.04°

= 100.47 0.05°

= 105.08 0.06°

[2003Pis]

* T, Mn2Mg3Al18 cF184

Fd3m

Cr2Mg3Al18

a = 1452.9

a = 1451.7

[1973Ohn1]

[1994Fun]


Al Mg Mn

L + Mn4Al11(r) + - U1 L

Mn4Al11(r)

67.7

73.3

81.5

62.0

30.6

0

0

37.5

1.7

26.7

19.5

0.5

L + MnAl6 + - U2 L

MnAl6

69.3

81.5

85.7

61.5

29.5

0

0

38.0

1.2

19.5

14.3

0.5

L (Al) + + MnAl6 437 E1 L

Al

MnAl6

70.7

84.5

61.0

85.7

28.3

15.0

38.5

0

1.0

0.5

0.5

14.3

L + (Mg) + ( Mn)(?) ~437 U3 L 30 69.5 0.5

L + ( Mn)(?) + X a) ~430 E2 L 34 64.6 1.4

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

147


MSIT®

Al–Mg–Mn

20

40

60

80

20 40 60 80

20

40

60

80

Mg Mn


Axes: at.%(Al)

MnAl6

µ

U2

U1

E1

P

X

E2

βMn

αMn

β

γ

U3

Mg

Fig. 1: Al-Mg-Mn.

Liquidus surface on

the Al-Mg side

10

10

90

Mg 20.00Mn 0.00Al 80.00

Mg 0.00Mn 20.00Al 80.00


Axes: at.%

(Al)

(Al)+L

L

(Al)+MnAl6

(Al)+MnAl6+L

MnAl6+L

MnAl6

Fig. 2: Al-Mg-Mn.

Isothermal section of

the Al corner at 630°C

[1945But]

148


MSIT®

Al–Mg–Mn

10

20

30

10 20 30

70

80

90

Mg 40.00Mn 0.00Al 60.00

Mg 0.00Mn 40.00Al 60.00


Axes: at.%

(Al)

(Al)+MnAl6

(Al)+MnAl6+T

(Al)+T

(Al)+T+β

T+β+ε

T

MnAl6

β

Mg 5.00Mn 0.00Al 95.00

Mg 0.00Mn 5.00Al 95.00


Axes: at.%

400°C

500°C

550°C

600°C

Fig. 3: Al-Mg-Mn.


the Al corner at 400°C

Fig. 4: Al-Mg-Mn.

Joint solubility of Mg

and Mn in solid (Al)

149


MSIT®

Al–Mg–Mn

Mg Mg 90.00Mn 10.00Al 0.00

Mg 90.00Mn 0.00Al 10.00 Data / Grid: at.%

Axes: at.%

(Mg)

(Mg)+(βMn)(?)

(Mg)+(βMn)(?)+X

(Mg)+X

(Mg)+γ+X

(Mg)+γ

90

10

10

Mg Mg 80.00Mn 20.00Al 0.00

Mg 80.00Mn 0.00Al 20.00 Data / Grid: at.%

Axes: at.%

(Mg)

(Mg)+(βMn)(?)

(Mg)+(βMn)(?)+X

(Mg)+X

(Mg)+γ+X

(Mg)+γ

Fig. 6: Al-Mg-Mn.


the Mg corner at

200°C, X is an Al-Mn

binary compound

Fig. 5: Al-Mg-Mn.


the Mg corner at

400°C, X is a Al-Mn

binary compound

150


MSIT®

Al–Mg–Ni

Aluminium – Magnesium – Nickel

Elena L. Semenova

Literature Data

The Al-Mg-Ni system has been examined first in 1924. From the results of thermal analysis and

metallography [1924Fus] concluded that the Mg2Al3-NiAl3 section is a quasibinary one. In [1934Fus] Fuss

presented a projection of the liquidus surface in the Al-Mg2Al3-NiAl area showing the lines of double

saturation on it. An essential conclusion was that a ternary eutectic equilibrium does not exist in the shown

part of the phase diagram. However, [1943Mon, 1944Cha, 1952Han] reported that the invariant eutectic

equilibrium exists and is reached independently of the heat treatment and the compositions of the phases,

except of solid solution of magnesium in aluminium. These conclusions were based on experimental data

obtained on as-cast, annealed and rapidly quenched alloys; their liquidus projection is essentially different

from the one without the eutectic invariant reaction proposed by [1934Fus].

[1968Var] studied the structure of the Al-Mg-Ni alloys containing 1 at.% Ni in as-cast conditions. The

intermetallic phases were separated by high temperature centrifuging and identified by X-ray analysis. As

a result, the AlNi3 and Al3Ni2 phases were found to coexist in the alloy 1Ni-15Mg-Al (at.%).

The assessment by [1993Pri] took into account the works published up to 1991 and deals with the Al-rich

part of the Al-Mg-Ni ternary system Al-Mg2Al3-Ni2Al3.

Later experimental investigations of the ternary system were mainly motivated by the search for new

hydrogen storage materials [1998Ori, 2000Yua, 2000Aiz, 2001Gua]. From these studies information on

new ternary phases was obtained. [1998Ori] examined the crystallization processes of Alx-Mg1-x-Ni alloys

which were mechanically alloyed under an argon atmosphere by planetary ball milling for 4800 min at

ambient temperature and 400 rpm. A phase with CsCl type crystal structure was found in alloys with

x = 0.3-0.5 and an amorphous phase formed in alloys with x < 0.2.

[2000Yua] synthesized Alx-Mg2-x-Ni (x = 0, 0.1, 0.2, 0.3, 0.4, 0.5) samples by a diffusion method. Mixtures

of pure Al, Mg and Ni powders were grounded and pressed into pellets under a pressure of 30 MPa. The

pellets were annealed at 540-550°C for 4 h and then cooled to room temperature. X-ray diffraction and SEM

were applied to investigate their structure. A new phase of cubic crystal structure of Ti2Ni type was

observed in the alloys, so that with x = 0.5 only this phase and a trace of magnesium were detected.

[2001Gua] studied by X-ray diffraction the Ni2Mg3Al ternary alloy prepared from components of purities

better than 99.95 % by compacting their mixtures at 30 MPa and annealing them at 540-550°C for 4 h under

0.5 MPa argon atmosphere. The composition of the alloy prepared coincided actually with the composition

of a new ternary phase found in the investigation by [2000Yua]. [2001Gua] confirmed the existence of the

new ternary phase with the composition Ni2Mg3Al and studied its crystal structure using more advanced

X-ray techniques. As a result, the crystal structure of Ni2Mg3Al is established and described in more detail

than by [2000Yua].

[1991Han] addressed some thermodynamic aspects on the effect that aluminium has on magnesium-nickel

melts in presence of 3.8-8.6 10-4 mass% O. [2000Aiz] studied the effect that the substitution of aluminium

by magnesium has on hydrogen absorption by a material based on Mg2Ni.

Binary Systems

The Al-Mg and Al-Ni binary phase diagrams are accepted from [2003Luk], [2003Sal], respectively. The

Mg-Ni phase diagram is accepted from [1998Jac]. [1998Jac] made a thermodynamic assessment of the

Mg-Ni binary system using the experimental characteristics of the Mg-Ni phase diagram from [1934Hau,

1978Bag, 1996Mic]. The calculated phase diagram is in a good agreement with the data from the

experimental works.

151


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Al–Mg–Ni

Solid Phases

The data on the relevant binary phases and ternary phases are listed in Table 1. [2001Gua, 2000Yua] found

a new ternary phase of the same stoichiometry Ni2Mg3Al; its structural characteristics were determined and

described in detail by [2001Gua]. Although the ternary alloys in both works were prepared in similar ways

the Ni2Mg3Al alloy contained different phases in addition to the main phase. Therefore, the real

composition of the compound discovered may differ slightly from the stoichiometry given.


At least one invariant four-phase equilibrium and one three-phase equilibrium exist in the ternary Al-Mg-Ni

system, besides those in the adjacent binary systems. They are in the region of aluminium-rich alloys. The

four-phase equilibrium is of eutectic type at a temperature of 449°C [1944Cha, 1952Han, 1993Pri]. The

temperature of this equilibrium is assumed to be only by a few tenths of a degree lower than that of the

binary eutectic reaction L (Al)+Mg2Al3, which is reliably confirmed to be at 450.5°C [2003Luk]. Type and

temperature of the three-phase equilibrium however are not firmly established. It is of eutectic nature and

takes place at a temperature between 449°C, where the four-phase eutectic equilibrium is, and 552°C the

melting temperature of Mg2Al3, [1993Pri].

The characteristics of the three-phase and four-phase invariant equilibria are listed in Table 2 according to

[1993Pri] with some correction for (Al) and Mg2Al3 by [2003Luk]. Concentration of the liquid phase in the

three-phase invariant equilibrium is not determined exactly, but taking into account its temperature it is

reasonable to assume that it is close to the L (Al)+Mg2Al3 eutectic point in the binary Al-Mg system. The

reaction scheme for Al-NiAl3-Mg2Al3 region is shown in Fig. 1.

Liquidus, Solidus Surfaces

The liquidus surface of the Al-Mg-Ni system in Al-NiAl-Mg2Al3 region is shown in Fig. 2. It is a

compilation of the [1952Han, 1934Fus] data with some corrections drawn out that the next phase after

NiAl3 should be Ni2Al3 [1968Var, 2003Sal], rather than NiAl2, as it was proposed by [1934Fus]. The

temperatures of the invariant reactions in the binary systems are also corrected to comply with the today

accepted binary descriptions of Al-Mg and Al-Ni [2003Luk, 2003Sal].

The projection of the solidus surface in the Al-Mg2Al3-NiAl3 region is plotted in Fig. 3 based on [1952Han]

with correction of the (Al) and Mg2Al3( ) homogeneity ranges by [2003Luk]. The Ni2Al3 homogeneity

range is shown according to [2003Sal].


The statement of [1924Fus] that the Mg2Al3-NiAl3 section is a quasibinary one can not be correct taking

into account the Al-Ni phase diagram [2003Sal], where the NiAl3 phase is shown to form by a peritectic

reaction from liquid and Ni2Al3.

Figure 4 gives the NiAl3-Mg2Al3 temperature-concentration cut constructed using the data of [1952Han,

2003Luk, 2003Sal]. It can be considered as a quasibinary one only below the solidus temperature of the

alloys and within the part between Mg2Al3 and the edge of the Ni2Al3 primary crystallization surface

including the e3 eutectic point.

Thermodynamics

[1991Han] showed that activity of magnesium, containing 3.8-8.6·10-4 % O, in nickel melts increases with

addition of aluminium.


NiMg2 base alloys with addition of Al are candidate materials for hydrogen storage [1998Ori].

Electrochemical capacity and live-cycles of NiMg2-xAlx (0 x 0.5) alloys during absorption and

desorption of hydrogen increase with increasing Al contents, due to increasing amount of the Ni2Mg3Al

152


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Al–Mg–Ni

phase in the alloy [2000Yua]. Addition of Al also improves the corrosion resistance of the NiMg2-xAlxalloys to a certain degree because an Al2O3 oxide layer forms on the surface. The corrosion rate of the

ternary alloys is lower than that of NiMg2 [2000Yua]. Chemical modification of NiMg2 alloy by aluminium

addition to (NiMg1.8Al0.2) is expected not to lead to significant reduction of onset temperature for hydrogen

absorbing [2000Aiz].

NiMg1-xAlx phase with CsCl type crystal structure dissolves hydrogen interstitially without any structural

transformation [1998Ori].

References

[1924Fus] Fuss, V., “On the Constitution of Ternary Al Alloys” (in German), Z. Metallkd., 16, 24,


[1934Fus] Fuss, V., “Metallography of Al and its Alloys”, Berlin, The Sherwood Press. Inc.,

Cleveland, 142-143 (1934) (Equi. Diagram, Experimental, 1)

[1934Hau] Haughton, J.L., Payne, R.I., J. Inst. Met., 54, 275-283 (1934) quoted by [1998Jac]

(Thermodyn.)

[1943Mon] Mondolfo, L., “Al-Mg-Ni, Aluminium-Magnesium Nickel”, in “Metallography of

Aluminium Alloys”, John Wiley and Sons, Inc., New-York - London, 101-102 (1943) (Equi.

Diagram, Review, 1)

[1944Cha] Chao, H.L., “On the Ternary System Al-Mg-Ni”, Thesis, Berlin Techn. Hochschule (1944)


[1952Han] Hanemann, H., Schrader, A., “Examples for the Crystallization of Ternary Systems” (in

German), Atlas Metallographicus, 3(2), 120-122 (1952) (Equi. Diagram,

Experimental, #, *)

[1968Var] Varich, N.I., Litvin, B.N., “Structure of Phases in the Aluminium-Magnesium System

Containing Transition Metals” (in Russian), Izv. Akad. Nauk SSSR, Met., 6, 179-182 (1968)

(Experimental, 4)

[1978Bag] Bagnoud, P., Feschotte, P., “The Binary Systems Magnesium-Copper and Magnesium -

Nickel, Especially the Nod-Stoechiometry of the MgCu2 and MgNi2 Laves Phases” (in

French), Z. Metallkd., 69, 114-120 (1978) (Crys. Structure, Equi. Diagram.

Experimental, 24)

[1991Han] Han, Q., Wang, C., “Equilibrium of Mg-O and the Effect of Fe, Al and Cr on the Activity

of Mg in Molten Nickel”, Beijing Keji Dexue Xuebao, 13(5), 461-466 (1991) (Experimental,

Thermodyn., 4)

[1993Pri] Prima, S., “Aluminium-Magnesium-Nickel”, MSIT Ternary Evaluation Program, in MSIT



Assessment, 10)

[1996Mic] Micke, K., Isper, H., “Thermodynamic Properties of Liquid Magnesium-Nickel Alloys”,

Monatsh. Chem., 127, 7-13 (1996) (Equi. Diagram, Experimental, Thermodyn., 18)

[1998Jac] Jacobs, M.H.G., Spencer, P.I., “A Critical Thermodynamic Evaluation of the System

Mg-Ni”, Calphad, 22(4), 519-525 (1998) (Equi. Diagram, Review, Thermodyn., #, *, 30)

[1998Ori] Orimo, I.S., Ikeda, K., Fujii, H., “B2-Phase Formation and Hydriding Properties of

(Mg1-xAlx)Ni (x = 0~0.5)”, J. Alloys Compd., 266, L1-L3 (1998) (Crys. Structure,

Experimental, 10)

[2000Aiz] Aizawa, T., “Solid-State Synthesis of Magnesium Base Alloys”, Mater. Sci. Forum,

350-351, 299-310 (2000) (Experimental, 22)

[2000Yua] Yuan, H.T., Wang, L.B., Cao, R., Wang, Y.J., Zhang, Y., Yan, D.Y., Zhang, W.H.,

Gong, W.L., “Electrochemical Characteristics of Mg2-xAlxNi (0<x<0.5) Alloys”, J. Alloys

Compd., 309, 208-211 (2000) (Crys. Structure, Electrochem. Prop., Experimental, 8)

153


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Al–Mg–Ni

[2001Gua] Guanglie, L., Linshen, C., Lianbang, W., Huantang, Y., “Study on the Phase Composition

of Mg2-xMxNi (M = Al, Ti) Alloys”, J. Alloys Compd., 321(1), L1-L4 (2001) (Crys.


[2003Luk] Lukas, H.L., Lebrum, N., “Al-Mg (Aluminium-Magnesium)”, MSIT Binary Evaluation



Assessment, 49)

[2003Sal] Saltykov, P., Cornish, L., Cacciamani, G., “Al-Ni (Aluminium-Nickel)”, MSIT Binary





Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Al)

< 660.452

cF4

Fm3m

Cu

a = 404.96

a = 410.5 0.8

at 25°C [Mas2]

dissolves 0.01 at.% Ni at 639.9°C

[2003Sal] and 18.6 at.% Mg at 450.5°C

[2003Luk]

at 445°C in the alloy with 18.6 at.% Mg

[1952Han]

(Mg)

< 650

hP2

P63/mmc

Mg

a = 320.94

c = 521.07

at 25°C [Mas2]

dissolves 11.5 at.% Al at 436°C

[2003Luk] and < 0.04 mol% Ni at 500°C

[1934Hau]

(Ni)

< 1455

cF4

Fm3m

Cu

a = 352.40 at 25°C [Mas2]


[2003Sal] < 0.2 mol% Mg at 1100°C

[1998Jac]

, Mg17Al12

< 458

cI58

I43m

Mn

a = 1054.38 at 41.4 at.% Al [V-C2]

39.5 to 51.7 at.% Al [2003Luk]

, Mg2Al3< 452

cF1168

Fd3m

Mg2Al3

a = 2816 - 2824 60 to 62 at.% Al [2003Luk]

, Mg23Al30

410 - 250

hR159

R3

Mn44Si9

a = 1282.54

c = 2174.78

56.3 at.% Al [2003Luk]

NiAl3< 856

oP16

Pnma

NiAl3

a = 661.3 0.1

b = 736.7 0.1

c = 481.1 0.1

[2003Sal]

Ni2Al3< 1138

hP5

P3m1

Ni2Al3

a = 402.8

c = 489.1

36.1 to 39.8 at.% Ni [2003Sal]

Ni3Al4< 702

cI112

Ia3d

Ni3Ga4

a = 1140.8 0.1 [2003Sal]

154


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Al–Mg–Ni


NiAl

< 1651

cP2

Pm3m

CsCl

a = 286.0

42.1 to 71.3 at.% Ni [2003Sal]

Ni5Al3< 723

oC16

Cmmm

Pt5Ga3

a = 753

b = 661

c = 376

63 to 68 at.% Ni

at 63 at.% Ni

[2003Sal]

Ni3Al

< 1372

cP4

Pm3m

AuCu3

a = 356.77 75.4 to 76.3 at.% Ni [2003Sal]

NiMg2

< 759.31

hP18

P6222

NiMg2

a = 520.5 0.1

c = 1320 6

[V-C2]

[1998Jac]

Ni2Mg

< 1147.60

hP24

P63/mmc

Ni2Mg

a = 482.4 0.2

c = 1582.6

66.2 at.% at 759.31°C to 67.34 at.% Ni

at 1095.28°C [1998Jac], [V-C2]

* NiMg1-xAlx cP2

Pm3m

CsCl

In alloys with x = 0.3 - 0.5 prepared by

mechanical alloying [1998Ori]

* NiMg1-xAlx amorphous phase In the alloys with x < 0.2 prepared by

mechanical alloying [1998Ori]

* Ni2Mg3Al cF96

Fd3 m

derived from

Ti2Ni

a = 1154.74 0.02 [2001Gua]


Al Mg Ni

L Mg2Al3 + NiAl3 449 - 552 e3 L

Mg2Al3NiAl3

~60

61

75

~40

39

0

?

0

25

L Mg2Al3 + NiAl3 + (Al) 449 E L

Mg2Al3NiAl3(Al)

64.6

61

75

81.39

34.6

39

0

18.6

0.8

0

25

0.01

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

155


MSIT®

Al–Mg–Ni

Fig. 1: Al-Mg-Ni. Reaction scheme in the partial Al-Mg2Al

3-NiAl

3system

Al-Mg A-B-C

l β + (Al)

450.5 e2

L β + NiAl3

449<T<452 e3

Al-Mg-Ni

L β+NiAl3+(Al)449 E

β+NiAl3+(Al)

Al-Ni

l NiAl3+(Al)

644 e1

10

20

30

40

10 20 30 40

60

70

80

90

Ni 50.00Mg 0.00Al 50.00

Ni 0.00Mg 50.00Al 50.00


Axes: at.%

(Al)e

1,644

p2,856

p1,1138

NiAl

Ni2Al

3

NiAl3

β

E

e3

e2,450.5°C

Fig. 2: Al-Mg-Ni.

Liquidus surface in

the Al-rich alloys

156


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Al–Mg–Ni

10

20

30

40

10 20 30 40

60

70

80

90

Ni 50.00Mg 0.00Al 50.00

Ni 0.00Mg 50.00Al 50.00


Axes: at.%

Ni2Al

3

NiAl3

β

β+(Al)+NiAl3

(Al)

449°C

20 10

400

500

600

700

800

900

1000

1100

1200

Mg 0.00Ni 25.00Al 75.00

Mg 40.00Ni 0.00Al 60.00Ni, at.%

Te

mp

era

ture

, °C

Mg2Al3+NiAl3

L+NiAl3

L+NiAl3+Ni2Al3

L+Ni2Al3856°C

449<T<452 452°C

L

e3

~1116°C

Fig. 3: Al-Mg-Ni.

Solidus surface of the

partial

Al-Mg2Al3-NiAl3system

Fig. 4: Al-Mg-Ni.

The NiAl3-Mg2Ni3vertical section

157


MSIT®

Al–Mg–Sc

Aluminium – Magnesium – Scandium

Evgeniya V. Lysova, updated by Rainer Schmid-Fetzer and Alexander Pisch

Literature Data

[1976Tur] investigated Al rich alloys of the Al-Mg-Sc system in the range up to 26.0 mass% Mg and 3.0

mass% Sc by thermal and metallographic analysis. The starting materials were 99.99% pure Al, 99.91%

pure Mg, and scandium which contained 0.03% Cu, 0.01% Fe, 0.01% Ca, and 0.01% Si. The alloys were

melted in an electrical resistance furnace in corundum crucibles under a layer of flux composed of 50% LiCl

and 50% KCl and cast into thick-walled Cu moulds. The castings were homogenized at 400°C for 30 h,

deformed under various conditions, annealed in evacuated ampoules at 640°C, 550°C and 430°C for 100,

360 and 600 h, respectively and finally quenched in water. Etching of the specimens was possible with a

solution of 25 ml HNO3+1.5 ml HCl+1 ml HF per 100 ml of water. A partial isothermal section of the

Al-Mg-Sc system at 430°C and two vertical sections have been established. There are indications that an

invariant equilibrium exists at approximately 447°C.

One alloy Al-2.5Mg-0.4Sc (mass%) has been studied by extensive X-ray diffraction after annealing at

150°C (10 h) and 350°C (2 h) and additional irradiation by electrons with energy 2.3 MeV [1984Rep].

[1989Odi, 1991Odi] examined the Al-Mg rich part of the system by differential thermal analysis,

micrographic and X-ray analysis. An isothermal section at 400°C as well as a partial liquidus surface has

been proposed for this region.

The Al-Mg rich part of this system has also been investigated by [1999Gro]. Based on thermodynamic

equilibrium calculations, key samples have been defined to determine the isothermal section at 350°C, the

invariant reactions in the Mg rich part as well as the liquidus surface. Starting materials were 99.999% pure

Al and 99.99% pure Mg and Sc. Two types of samples have been prepared: binary Al-Sc master alloys by

levitation or e-beam melting mixed with Mg for solid state reactions and ternary Al-Mg-Sc samples, melted

in an induction furnace in sealed Ta crucibles under Ar atmosphere. Both types of samples have been

annealed for 2 months to reach equilibrium at 350°C. The samples were analyzed by X-ray diffraction,

optical and electronic microscope, electron microprobe analysis as well as differential thermal analysis.

Binary Systems

For the binary systems the following versions have been adopted: Al-Mg [1998Lia], which is essentially the

same as [2003Luk], Al-Sc [1999Cac] and Mg-Sc [1998Pis].

Solid Phases

No ternary phases have been detected in the investigated range of concentrations. Although magnesium

dissolves a considerable amount of Al or Sc, the ternary solubility is extremely small. There is a small

solubility of magnesium in the four intermetallic Al-Sc compounds at 350°C: 4 at.% Mg in Sc2Al, 12 at.%

in ScAl, 1-2 at.% in ScAl2 and 5 at.% in ScAl3 [1999Gro]. The solubility of scandium in the binary Al-Mg

phases is small. A value of 1 at.% at 350°C has been found for (Mg17Al12) [1999Gro]. No information has

been given for the solubility of aluminium in the binary Mg-Sc phases. Crystallographic data on the binary

compounds are given in Table 1.


Based on thermodynamic calculations by extrapolation of the three binary systems, [1999Gro] identified 14

invariant reactions. The Al-Mg-Sc system is characterized by a liquid miscibility gap in the ternary with

associated invariant reactions of the eutectic type and a series of U type reactions in the Mg-rich corner. The

measured temperature of 1165 40°C for the reaction L L´+ScAl+ScAl2 is reproduced by the calculations.

Two of the five U type reactions (U3: L+ScAl ScAl2+(Mg) and U5: L+ScAl2 ScAl3+(Mg)) have been

measured and the temperatures agree also well with calculated ones (U3: 590 10°C; U5: 660 10°C). All

158


MSIT®

Al–Mg–Sc

other ternary invariant reactions including the binary Al-Mg are degenerated and close to the binary values.

The reported L (Al)+ScAl3+ (Mg2Al3) at 447°C by [1976Tur] is in agreement with [1999Gro]. The

values of [1991Odi] have not been considered due to an incorrect isothermal section at 400°C with

subsequent erroneous invariant equilibrium reactions. The invariant equilibria according to [1999Gro] are

summed up in Table 2.

Liquidus Surface

The liquidus surface, based on thermodynamic calculations and experimental DTA results [1999Gro] is

drawn in Figs. 1a, 1b. The results of [1991Odi] differ considerably and are not considered. One possible

explanation for this discrepancy is the use of sealed quartz ampoules by these authors. Quartz is known to

easily react with magnesium.

Isothermal Sections

Figure 2 shows the isothermal section at 350°C as determined by [1999Gro]. (Mg), depending on the mutual

solubilities of Sc and Al, is in equilibrium with ScAl, ScAl2 and ScAl3. All binary Al-Mg phases are only

in equilibrium with ScAl3. The isothermal section at 430°C for the Al rich region according to [1976Tur]

is reproduced in Fig. 3. Additions of scandium substantially decrease the solubility of magnesium in

aluminium. The point of ultimate saturation of aluminium with scandium and magnesium was originally

observed at 10.5 mass% (11.6 at.%) Mg and 0.01 mass% (0.006 at.%) Sc [1976Tur]. Due to the low Mg

solubility in Al, as shown by [1976Tur] at 430°C (~11 mass% Mg), the (Al)+ScAl3+ (Mg2Al3) vertex has

been shifted to 12.5 mass% Mg and 0.01 mass% Sc in order to meet the accepted binary value of 13 mass%

Mg in Al given by [1981Sch]. The two phase (Al)+ (Mg2Al3) region is very narrow and closely adjoins the

Al-Mg side, whereas the two phase (Al)+ScAl3 region is quite wide.

The proposition of [1984Rep] does not fit with Fig. 2. In both the irradiated and non-irradiated alloys faint

X-ray reflections were observed in addition to those of the (Al) matrix. These were attributed to phases

Mg2Al3, “Mg3Al4” and “Mg5Al8”. This is a strange proposition since the alloy Al-2.5Mg-0.4Sc (mass%)

is located clearly inside the (Al)+ScAl3 two phase field, which was confirmed in over 50 alloys by

[1976Tur]. The work of [1984Rep] does not provide preparation details and does not refer to the prior

results of [1976Tur].


Figures 4 and 5 show two calculated isopleths [1999Gro] in agreement with the experimental sections from

[1976Tur]. The two diagrams correspond to the sections between 17Mg-Al (mass%) and 1.0Sc-Al (mass%)

and 22.0Mg-Al (mass%) and 2.0Sc-Al (mass%), respectively. Both sections intersect two regions of

primary crystallization: of the aluminium solid solution (Al) and of the ScAl3 compound. The calculated

ScAl2-Mg section from [1999Gro], confirmed by selected experiments, is drawn in Fig. 6. One notices the

steep liquidus in the Mg-rich part of the diagrams.


Additions of small amounts of scandium has been found to significantly improve yield stress, fatigue

strength and resistance against microcrack growth of Al-Mg alloys [1981Dri, 1984Dri, 1990Saw, 1992Ela,

1997Rod]

References

[1976Tur] Turkina, N.I., Kuzmina, V.I., “Phase Reactions in Al-Mg-Sc Alloys (up to 26 % Mg and

3 % Sc)”, Russ. Metall., (4), 179-183 (1976), translated from Izv. Akad. Nauk SSSR, Met.,

(4), 208-212 (1976) (Equi. Diagram, Experimental, #, 9)

159


MSIT®

Al–Mg–Sc

[1981Dri] Dritz, M.E., Pavlenko, S.G., Toropova, L.S., Bykov, Yu.G., Ber, L.B., “Mechanism of

Scandium Effect to Increasing Strength and Termal Stability of Al-Mg Alloys”, Dokl. Akad.

Nauk SSSR, 257(2), 353-356 (1981) (Crys. Structure, Experimental, 11)

[1981Sch] Schuermann, E., Voss, H.J., “Investigation of the Liquid Equilibria of Mg-Li-Al Alloys:

Part 4. Liquid Equilibria of the Mg-Al Binary System” (in German), Giessereiforschung,

33, 43-46 (1981) (Equi. Diagram, Experimental, #, 17)

[1984Dri] Dritz, M.E., Ber, L.B., Bykov, Yu.G., Toropova, L.S., Anastaseva, G.K., “Ageing Alloy

Al-0.3 at.% Sc”, Phys. Met. Metallogr., 57, 118-126 (1984) (Experimental)

[1984Rep] Repnikova, Ye.A., Malinenko, I.A., Chudinova, S.A., Toropova, L.S.,

Ustinovshchikov, V.M., “Influence of Electron Irradiation on Decomposition of Alloy

Al-Mg-Sc”, Phys. Met. Metallogr., 57(3), 103-106 (1984) (Crys. Structure, Experimental,

3)

[1985Sch] Schuster, J.C., Bauer, J., “The Ternary System Sc-Al-N and Y-Al-N”, J. Less-Common

Met., 109, 345 (1985) (Experimental, Crys. Structure)

[1989Gsc] Gscheidner K.A., Calderwood, F.M., “ The Al-Sc (Aluminium-Scandium) System”, Bull.

Alloy Phase Diagrams, 10, 34-36, (1989) (Crys. Structure, Equi. Diagram, Review, #, 18)

[1989Odi] Odinaev, K.O., Ganiev, I.N., Kinzhibalo, V.V., Kotur, B.Y., “Phase Diagram of the

Aluminum-Magnesium-Scandium System in the 0-33.3 at.% Sc Interval at 673K”, Dokl.

Akad. Nauk Tadzh. SSR, 32(1), 37-38 (1989) (Experimental, Equi. Diagram, 4)

[1990Saw] Sawtell, R.R., Jensen, C.L., “Mechanical Properties and Microstructures of Al-Mg-Sc

Allyos”, Met. Trans., 21A, 421-430 (1990) (Crys. Structure, Mechan. Prop.)

[1991Odi] Odinaev, K.O., Ganiev, I.N., Vakhobov, A.V., “Quasi Binary Sections and the Liquids

Surface of the Al-Mg-Sc System”, Dokl. Akad. Nauk SSSR, Met., (4), 195-197 (1991)

(Experimetnal, Equi. Diagram, 10)

[1992Ela] Elagin, V.I., Zakharov, V.V., Rostova, T.D., “Scandium-Alloyed Aluminium Alloys”, Met.

Sci. Heat Treat., 34, 37-45 (1992) (Experimental)

[1997Rod] Roder, O., Wirtz, T., Gysler, A., Luetjering, G., “Fatigue Properties of Al-Mg Alloys with

and without Scandium”, Mater. Sci. Eng. A, A234-236, 181-184 (1997) (Experimental, 5)

[1997Su] Su, H.-L., Harmelin, M., Donnadieu, P., Baetzner, C., Seifert, H.J., Lukas, H.L.,

Effenberg, G., Aldinger, F., “Experimental Investigation of the Mg-Al Phase Diagram from

47 to 63 at.% Al”, J. Alloys Compd., 247, 57-65 (1997) (Crys. Structure, Experimental,

Equi. Diagram, #, *, 20)

[1998Lia] Liang, P., Su, H.-L., Donnadieu, P., Harmelin, M., Quivy, G., Ochin, P., Effenberg, G.,

Seifert, H.J., Lukas, H.-L., Aldinger, F., “Experimental Investigation and Thermodynamic


(1998) (Equi. Diagram, Thermodyn., Experimental, Assessment, *, 33)

[1998Pis] Pisch, A., Schmid-Fetzer, R., Cacciamani, G., Riani, P., Saccone, A., Ferro, R., “Mg-rich

Phase Equilibria and Thermodynamic Assessment of the Mg-Sc System”, Z. Metallkd., 89

(7), 474-477 (1998) (Equi. Diagram, Experimental, *,11)

[1999Cac] Cacciamani, G., Riani, P., Saccone, A., Ferro, R., Pisch, A., Schmid-Fetzer, R.,

“Thermodynamic Measurements and Assessment of the Al-Sc System”, Intermetallics, 7,

101-108 (1999) (Experimental, Equi. Diagram, Thermodyn., 26)

[1999Gro] Groebner, J., Schmid-Fetzer, R., Pisch, A., Cacciamani, G., Riani, P., Parodi, N.,

Borzone, G., Saccone, A., Ferro, R., “Experimental Investigations and Thermodynamic

Calculations in the Al-Mg-Sc System”, J. Phase Equilib., 90(II), 872-880 (1999)

(Experimental, Calculation, Themodyn., Equi. Diagram, 23)




Assessment, 49)

160


MSIT®

Al–Mg–Sc


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Al)

< 660.452

cF4

Fm3m

Cu

a = 404.96 pure at 25°C [Mas2]

(Mg)

< 650

hP2

P63/mmc

Mg

a = 320.94

c = 521.05

pure at 25°C [Mas, V-C]

( Sc)

1541 - 1337

cI2

Im3m

W

a = 373 [1989Gsc]

( Sc)

< 1337

hP2

P63/mmc

Mg

a = 330.90

c = 527.33

at room temperature [Mas, V-C]

, Mg2Al3 452

cF1168

Fd3m

Mg2Al3

a = 2816 to 2824 60-62 at.% Al [2003Luk]


[2003Luk]

, Mg17Al12

< 458

cI58

I43m

Mn

a = 1054.38 at 41.4 at.% Al [V-C2]

39.5 to 51.5 at.% Al [2003Luk]

, Mg23Al30

410 - 250

hR159

R3

Mn44Si9

a = 1282.54

c = 2174.78

54.5 to 56.5 at.% Al [2003Luk]

ScAl3< 1320

cP4

Pm3m

AuCu3

a = 410.3

a = 411.6

[1989Gsc]

3.7 at.% Mg [1999Gro]

ScAl2< 1370

cF24

Fd3m

Cu2Mg

a = 758.2

a = 757.8

a = 757.5

[1989Gsc]

[1999Cac]

1.0 at.% Mg [1999Gro]

ScAl

< 1300

oP8

Cmcm

CrB

cP2

Pm3m

CsCl

a = 398.8

b = 988.2

c = 365.2

a = 345.0

a = 339.1

a = 343.2

a = 344.7

[1985Sch]

[1989Gsc]

4.9 at.% Mg [1999Gro]

9.7 at.% Mg

11.5 at.% Mg

Sc2Al

< 1300

hP6

P63/mmc

Ni2In

a = 488.8

c = 616.6

a = 488.5

c = 615.7

[1989Gsc]

4 at.% Mg [1999Gro]

(Mg-Sc system)

< 520

cP2

Pm3m

CsCl

- [Mas]

161


MSIT®

Al–Mg–Sc

Table 2: Invariant Equilibria According to [1999Gro]

Invariant Reaction T [°C] Type

L' L + ScAl + ScAl2 1165 M1

L' L + ScAl + Sc2Al 1144 (calc.) M2

L + Sc2Al ScAl + ( Sc) 787 (calc.) U1

L + ( Sc) ScAl + (Mg) 709 (calc) U2

L + ScAl ScAl2 + (Mg) 660 U3

( Sc) + ScAl Sc2Al + (Mg) 624 (calc.) U4

L + ScAl2 ScAl3 + (Mg) 590 U5

( Sc) (Mg) + ( Sc), Sc2Al 481 (calc.) D1

L (Al) + (Mg2Al3), ScAl3 450 (calc.) D2

L (Mg17Al12) + (Mg2Al3), ScAl3 449 (calc.) D3

L (Mg) + (Mg17Al12), ScAl3 436 (calc.) D4

Sc2Al + (Mg) ScAl + (ScMg) 410 (calc.) U6

(Mg17Al12) + (Mg2Al3) , ScAl3 410 (calc.) D5

(Mg17Al12) + (Mg2Al3), ScAl3 250 (calc.) D6

20

40

60

80

20 40 60 80

20

40

60

80

Al Mg

Sc Data / Grid: at.%

Axes: at.%

p1

e4

e1

e2

e3

U1

U2

U3

U5D

4D

3D

2

M1'

M2'

M2

M1

L, + L

1140

1220

1140

1220

(βSc)

AlSc

AlSc2

Al3Sc

Al2Sc

AlSc

Fig. 1a: Al-Mg-Sc.

Calculated liquidus

surface

162


MSIT®

Al–Mg–Sc

10

20

80 90

10

20

Al 30.00Mg 70.00Sc 0.00

Mg

Al 0.00Mg 70.00Sc 30.00 Data / Grid: at.%

Axes: at.%

U5

Al2Sc

Al3Sc

AlSc

(Mg)

(βSc)

U2

U1

U3

AlSc2

Fig. 1b: Al-Mg-Sc.

Enlarged schematic

Mg rich corner of the

calculated liquidus

surface

20

40

60

80

20 40 60 80

20

40

60

80

Al Mg

Sc Data / Grid: at.%

Axes: at.%

MgSc

AlSc2

AlSc

Al2Sc

Al3Sc

β ε γ

(Mg)not investigated

AlSc+AlSc2 +MgSc

β+Al3Sc+ε

AlSc+MgSc+(M

g)

AlSc2+AlSc+(Mg)Al

3Sc+AlSc2+(Mg)

γ+Al3Sc+(Mg)

ε+Al3 Sc+γ

Fig. 2: Al-Mg-Sc.

Partial isothermal

section at 350°C

163


MSIT®

Al–Mg–Sc

1.50

40

α + ScAl3α

α + Mg Al2 3

α + Mg Al + ScAl2 3 3

Sc, at. %

Mg,at.

%

Al

0

90

400

500

600

700

800

Al 99.40Mg 0.00Sc 0.60

Al 81.50Mg 18.50Sc 0.00Al, at.%

Te

mp

era

ture

, °C L + Al3Sc

L + Al3Sc + (Al)

Al3Sc + (Al)

450

Thermal arrest [1976Tur]

L

Fig. 3: Al-Mg-Sc.

Partial isothermal

section at 430°C

Fig. 4: Al-Mg-Sc.

Calculated vertical

section from

0.6Sc-99.4Al to

18.5Mg-81.5Al

(at.%)

164


MSIT®

Al–Mg–Sc

60 50 40 30 20 10

250

500

750

1000

1250

Al 66.66Mg 0.00Sc 33.33

Mg

Al, at.%

Te

mp

era

ture

, °C

Thermal arrest [1999Gro]

L + Al2Sc

(Mg) + Al2Sc

(Mg) + Al2Sc + Al3Sc

L + Al2Sc + (Mg)

616

90 80

400

500

600

700

800

Al 98.80Mg 0.00Sc 1.20

Al 76.20Mg 23.80Sc 0.00Al, at.%

Te

mp

era

ture

, °C

Al3Sc + (Al)

L + Al3Sc + (Al)

L + Al3Sc

450

Thermal arrest [1976Tur]

Fig. 6: Al-Mg-Sc.

Calculated vertical

section from Al2Sc to

the Mg corner

Fig. 5: Al-Mg-Sc.

Calculated vertical

section from 98.8

Al-1.2 Sc to 76.2

Al-23.8 Mg (at.%)

165


MSIT®

Al–Mg–Si

Aluminium – Magnesium – Silicon

K.C. Hari Kumar, Nirupan Chakraborti, Hans-Leo Lukas, Oksana Bodak, Lazar Rokhlin

Literature Data

Al-Mg-Si alloys are being increasingly used in automotive and aerospace industries for critical structure

applications because of their excellent castability and corrosion resistance and, in particular, good

mechanical properties in the heat treated condition. These are known as 4xxx series of wrought alloys and

3xx.0 and 4xx.0 series of casting alloys. In these alloys, Mg is intentionally added to induce age hardening

through precipitation of Mg2Si, metastable phases or Guinier-Preston zones.

Several studies pertaining to the liquidus of the system are reported in the literature. [1921Han] studied

alloys containing 0 to 11 mass% Si by thermal analysis in the Al-rich corner. [1930Ota] studied 24 alloys

in the Al-rich corner, up to 20 mass% Si and 15 mass% Mg. [1931Dix] examined Al-rich alloys by thermal

analysis and metallography, while [1931Los] reported thermal analysis data for a total of 150 alloys

covering the entire composition range. [1935Saw] measured liquidus temperatures of 29 Mg-rich alloys (50

to 100 mass%) with Si contents up to 12 mass%. [1941Phi] examined the liquidus of certain Al-rich alloys.

[1958Gul] gave data on 4 alloys with 1 to 7 at.% Si and approximately equal amounts of Mg and Al.

[1976Fis, 1977Sch] again studied the liquidus of the whole system, primarily by thermal analysis and 15

vertical sections were reported along with the isotherms superimposed on the complete liquidus surface.

More recent investigations of the liquidus are due to [2001Goe, 2001Li, 2001Bar, 2002Bar], employing

thermal analysis.

The homogeneity range of the (Al) solid solution was measured using metallography by [1921Han,

1931Dix] and [1943Wes], who reported the data as solubilities of Mg2Si in (Al). [1936Kel] also gave

solubility values for Mg2Si and excess Mg. [1940Kuz] measured lattice parameters of the (Al) solid

solutions along different lines in the Gibbs triangle and deduced the boundaries of the homogeneity range.

Using dilatometry [1997Feu] reported two data points on Al solvus corresponding to the Al-Mg2Si section,

which is in reasonable agreement with data reported by [1931Dix] and [1940Kuz]. [1941Phi] constructed

several isothermal and vertical sections of the Al corner from metallographic measurements. An isothermal

section of the Al corner at 460°C was reported by [1948Axo] based on metallographic experiments.

[1988Rok] presented the (Mg) corner of the (Mg)+Mg2Si+Mg17Al12 three-phase field at 430, 400 and

300°C. The purity of the starting materials was 99.8 to 99.99 % for Mg and Al, but the Si in all investigations

contained some impurities (1 % Fe in an 80 % Al, 20 % Si master alloy [1921Han], 0.66 % Fe+0.1 % Ti

[1931Dix], 0.3 % Fe in Si [1936Kel], 0.28 % Fe+0.17 % Ca [1931Los], 1.5 % Fe [1941Phi], 1 % mainly

Fe [1977Sch] or 0.5 % of impurities not specified [1935Saw]).

Reports of thermodynamic measurements for this system are rather limited. The activity of Mg in liquid

phase was determined employing emf technique by [1979Seb] using alloys of nine different compositions

near the Al-rich corner at three temperatures, 700, 750 and 800°C. The enthalpy change for the ternary

eutectic reaction at 560°C was determined by [1980Bir] using DSC. [1986Lue] made a thermodynamic

optimization of the system and reported a single ternary interaction parameter for the liquid phase, based

mainly upon results of [1977Sch]. The data of [1977Sch], however, failed to extrapolate to the currently

accepted melting point of Si [1986Bul] due to the limited purity of the Si used in the experiments. The

deviation is about 12°C. [1991Zhi] determined the specific heats and enthalpies of the phase

transformations of Al-Si and Al-Mg-Si alloys by isothermal calorimetry measurements. The ternary system

was recalculated by [1992Cha] using phase stability values for pure elements recommended by [1989Din].

The calculated diagram agrees with the experimental data reasonably [1921Han, 1930Ota, 1931Dix,

1940Kuz, 1941Phi, 1943Wes, 1948Axo, 1958Gul, 1977Sch, 1979Seb]. [1993Rei] determined temperatures

of the secondary phase particle formation in two Al-Mg-Si alloys. One of the alloys had a composition

corresponding to the section Al-Mg2Si and the other alloy had composition with some excess of Si as

compared with Al-Mg2Si. The results of the experiments confirmed the phase diagram presented by

[1992Cha]. [1997Feu] carried out experiments in the Al corner of the system and updated the

166


MSIT®

Al–Mg–Si

thermodynamic assessment of the whole system. They also measured enthalpy of formation and fusion as

well as the heat capacity of Mg2Si employing calorimetry.

Binary Systems

Binary systems Al-Si [2003Luk1] and Al-Mg [2003Luk2] are from the MSIT Binary Evaluation Program.

The binary system Mg-Si is from [1997Feu].

Solid Phases

No stable ternary compounds have been reported. Several metastable phases were reported to form during

annealing of quenched supersaturated (Al) solid solutions [1999Mat, 2000Cay, 2001Mar, 2002Der,

2003Mar]. The stable and some metastable phases are listed in Table 1. The solid phase (Si) has only

negligible solubility for Al and Mg.


The Al-Mg2Si section is approximately pseudobinary. It is shown in Fig. 1, calculated from the dataset of

[1997Feu]. The calculation shows that the (Al) phase in the eutectic maximum contains more than twice as

much Mg than Si. Therefore, the section Al-Mg2Si is not exactly pseudobinary and shows an extended

three-phase field L+(Al)+Mg2Si [1997Feu, 2001Zha]. In fact it is reported that the real pseudobinary is

shifted more towards the Mg-rich region [1997Feu, 2001Bar, 2001Li, 2002Bar], located along the section

Mg2Si-Al97.2Mg2.4 [2001Goe]. The temperature of the so-called pseudobinary eutectic was given as 595°C

by [1930Ota, 1931Dix, 1997Feu, 2001Goe, 2001Li], as approximately 593°C by [1977Sch] and as

approximately 590°C by [1921Han, 1931Los, 1941Phi], and 597°C [2001Bar, 2002Bar]. Using the

thermodynamic model parameters reported by [1997Feu] it is calculated to be 594°C.


The reaction scheme (Fig. 2) proposed by [1992Cha] is calculated from the dataset of [1997Feu] with the

data of the binary Al-Mg intermediate phases replaced by those of [1998Lia] assuming zero solubilities of

Si in these phases. The calculated compositions of the phases in the reactions containing liquid are listed in

Table 2. The temperature of E1 was reported as 551°C [1930Ota, 1941Phi], [1931Los] ~557°C, ~550°C

[1977Sch], and 560°C [1980Bir]. Calculations by [1992Cha, 1997Feu] indicate E1 to be at 557°C. E4 was

reported to be at 437°C by [1935Saw] and 435°C by [1977Sch, 1988Rok], which are nearly identical to the

binary L (Mg)+ eutectic (e7). In the experimental investigation by [1997Feu], E1 is reported to be at

558°C.

Liquidus Surface

The liquidus surface is shown in Fig. 3, calculated from the dataset of [1997Feu]. Only the ternary equilibria

e3, E1, E4 are indicated and not E2, E3, which virtually coincide with the binary reactions e7, e8, respectively.

In Fig. 4 the Al corner of the liquidus surface is shown, calculated from the dataset of [1997Feu].

Isothermal Sections

Figures 5 and 6 show the isothermal sections at 600 and 550°C. At room temperature, all solid phases are

in equilibrium with Mg2Si. The calculated solidus and solvus isotherms of the (Al) solid solution are given

in Figs. 7 and 8, respectively, after [1997Feu]. The results of [1988Rok] could not be reproduced by the

calculation of [1997Feu] as it contradicts Henry’s rule for dilute solutions. Due to the small Si solubility in

(Mg) this rule predicts the Al solubility of (Mg) in the three-phase field (Mg)+ +Mg2Si to be very near to

that in the binary two-phase field (Mg)+ .

167


MSIT®

Al–Mg–Si


Two calculated vertical sections are shown in Figs. 9 and 10, originally reported by [1992Cha], but

recalculated here using the thermodynamic data set of [1997Feu]. Figures 11 to 13 show calculated vertical

sections in the Al corner as experimentally investigated and calculated by [1997Feu].

Thermodynamics

As mentioned in the section “Literature Data” thermodynamic information on the system is limited to the

emf studies [1979Seb] in the Al-rich liquid and the enthalpy of the eutectic reaction E1 [1980Bir]. The

enthalpy of melting of the eutectic reaction E1 is reported to be +26.4 kJ mol-1 [1980Bir].

References

[1921Han] Hanson, D., Gayler, M.L.V., “The Constitution and Age-Hardening of the Alloys of

Aluminium with Magnesium and Silicon”, J. Inst. Met., 26, 321-359 (1921) (Equi. Diagram,

Experimental, 4)

[1930Ota] Otani, B., “Siliciumin and its Structure” (in Japanese), Kinzuku no Kenkyu, 7, 666-686


[1931Dix] Dix, E.H., Keller, F., Graham, R.W., “Equilibrium Relations in Aluminium-Magnesium

Silicide Alloys of High Purity”, Trans. A.I.M.M.E., 404-420 (1931) (Equi. Diagram,

Experimental, 10)

[1931Los] Losana, L., “The Ternary System Al-Mg-Si” (in Italian), Metall. Ital., 23, 367-382 (1931)


[1935Saw] Sawamoto, H., “Equilibrium Diagram of the Magnesium-Rich Magnesium-Aluminium-

Silicon Ternary System” (in Japanese), Suiyokwai Shi, 8, 713-727 (1935) (Equi. Diagram,

Experimental, 25)

[1936Kel] Keller, F., Craighead, C.M., “Equilibrium Relations in Aluminium-Magnesium Silicide

Alloys Containing Excess Magnesium”, Trans. A.I.M.M.E., 122, 315-323 (1936) (Equi.


[1940Kuz] Kuznetsov, V.G., Makarov, E.S., “X-Ray Investigation of the Structure of Ternary Solid

Solutions of Magnesium and Silicon in Aluminium” (in Russian), Izv. Sekt. Fiz.-Khim.

Anal., 13, 177-190 (1940) (Equi. Diagram, Experimental, 18)

[1941Phi] Phillips, H.W.L., “The Constitution of Alloys of Aluminium with Magnesium and Silicon”,

J. Inst. Met., 67, 257-273 (1941) (Equi. Diagram, Experimental, 9)

[1943Wes] Westlinning, H., Klemm, W., “The Solubility of Mg2Si, Mg2Ge, Mg2Sn and Mg2Pb in

Aluminium” (in German), Z. Elektrochem., 49, 198-200 (1943) (Equi. Diagram,

Experimental, 4)

[1948Axo] Axon, H.J., Hume-Rothery W., “The Effect of 1% Silicon on the Constitution of Aluminium

Magnesium Manganese Zinc Alloys at 460°C”, J. Inst. Met., 74, 315-329 (1948) (Equi.


[1958Gul] Gul'din, I.T., Dokokina, N.V., “The Aluminium-Magnesium-Iron-Silicon System”, Russ. J.

Inorg. Chem., 3, 359-379 (1958), translated from Zh. Neorg. Khim., 3, 799-814 (1958)


[1976Fis] Fischer, A., “Investigations on Equilibria with Liquid and Mechanical Properties of the

Binary Al-Mg and Mg-Si Systems as well as of the Aluminium-Magnesium-Silicon Ternary

System” (in German), Thesis, University of Clausthal, F.R. Germany (1976) (Equi.


[1977Sch] Schürmann, E., Fischer, A., “Equilibria with Liquid in the Aluminium- Magnesium-Silicon

Ternary System, Part 3, Al-Mg-Si System” (in German), Giessereiforschung, 29, 161-165

(1977) (Equi. Diagram, Experimental, #, *, 14)

[1979Seb] Sebkova, J., Beranek, M., Halamkova, P., “Thermodynamic Properties of Liquid Al-Mg-Si

Alloys” (in Czech), Kovove Mater., 17, 137-143 (1979) (Experimental, Thermodyn., 12)

168


MSIT®

Al–Mg–Si

[1980Bir] Birchenall, C.E., Riechmann, A.F., “Heat Storage in Eutectic Alloys”, Metall. Trans. A,

A11, 1415-1420 (1980) (Thermodyn., Experimental, 13)

[1986Bul] “Melting Points of the Elements”, Bull. Alloy Phase Diagrams, 7, 602 (1986) (Review, 0)

[1986Lue] Lüdecke, D., “Phase Diagram and Thermochemistry of the Al-Mg-Si System”, Z. Metallkd.,

77, 278-283 (1986) (Equi. Diagram, Theory, Thermodyn., 33)

[1988Rok] Rokhlin, L.L., Pepelyan, A.G., “Phase Equilibria in the Mg-Al-Si System in the Magnesium

Rich Area” (in Russian), Izv. Akad. Nauk SSSR, Met., 176-179 (1988) (Equi. Diagram,

Experimental, 4)

[1989Din] Dinsdale, A.T., “SGTE Data for Pure Elements”, NPL Report DMA(A) 195, September

(1989) (Review, 20)

[1991Zhi] Zhiguang, H., Sinong, X., Guangzhong, W., Shaohua, M., “Measuring Heat of Thermal

Storage of Phase Change Metal” (in Chinese), Gongcheng Rewuli Xuebao (J. Eng.

Thermophys.), 12(1), 46-49 (1991) (Thermodyn., Experimental, 5)

[1992Cha] Chakraborti, N., Lukas, H.L., “Thermodynamic Optimization of the Mg-Al-Si Phase

Diagram”, Calphad, 16, 79-86 (1992) (Equi. Diagram, Review, Thermodyn., 27)

[1993Rei] Reiso, O., Ryum, N., Strid, J., “Melting of Secondary-Phase Particles in Al-Mg-Si Alloys”,

Metall. Trans. A, 24A, 2629-2641 (1993) (Equi. Diagram, Experimental, 13)

[1997Feu] Feufel, H., Gödecke, T., Lukas, H.L., Sommer, F., “Investigation of the Al-Mg-Si System

by Experiments and Thermodynamic Calculations”, J. Alloys Compd., 247, 31-42 (1997)

(Equi Diagram, Thermodyn., Experimental, Theory, 38)


Seifert, H. J., Lukas, H.L., Aldinger, F., “Experimental Investigation and Thermodynamic


(1998) (Equi. Diagram, Thermodyn, Experimental, Theory, 33)

[1999Mat] Matsuda, K., Naoi, T., Fujii, K., Uetani, Y., Sato, T., Kamio, A., Ikeno, S., “Crystal

Structure of the ´´ Phase in an Al-1.0 mass% Mg2Si-0.4 mass% Si Alloy”, Mater. Sci.

Eng. A, A262, 232-237 (1999) (Crys. Structure, Experimental, 14)

[2000Cay] Cayron, C., Buffat, P.A., “Transmission Electron Microscopy Study of the Phase

(Al-Mg-Si Alloys) and QC Phase (Al-Cu-Mg-Si Alloys) Ordering Mechanism and

Crystallographic Structure”, Acta Mater., 48, 2639-2653 (2000) (Crys. Structure,

Experimental, 38)

[2001Bar] Barabash, O.M., Sulgenko, O.V., Legkaya, T.N., Korzhova, N.P., “Experimental Analysis

and Thermodynamic Calculation of the Structural Regularities in the Fusion Diagram of the

System of Alloys Al-Mg-Si”, J. Phase Equilib., 22(1), 5-11 (2001) (Calculation, Equi.


[2001Goe] Goedecke, T., “Direction of Crystallisations Paths in Ternary As-cast Alloys” (in German),

Z. Metallkd., 92(8), 966-978 (2001) (Equi. Diagram, Experimental, 37)

[2001Li] Li, S.-P., Zhao, S.-X., Pan, M.-X., Zhao, D.-Q., Chen, X.-C., Barabash, O.M., “Eutectic

Reaction and Microstructural Characteristics of Al(Li)-Mg2Si Alloys”, J. Mater. Sci., 36,

1569-1575 (2001) (Equi. Diagram, Experimental, 10)

[2001Mar] Marioara, C.D., Andersen, S.J., Jansen, J., Zandbergen, H.W., “Atomic Model for

GP-Zones in a 6082 Al-Mg-Si System”, Acta Mater., 49, 321-323 (2001) (Crys. Structure,

Metastable, Experimental, 12)

[2001Zha] Zhang, J., Fan, Z., Wang, Y.Q., Zhou, B.L., “Equilibrium Pseudobinary Al-Mg2Si Phase

Diagram”, Mater. Sci. Technol., 17, 494-496 (2001) (Calculation, Equi. Diagram,

Experimental, 17)

[2002Bar] Barabash, O.M., Milman, Yu.V., Korzhova, N.P., Legkaya, T.N., Podrezov, Yu.N.,

“Design of New Cast Aluminium Materials Using Properties of Monovariant Eutectic

Transformation L -Al+Mg2Si”, Mater. Sci. Forum, 396-402, 729-734 (2002) (Equi.

Diagram, Mechan. Prop., 9)

169


MSIT®

Al–Mg–Si

[2002Der] Derlet, P.M., Andersen, S.J., Marioara, C.D., Froseth, A., “A First Principles Study of the

”-Phase in Al-Mg-Si Alloys”, J. Phys.: Condens. Matter, 14, 4011-4024 (2002) (Crys.

Structure, Theory, 19)

[2003Luk1] Lukas, H.L., “Al-Si (Aluminium-Silicon)”, MSIT Binary Evaluation Program, in MSIT



[2003Luk2] Lukas, H.L., Lebrum, N., “Al-Mg (Aluminium-Magnesium)”, MSIT Binary Evaluation



Assessment, 49)

[2003Mar] Marioara, C.D., Andersen, S.J., Jansen, J., Zandbergen, H.W., “The Influence of

Temperature and Storage Time at RT on Nucleation of the ” Phase in a 6082 Al-Mg-Si

Alloy”, Acta Mater., 51, 789-796 (2003) (Crys. Structure, Experimental, 13)


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Mg)

< 650

hP2

P63/mmc

Mg

a = 320.94

c = 521.07

at 25°C [Mas2]

(Al)

< 660.452

cF4

Fm3m

Cu

a = 404.96 pure Al, 25°C [Mas2]

(Si)

< 1414

cF8

Fd3m

C (diamond)

a = 543.06 pure Si, 25°C [Mas2]

~0 at.% Al,

~0 at.% Mg [Mas2]

, Mg2Al3< 452

cF1832

Fd3m

Mg2Al3

a = 2823.9 60-62 at.% Al [V-C2]

, Mg23Al30

410 - 250

hR53

R3

Mg23Al30

a = 1282.54

c = 2174.78

[V-C2, 1998Lia]

, Mg17Al12

< 458

cI58

I43m

Mn

a = 1048.11

a = 1053.05

a = 1057.91

52.58 at.% Mg [L-B]

56.55 at.% Mg [L-B]

60.49 at.% Mg [L-B]

Mg2Si

< 1076

cF12

Fm3m

CaF2

a = 633.8 [Mas2, V-C2]

MgAl3Si6”

mP*

P2/m

?

a = 770 20

b = 670 10

c = 203

= 75 0.5°

[1999Mat]

metastable precipitate in (Al), aged at

150°C

(Mg,Al)5Si6”

mC*

C2/m

?

a = 1516

b = 405

c = 674

= 105.3°

[2001Mar, 2002Der]

metastable precipitate in (Al)

´ hP*

P62m

a = 710

c = 405

[2000Cay] metastable precipitate in (Al)

170


MSIT®

Al–Mg–Si


Reaction T [°C] Type Phase Composition (at. %)

Al Mg Si

L (Al) + Mg2Si 594 e3 L

(Al)

Mg2Si

85.3

97.1

0

10.8

2.7

66.7

3.9

0.2

33.3

L (Al) + (Si) + Mg2Si 557 E1 L

(Al)

(Si)

Mg2Si

81.5

98.0

0

0

5.4

0.70

0

66.7

13.1

1.3

100.0

33.3

L + Mg2Si 462.5 e5 L

Mg2Si

46.1

46.1

0

53.8

53.9

66.7

0.1

0

33.3

L + Mg2Si 451.2 e6 L

Mg2Si

61.0

61.1

0

38.9

38.9

66.7

0.1

0

33.3

L (Al) + + Mg2Si 450 E2 L

(Al)

Mg2Si

64.0

83.4

61.1

0

36.3

16.5

38.9

66.7

0.1

4.0 10-6

0

33.3

L + + Mg2Si 449 E3 L

Mg2Si

57.4

61.1

51.9

0

42.5

38.9

48.1

66.7

0.1

0

0

33.3

L (Mg) + + Mg2Si 435.6 E4 L

(Mg)

Mg2Si

30.9

11.6

39.9

0

69.0

88.4

60.1

66.7

0.1

5.5 10-5

0

33.3

20 40 60

250

500

750

1000

Al Mg 66.67Al 0.00Si 33.33Mg, at.%

Te

mp

era

ture

, °C

(Al) L+(Al)+Mg2Si

(Al)+Mg 2Si

L+Mg2Si

L+(Al)

L

Mg2Si

660.452°C

1076°CFig. 1: Al-Mg-Si.

Section from Al to

Mg2Si

171


MSIT®

Al–Mg–Si

Fig

. 2

: A

l-M

g-S

i. R

eact

ion

sch

eme

Al-

Mg

Al-

Si

A-B

-CA

l-M

g-S

iM

g-S

i

l (

Si)

+ M

g2S

i

94

1e 1

l (

Mg

) +

Mg

2S

i

63

9e 2

l (

Al)

+ (

Si)

57

7e 4

L (

Al)

+ M

g2S

i

59

4e 3

L (

Al)

+ (

Si)

+ M

g2S

i5

57

E1

Lγ,

Mg

2S

i

46

2e 5

Lβ,

Mg

2S

i

45

1e 6

l (

Al)

+ β

45

0.5

e 7

L (

Al)

+ β

+ M

g2S

i4

50

E2

lβ

+ γ

44

9.5

e 8

l (

Mg

) +

γ4

36

e 9

γ +

β ε

ca.4

10

p1

εγ

+ β

ca.2

50

e 10

Lβ

+ γ

+ M

g2S

i4

49

E3

L (

Mg

) +

γ +

Mg

2S

i4

35

.6E

4

γ β

+ ε,

Mg

2S

ica

.41

0D

1

εγ

+ β

, Mg

2S

ica

.25

0D

2

β +

γ +

Mg

2S

i ε +

β +

Mg

2S

iε

+ γ

+ M

g2S

i

(Al)

+ β

+ M

g2S

i

(Al)

+ (

Si)

+ M

g2S

i

(Mg

) +

γ +

Mg

2S

i

γ +

β +

Mg

2S

i

172


MSIT®

Al–Mg–Si

10

90

10

Mg 20.00Al 80.00Si 0.00

Al

Mg 0.00Al 80.00Si 20.00 Data / Grid: at.%

Axes: at.%

640

620

600

580

660

(Al)Mg

2Si

(Si)

e3

E1

600

600

620

620

640

580

580

20

40

60

80

20 40 60 80

20

40

60

80

Mg Al


Axes: at.%

e1

e2

e4

e3

(Mg)

(Al)

(Si)

Mg2Si1000

600

700

800

900

1000

1100

1200

1300°C

600

e9

E4

e8

e7

E1

γ β

E2

E3

Fig. 4: Al-Mg-Si.

Calculated liquidus

surface in the Al

corner

Fig. 3: Al-Mg-Si.

Liquidus surface

173


MSIT®

Al–Mg–Si

20

40

60

80

20 40 60 80

20

40

60

80

Mg Al


Axes: at.%

L+(Mg)+Mg2Si

L+Mg2Si

L

L+(Si)

L+(Si)+Mg2Si

(Al)(Mg)

Mg2Si

(Si)

(Mg)+Mg2Si

20

40

60

80

20 40 60 80

20

40

60

80

Mg Al


Axes: at.%

L+(Mg)+Mg2Si L+Mg

2Si

L

L+(Al)+Mg2Si

(Al)+Mg2Si+(Si)

(Al)(Mg)

Mg2Si

(Si)

(Mg)+Mg2Si

Fig. 5: Al-Mg-Si.


600°C

Fig. 6: Al-Mg-Si.


550°C

174


MSIT®

Al–Mg–Si

(Al) at E1

(Al) at e3

10 2 3 4 5 6 7 8 9 10

0.0

0.2

0.4

0.6

0.8

1.0

1.2

1.4

1.6

650°C

625°C

600°C

575°C

Mg, at.%

Si,

at.%

Al

(Al) at E1

(Al) at e3

10 2 3 4 5 6 7 8 9 10

0.0

0.2

0.4

0.6

0.8

1.0

1.2

1.4

1.6

575°C

Mg, at.%

Si,

at.%

Al

550°C

450°C

500°C

Fig. 7: Al-Mg-Si.

Solidus of the (Al)

phase

Fig. 8: Al-Mg-Si.

Solvus of the (Al)

phase

175


MSIT®

Al–Mg–Si

20 40

250

500

750

1000

1250

Si Mg 50.90Al 49.10Si 0.00Mg, at.%

Te

mp

era

ture

, °C

L+(Al)+Mg2Si

L+Mg2Si

L

L+(Si)

L+Mg2Si+(Si)

(Al)+Mg2Si+(Si)

(Al)+Mg2Si

β+(Al)+Mg2Si

β+γ+Mg2Si

(Si)

557°C

1414°C

L+Mg2Si+γ

γ+ε+Mg2Si

ε+β+Mg2Si

γ+Mg2Si

β+γ+Mg2Si

60 40 20

500

750

1000

Mg 66.67Al 0.00Si 33.33

Mg 0.00Al 53.00Si 47.00Mg, at.%

Te

mp

era

ture

, °C

(Si)+(Al)(Al)+Mg2Si+(Si)

L+Mg2Si+(Si)

L+Mg2Si

L+(Al)+(Si)

L+(Si)

L

Mg2Si

557°C

577°C

1076°C

1029°C

Fig. 9: Al-Mg-Si.


Si to Mg50.9Al49.1

Fig. 10: Al-Mg-Si.


Mg2Si to Al53Si47

176


MSIT®

Al–Mg–Si

10

500

600

700

Mg 0.00Si 9.64Al 90.36

Mg 10.98Si 0.00Al 89.02Mg, at.%

Te

mp

era

ture

, °C

L

L+(Al)

(Si)+(Al)+Mg2Si(Al)+Mg2Si

L+(Al)+Mg2Si

L+(Si)+(Al)

(Si)+(Al)

L+(Al)+Mg2Si

610.3°C

(Al)

596°C

577°C

557°C

10

500

600

700

Mg 0.00Si 14.49Al 85.51

Mg 16.38Si 0.00Al 83.62Mg, at.%

Te

mp

era

ture

, °C

(Si)+(Al)+Mg2Si (Al)+Mg2Si

L

L+(Al)+Mg2SiL+(Si)+(Al)

L+(Al)

L+(Al)

(Si)+(Al)

577°C

L+(Al)

L+(Al)+Mg2Si557°C

Fig. 11: Al-Mg-Si.


mass% Al

Fig. 12: Al-Mg-Si.


mass% Al

177


MSIT®

Al–Mg–Si

10 20

500

600

700

Mg 0.00Si 19.36Al 80.64

Mg 21.72Si 0.00Al 78.28Mg, at.%

Te

mp

era

ture

, °C

(Si)+(Al)+Mg2Si (Al)+Mg2Si

L+(Al)+(Si)

L+(Al)+Mg2Si

L+(Al)+Mg2Si

L+Mg2SiL+(Si)

L

L+(Si)+Mg2Si

(Si)+(Al)

L+(Al)

577°C

Fig. 13: Al-Mg-Si.


mass% Al

178


MSIT®

Al–Mg–Sn

Aluminium – Magnesium – Tin

Lazar Rokhlin, updated by Hans Leo Lukas

Literature Data

[1958Bad] investigated the partial equilibrium diagram Al-Mg2Al3( )-Mg2Sn-Sn using thermal and

microscopic analyses. They studied nine vertical sections, determined the phase equilibria in the solid state,

the nature and temperatures of invariant equilibria and constructed the liquidus surface. [1968Kop,

1969Sem, 1973Sem] investigated the magnesium rich corner of the equilibrium diagram limited by the join

Mg17Al12( ) to Mg2Sn using thermal and microscopic analyses. The authors determined the joint solubility

of Al and Sn in solid Mg at 400 and 200°C and the temperature and compositions of liquid of the ternary

eutectic occurring in this part of the system, constructed four vertical sections and a projection of the

liquidus surface. [1938Hum, 1943Wes] determined the solubility of the Mg2Sn compound in solid Al by

microscopy or combined microscopy and X-ray diffraction, respectively. [1977Ray] reviewed the papers of

[1968Kop, 1969Sem] regarding the Mg corner. A progress report [1950Dow] shows the phase diagram of

the pseudobinary section Mg17Al12( ) to Mg2Sn.

Binary Systems

The binary Al-Sn and Mg-Sn systems are accepted from [Mas2]. Al-Mg is taken from [2003Luk] out of the

MSIT collection of binary systems. It is based on the assessment of [1998Lia].

Solid Phases

No ternary phases have been detected. The stable binary phases are summarized in Table 1. The (Al) solid

solution dissolves up to 16.6 at.% Mg [1998Lia]. The solubility of Sn in solid (Al) is retrograde with a

maximum of nearly 0.026 at.% Sn at 625°C [Mas2]. The maximum solubility of Sn in solid Mg is 3.35 at.%

at the eutectic temperature, 561.2°C [Mas2]. The solubility of Al in solid (Mg) is 11.6 at.% in the binary

Al-Mg system [1998Lia]. Sn and Al decrease somewhat the solubility of each other in solid Mg [1968Kop,

1973Sem]. Solid (Sn) dissolves about 1 at.% Al and nearly no Mg [Mas2]. There is no detectable solubility

of Al in Mg2Sn [1938Hum]. It has not been established if some Sn is soluble in the phases , and of the

Al-Mg binary system [1958Bad, 1969Sem].


Three pseudobinary sections have been established: - Mg2Sn [1950Dow, 1969Sem], - Mg2Sn

[1958Bad] and (Al) - Mg2Sn [1958Bad]. They are illustrated in Figs. 1 to 3. The two versions of the section

- Mg2Sn disagree in the position and temperature of the eutectic, 11.5 mass% Mg2Sn and 450°C

[1969Sem] or 2.5 mass% Mg2Sn and 455°C [1950Dow], respectively. [1969Sem] assumes a large vertical

part (450 to 600°C) in the liquidus line, which postulates zero heat of reversible solution of Mg2Sn in liquid

[1981Goo], which is not very likely along such a large distance as the liquidus becomes very flat not far

above, indicating an appreciable reversible heat of solution there. In constructing the - Mg2Sn section

[1969Sem] assumed certain mutual solubilities of the terminating solid phases, which have not been

confirmed experimentally. After [Mas2, 1938Hum], however, the solubilities of Al as well of Mg in Mg2Sn

are negligible. Therefore in the section in Fig. 1 no solubility is assumed in the terminating solid phases.

The eutectic is accepted from [1950Dow].


Three invariant four-phase equilibria have been reported in the Al-Mg-Sn system: L (Mg)+ +Mg2Sn

[1969Sem], L (Al)+ +Mg2Sn [1958Bad] and L (Sn)+(Al)+Mg2Sn [1958Bad]. In addition, three

invariant three-phase equilibria take place in the pseudobinary systems: L +Mg2Sn [1969Sem],

179


MSIT®

Al–Mg–Sn

L +Mg2Sn [1958Bad] and L (Al)+Mg2Sn [1958Bad]. To complete the reaction scheme, an additional

four-phase equilibrium L + +Mg2Sn must be assumed. The compositions are given in Table 2. Those of

the phases and are taken from the Al-Mg binary system assuming insignificant solubility of Sn in these

phases. The equilibria L +Mg2Sn and L (Al)+ +Mg2Sn are nearly degenerate and the compositions of

the liquid are very near to the Al-Mg binary system [1958Bad]. The temperature of L +Mg2Sn is given

453°C by [1958Bad], but it cannot be higher than that of the maximum L in the binary Al-Mg system.

Thus, 451°C is estimated here for this reaction. Also, the equilibrium L + +Mg2Sn is assumed to be

degenerate. For all the degenerate equilibria, the corresponding concentrations of the binary liquid are

adopted for Table 2. The composition of (Sn) in L (Sn)+(Al)+Mg2Sn is adopted from the accepted Mg-Sn

[Mas2] and Al-Sn [Mas2] binary systems. The three-phase equilibria in solid state of the Al-Mg binary

system are expected to form degenerate four-phase equilibria in the ternary system with Mg2Sn as forth

(inert) phase, since the Sn-solubility in all intermediate Al-Mg phases is assumed to be negligible.

Figure 4 shows the reaction scheme. The concentration range of the triangle - -Mg2Sn has not yet been

investigated, but, as no ternary phases exist, it can easily be interpolated between the known parts.

Liquidus Surface

Figure 5 shows the liquidus surface. The isotherms within the region Mg2Sn-(Sn)-(Al)- are drawn

according to [1958Bad] with small corrections due to the accepted binary Al-Sn and Mg-Sn systems. The

isotherms within the region (Mg) - Mg2Sn - are constructed from the vertical sections given by [1950Dow,

1968Kop], partially from those of [1969Sem] and from the binary Mg-Sn and Al-Mg systems. The very flat

part at 600°C in the field of primary crystallization of (Al) indicates the presence of a metastable miscibility

gap in liquid just below the liquidus surface.

Isothermal Sections

Figure 6 shows the isothermal section at 250°C. It is constructed from the extensions of homogeneity

regions of the phases liquid, (Al), , , (Mg) and (Sn) in the accepted binary systems, assuming negligible

ternary solubilities in the intermediate phases and .


Several temperature-composition sections are given in literature. Besides the pseudobinary system

- Mg2Sn [1969Sem] reported sections at constant Al content of 15 mass% and at constant Sn content of

18 mass%. [1968Kop] constructed a section at constant Mg content of 75 mass%, which is converted into

at.% and redrawn in Fig. 7. The field (Mg)+ is corrected, in the original publication it is drawn too large,

contradicting an isothermal section at the Mg corner given in the same paper. [1958Bad] reported 9 sections

through the Al corner, at high Mg:Sn ratios only partially until the tie line - Mg2Sn. Figure 8 shows the

section at constant atomic ratio Mg:Sn = 1:1.

Thermodynamics

[1983Som] measured the enthalpy of mixing of the liquid along the sections Mg2Sn - Al, Mg50Sn50 - Al,

Mg30Sn70 - Al and Mg50Al50 - Sn at 835°C and along Mg50Sn50 - Al and Mg50Al50 - Sn also at 735°C.

Complete thermodynamic datasets of the Al-Mg and Al-Sn binary systems were assessed in the COST 507

action [1998Ans]. Al-Sn in [1998Ans] contains a typing error, the parameter °LfccAl,Sn must be

45297.84+8.39814 T, not 45297.84 8.39814 T J mol, but was assessed together with the original value

from [1991Din] for °GfccSn - °Gbct

Sn = 4150 - 5.2 T J mol-1. Using the updated value for

°GfccSn - °Gbct

Sn = 5510 - 8.46 T J mol-1 given in [1998Ans] the parameter must be corrected to

°LfccAl,Sn = 43410.66 + 11.76812 T J mol-1. A thermodynamic dataset for Mg-Sn was assessed by

[1993Fri]. For the liquid phase in the Al-Sn and Mg-Sn (above 75 at.% Sn) Gibbs energy datasets are also

given by [1996Heu]. As the ternary solubilities in all solid phases are small, only the liquid phase needs

additional ternary terms for a complete thermodynamic description of the ternary system. A calculation

180


MSIT®

Al–Mg–Sn

using the above mentioned binary descriptions without ternary terms yields fairly good results in the Sn and

Mg rich parts, but fails for the Al rich liquidus surface.


There is information about favour effect of Sn additive on strength properties of Mg base alloys containing

Al at elevated (200-250°C) temperatures [1968Kop].

References

[1938Hum] Hume-Rothery, W., Raynor, G.V., “On the Nature of Intermetallic Compounds of the Type

Mg2Sn”, Philos. Mag., 25, Ser. VII, 335-339 (1938) (Experimental, 3)

[1943Wes] Westlinning, H., Klemm, W., “The Solubility of Mg2Si, Mg2Ge, Mg2Sn and Mg2Pb in

Aluminium” (in German), Z. Electrochem., 49, 198-200 (1943) (Experimental, 3)

[1950Dow] “Liquidus Determinations of Polynary Magnesium Alloys“, Final Status Report No. 15004;

Dow Chemical Company, Off. of Naval Res., Contract No. N9 ONR 85900, 1-20 (1950) (2)


[1958Bad] Badaeva, T.A., Kuznetsova, R.I., “The Structure of Alloys of Aluminium with Magnesium

and Tin”, Tr. Inst. Metall. im. A.A. Baikova, (3), USSR Academy of Science, Moscow,


[1968Kop] Kopetsky, Ch.V., Padezhnova, E.M., Semenova, E.M., “Investigation of the Equilibrium

Diagram of the Mg-Al-Sn in the Magnesium-Rich Region” (in Russian), Izv. Vyss. Uchebn.

Zaved., Tsvetn. Metall., (5), 78-82 (1968) (Equi. Diagram, Experimental, 10)

[1969Sem] Semenova, E.M., “Equilibrium Diagram of the Mg-Al-Sn System in the Magnesium-Rich

Region” (in Russian), Dokl. Akad. Nauk SSSR, 188, 1308-1310 (1969) (Equi. Diagram,

Experimental, 8)

[1973Sem] Semenova, E.M., “Phase Composition and Properties of Alloys of the Mg-Al-Sn System”

(in Russian), Tr. Inst. Metall. im. A.A. Baikova, Moscow, Nauka, 165-168 (1973)

(Experimental, 5)

[1977Ray] Raynor, G.V., “Constitution of Ternary and More Complex Alloys of Magnesium”, Int.

Met. Rev., (5), 65-95 (1977) (Equi. Diagram, Review, 83)

[1981Goo] Goodman, D.A., Cahn, J.W., Bennettt, L.H., “The Centennial of the Gibbs-Konovalov Rule

for Congruent Points”, Bull. Alloy Phase Diagrams, 2, 29-34 (1981) (Equi. Diagram,

Theory, 20)

[1983Som] Sommer, F., Rupf-Bolz, N., Predel, B., “Investigations on the Temperature Dependence of

the Enthalpy of Mixing of Ternary Alloy Melts” (in German), Z. Metallkd., 74, 165-171

(1983) (Experimental, Thermodyn., 15)

[1991Din] Dinsdale, A.T., “SGTE Data for Pure Elements”, Calphad, 15, 317-425 (1991)

(Thermodyn., Assessment)

[1993Fri] Fries, S., Lukas, H.L., “Optimisation of the Mg-Sn System”, J. Chim. Phys., 90, 181-187

(1993) (Equi. Diagram, Thermodyn., Assessment, 32)

[1996Heu] Heuzey, M.-C., Pelton, A.D., “Critical Evaluation and Optimization of the Thermodynamic

Properties of Liquid Tin Solutions”, Metall. Mater. Trans. B, 27B, 810-828 (1996) (Equi.

Diagram, Thermodyn., Assessment, 156)

[1998Ans] Ansara, I., Dinsdale, A.T., Rand, M.H., COST 507, Thermochemical Database for Light

Metal Alloys, Vol. 2, European Communities, Luxembourg, Vol. 2, Al-Mg: 48-54; Al-Sn:

81-83 (1998) (Equi. Diagram, Thermodyn., Assessment, Crys. Structure, 0)




(1998) (Equi. Diagram, Thermodyn., Experimental, Assesssment, 33)

181


MSIT®

Al–Mg–Sn




Assessment, 49)



Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Al)

< 660.452

cF4

Fm3m

Cu

a = 404.88 pure Al at 25°C [Mas2]

(Mg)

< 650

hP2

P63/mmc

Mg

a = 320.89

c = 521.01

pure Mg at 25°C [Mas2]

Sn(r)

231.97 - 13.05

tI4

I41/amd

Sn

a = 583.18

c = 318.18

pure Sn at 25°C [Mas2]

Sn(l)

< 13.05

cF8

Fd3m

C (diamond)

a = 648.92 pure Sn [Mas2]

, Mg2Al3 452

cF1168

Fd3m

Mg2Al3

a = 2816 to 2824 60-62 at.% Al [2003Luk]


[2003Luk]

, Mg17Al12

< 458

cI58

I43m

Mn

a = 1054.38 at 41.4 at.% Al [V-C2]

39.5 to 51.5 at.% Al [2003Luk]

, Mg23Al30

410 - 250

hR159

R3

Mn44Si9

a = 1282.54

c = 2174.78

54.5 to 56.5 at.% Al [2003Luk]

Mg2Sn

< 770

cF12

Fm3m

CaF2

a = 676.5 [V-C]


Al Mg Sn

L (Al) + + Mg2Sn 448 E1 L

(Al)

Mg2Sn

63

83

61

0

37

17

39

66.7

0.4

<0.01

0

33.3

L + + Mg2Sn 447 E2 L

Mg2Sn

57

61

52

0

43

39

48

66.7

0.4

0

0

33.3

182


MSIT®

Al–Mg–Sn

Note: Values in brackets () are estimated.

L (Mg) + + Mg2Sn 428 E3 L

(Mg)

Mg2Sn

31.6

8.3

(41)

0

66.0

91.3

(59)

66.7

2.4

0.4

(0)

33.3

+ , Mg2Sn 410 D1 50.6 49.4 0

+ , Mg2Sn 250 D2 46.4 53.6 0

L (Sn) + (Al) + Mg2Sn 198 E4 L

(Sn)

(Al)

Mg2Sn

0.9

0.7

100

0

8.8

(0)

(0)

66.7

90.3

99.3

<0.01

33.3

L (Al) + Mg2Sn 605 e1 L

(Al)

Mg2Sn

86.2

98

0

9.8

2

66.7

4.0

0.02

33.3

L + Mg2Sn 455 e3 L

Mg2Sn

(45.6)

(46)

(0)

(53.5)

(54)

66.7

0.9

(0)

33.3

L , Mg2Sn 451 e4 L

Mg2Sn

60.8

61.1

0

38.8

38.9

66.7

0.4

(0)

33.3


Al Mg Sn

10 20 30

400

500

600

700

800

Mg 54.00Al 46.00Sn 0.00

Mg 66.70Al 0.00Sn 33.30Sn, at.%

Te

mp

era

ture

, °C

γ+Mg2Sn

L+Mg2Sn

L

L+γ

455°C458°C

770.5°C

e3

Fig. 1: Al-Mg-Sn.

The pseudobinary

system - Mg2Sn

183


MSIT®

Al–Mg–Sn

10 20 30

400

500

600

700

800

Mg 39.20Al 60.80Sn 0.00

Mg 66.70Al 0.00Sn 33.30Sn, at.%

Te

mp

era

ture

, °C

β+Mg2Sn

L+Mg2Sn

L

451°C

770.5°CFig. 2: Al-Mg-Sn.

The pseudobinary

system - Mg2Sn

20 40 60 80

400

500

600

700

800

Mg 66.70Al 0.00Sn 33.30

Al

Al, at.%

Te

mp

era

ture

, °C

Mg2Sn+(Al)

L+Mg2Sn

L

605°C

660.452°C

770.5°C

e1

L+(Al)

Fig. 3: Al-Mg-Sn.

The pseudobinary

system Mg2Sn - Al

184


MSIT®

Al–Mg–Sn

Fig

. 4:

A

l-M

g-S

n.

Rea

ctio

n s

chem

e

Al-

Mg

Mg

-Sn

A-B

-CA

l-M

g-S

nA

l-S

n

L

(Al)

+ M

g2S

n

60

5e 1

l (A

l) +

β4

50

e 5

l β

+ γ

44

9e 6

l (

Mg)

+ γ

43

6e 7

εβ

+ γ

25

0e 8

l(M

g)

+ M

g2S

n

56

1e 2

l (S

n)

+ (

Al)

22

8e 9

L (

Al)

+ β

+ M

g2S

n 4

48

E1

l(S

n)

+ M

g2S

n

20

3e 1

0

β +

γε

41

0p

1

L

β +

Mg

2S

n

45

1e 4

Lβ

+ γ

+ M

g2S

n4

47

E2

β +

γε,

Mg

2S

n4

10

D1

L (

Mg)

+ γ

+ M

g2S

n4

28

E3

(Mg

)+γ+

Mg

Sn

Lγ

+ M

g2S

n

45

5e 3

L(S

n)

+ (

Al)

+ M

g2S

n1

98

E4

(Sn

)+(A

l)+

Mg

2S

n

ε β

+ γ

, M

g2S

n2

50

D2

β+γ+

Mg

2S

n

β+ε+

Mg

2S

nε+

γ+M

g2S

n

β+γ+

Mg

2S

n

(Al)

+β+

Mg

2S

n

185


MSIT®

Al–Mg–Sn

20

40

60

80

20 40 60 80

20

40

60

80

Mg Al

Sn Data / Grid: at.%

Axes: at.%

Mg2Sn

(Al)+L+Mg2Sn

(Al)βγ(Mg)

(Mg)+γ+Mg2Sn

L

β+γ+Mg2 Sn

(Al)+β+Mg2 Sn

20

40

60

80

20 40 60 80

20

40

60

80

Mg Al

Sn Data / Grid: at.%

Axes: at.%e

9

e10

E4

(Al)

e1

Mg2Sn

(Mg)

e2

e7 E

3

350

400

450

500

550

600610

620

650

700

750

650600550

600 500

e6

e5

E1

(Sn)

βγ E2

e3 e

4

Fig. 6: Al-Mg-Sn.


250°C

Fig. 5: Al-Mg-Sn.

Liquidus surface

186


MSIT®

Al–Mg–Sn

10 20

400

500

600

Mg 93.61Al 0.00Sn 6.39

Mg 76.90Al 23.10Sn 0.00Al, at.%

Te

mp

era

ture

, °C

L

L+(Mg)

L+(Mg)+Mg2Sn

(Mg)+Mg2Sn

(Mg)+γ+Mg2Sn

L+(Mg)+γ

(Mg)+γ

428°C

10 20 30 40

100

200

300

400

500

600

700

Al Mg 50.00Al 0.00Sn 50.00Sn, at.%

Te

mp

era

ture

, °C

(Al)+(Sn)+Mg2Sn

L+(Al)+Mg2Sn

L

L+(Al)

L+Mg2Sn

L+(Sn)+Mg2Sn

198°C

604°C 31.55%Mg635°C

203.5°C

660.452°C

Fig. 7: Al-Mg-Sn.

Temperature-

concentration section

at 75 mass% Mg

Fig. 8: Al-Mg-Sn.

Temperature-

concentration section

at the atomic ratio

Mg:Sn = 1:1

187


MSIT®

Al–Mg–Ti

Aluminium – Magnesium – Titanium

Frederick H. Hayes, updated by Andy Watson and Tatyana Dobatkina

Literature Data

There is general agreement that the solubilities of both Mg in Ti aluminides and Ti in liquid and solid Al-Mg

alloys are very small [1954Eis, 1968Var, 1971Dil, 1973Kol, 1984Rus, 1989Ker1]. The observation of

[1954Eis] that the solubility of Ti in liquid Mg, 0.04 mass% Ti at 750°C, is sharply decreased by Al

additions was later confirmed by [1971Dil]. [1971Dil] found that additions of Al to Mg-Ti alloys reduce the

solubility of Ti in the liquid phase as follows: at 720°C from 0.028 to 0.003 by 0.5 mass% Al, at 800°C from

0.042 to 0.004 by 1 mass% Al, at 900°C from 0.08 to 0.02 by 2 mass% Al. [1954Eis] reports that the

solubility of Ti in Al-Mg alloys decreases with increasing Mg content at 750°C to become vanishingly small

at 90 mass% Mg. At 12.5 mass% Mg [1954Eis] gives the solubility of Ti as 0.122 mass% at 750°C in

agreement with the later work [1973Kol]. [1968Var] studied intermetallic phases in Al-Mg alloys

containing traces of transition metals; TiAl3 was observed to be present in Al-Mg alloys containing up to

38 at.% Mg and 1 at.% Ti. The liquidus contours given by [1973Kol] from 700 to 850°C for the Al rich

corner containing up to 12 mass% Mg and 0.8 mass% Ti are given in Fig. 1. [1984Rus] constructed the

427°C isothermal section for the entire composition range using kinetic data from vapor-diffusion and

powder-sintering experiments. Electron-microprobe analysis, electron microscopy and X-ray

phase-analysis techniques were used to determine phase compositions. Only binary Al-Ti and Al-Mg

intermetallic phases were found. In contrast, the 487°C isothermal section of [1989Ker1] contains a ternary

phase in the Al-rich corner; reference to a ternary Al-Mg-Ti phase is also made in the report of [1948Fel].

[1989Ker1] gives the single-phase composition range of the ternary phase as 78.5 to 80.5 at.% Al, 11.4 to

12.8 at.% Mg and 8.7 to 9.1 at.% Ti. In other respects the sections of [1984Rus, 1989Ker1] are in good

agreement. No further information on solid-liquid equilibria or invariant reactions is available for this

system.

Binary Systems

The binary Al-Mg system is accepted from [2003Luk], and the Al-Ti system from [1995Hay]. Both differ

from those given by [Mas2] in the solid-solid equilibria. The binary Mg-Ti system is accepted from [Mas2].

Solid Phases

A ternary phase is reported by [1948Fel, 1987Ker, 1989Ker1, 1989Ker2] but was not mentioned by

[1968Var, 1984Rus]. [1987Ker] gave the lattice parameter and suggested it to be of the Cr2Mg3Al18 type.

The crystallographic data of all solid phases are given in Table 1.

Isothermal Sections

Figure 2 gives the 427°C isothermal section based on [1984Rus, 1989Ker1]; it is consistent with the

accepted binary phase diagrams.

Thermodynamics

From a thermodynamic analysis of the Gibbs energy of formation of binary phases in equilibrium with

Ti2Mg3Al18 [1989Ker2] concluded that the Gibbs energy of formation of the ternary compound was

approximately -15 kJ·mol-1.

188


MSIT®

Al–Mg–Ti

Miscellaneous

[1996Set] used X-ray electron probe analysis to define titanium concentrations throughout dendritic cells

of a cast aluminium solid solution of Al-Mg-Ti. [1996Set] indicated that Mg levels within the limits of its

solubility in aluminium, have no significant effect on titanium intercrystalline segregation.

References

[1948Fel] Feldman, W., Schrader, A., Seemann, J., “Structure of Primary Aluminium and Aluminium

Alloys” (in German), FIAT Rev. German Sci., 1939-1946, Non-Ferrous Metallurgy, 1,

153-155 (1948) (Review, 21)

[1954Eis] Eisenreich, H., Putter, H., “Magnesium-Titanium Ternary Systems” (in German), Metall, 8,

624-625 (1954) (Experimental, 0)

[1968Var] Varich, N.I., Litvin, B. N., “Structure of Phases in the Aluminium-Magnesium System

Containing Transition Metals” (in Russian), Izv. Akad. Nauk SSSR. Met., (6), 179-182


[1971Dil] Dilov, V.V., Sergeev, V.V., “Effect of Some Elements on Ti Solubility in Liquid Mg” (in

Russian), Tr. Vses. N.–I. Proekt. Inst. Alyum. Magn. El., (79), 100-113 (1971)

(Experimental, 4)

[1973Kol] Kolpachev, A.A., Medvedeva, N.D., Samoilova, Yu.A., Titova, I.A., “Solubility of Ti in

Al-Mg Alloys” (in Russian), Tekhnol. Legk. Splavov, (8), 15-17, (1973) (Experimental,

Equi. Diagram, #, 3)

[1984Rus] Rusnyak, V.D., Dunaev, S.F., Slyusarenko, E.M., Sokolskii, S.V., Sokolovskaya, E.M.,

“Study of Phase Equilibria in the Aluminium-Magnesium-Titanium System” (in Russian),

Deposited Doc., VINITI, 2189-89, Moscow, 15 pp. (1984) (Experimental, Equi.

Diagram, *, 11)

[1987Ker] Kerimov, K.M., Dunaev, S.F., Slyusarenko, E.M., “Investigation of the Structure of Ternary

Phases in Al-Mg-Ti, Al-Mg-V and Al-Mg-Cr Systems”, J. Less-Common Met., 133,

297-302 (1987) (Experimental, Crys. Structure, 9)

[1989Ker1] Kerimov, K.M., Dunaev, S.F., Slyusarenko, E.M., “Study of the Phase Diagrams of the

Systems: Aluminium-Magnesium-(Titanium, Zirconium, Hafnium)” (in Russian), Vestn.

Mosk. Univ., Ser. 2: Khim., 30(2), 156-161 (1989) (Experimental, Equi. Diagram, *, 8)

[1989Ker2] Kerimov, K.M., Dunaev, S.F., “The M2Mg3AL18 Phase in Al-Mg Transition Metal

Systems”, J. Less-Common Met., 153, 267-273 (1989) (Thermodyn., 10)


Z. Metallkd., 81, 389-396 (1990) (Equi. Diagram, Experimental, 33)

[1995Hay] Hayes, F.H., “The Al-Ti-V (Aluminium-Titanium-Vanadium) System”, J. Phase Equilib.,

16(2), 163-176 (1995) (Equi. Diagram, Review)

[1995Bra] Braun, J., Ellner, M., Predel, B., “Experimental Investigations of the Structure and Stability

of the TiAl Phase”, Z. Metallkd., 86(12), 870-876 (1995) (Experimental, Crys. Structure)

[1996Set] Setiukov, O.A., Fridlyander, I.N., “Peculiarities of Ti Dendritic Segregation in Aluminium

Alloys”, Mater. Sci. Forum, 217-222, 195-200 (1996) (Experimental, 2)




(1998) (Equi. Diagram, Thermodyn., Experimental, Theory, *, 33)




Assessment, 49)

189


MSIT®

Al–Mg–Ti


Phases/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Al)

< 660.45

cF4

Fm3m

Cu

a = 404.96 pure Al at 25°C [Mas2]

dissolves ~17.0 at.% Mg

(Mg)

< 650

hP2

P63/mmc

Mg

a = 320.94

c = 521.07


dissolves ~0.12 at.% Ti,

~12 at.% Al

,( Ti) (h)

1670 - ~865

cI2

Im3m

W

a = 330.65 pure Ti at 25°C [Mas2]

dissolves 45 at.%Al,

~2.4 at.%Mg

,( Ti) (r)

1490

hP2

P63/mmc

Mg

a = 295.06

c = 468.35

pure Ti at 25°C [Mas2]

dissolves 51.5 at.% Al,

~1.6 % Mg

, Mg2Al3 452

cF1168

Fd3m

Mg2Al3

a = 2816 to 2824 60-62 at.% Al [2003Luk]


[2003Luk]

, Mg17Al12

< 458

cI58

I43m

Mn

a = 1054.38 at 41.4 at.% Al [V-C2]

39.5 to 51.5 at.% Al [2003Luk]

, Mg23Al30

410 - 250

hR159

R3

Mn44Si9

a = 1282.54

c = 2174.78

54.5 to 56.5 at.% Al [2003Luk]

TiAl3< 1387

tI8

I4/mmm

TiAl3

a = 384.88

c = 859.82

[1990Sch]

TiAl2(r)

< 1216

tI24

I41/amd

HfGa2

a = 396.7

c = 2429.68

[1990Sch]

TiAl

< 1460

tP4

P4/mmm

AuCu

a = 398.69

c = 405.39

a = 398.8

c = 408.2

38.5 to 52 at.% Ti [1990Sch]

at 38.5 at.% Ti, 1000°C [1990Sch]

at 45 at.% Ti, 20°C [1995Bra]

2, Ti3Al

1180

hP8

P63/mmc

Ni3Sn

a = 580.6

c = 465.5

a = 574.6

c = 462.4

at 78 at.% Ti [L-B]

at 62 at.% Ti [L-B]

* , Ti2Mg3Al18 cF184

Fd3m

Cr2Mg3Al18

a = 1477 [1987Ker]

190


MSIT®

Al–Mg–Ti

850

800

7507

00

Mg

Ti

Al

15.000.00

85.00

Mg

Ti

Al

0.001.50

98.50

Al

10 5

Mg, at.%

Al,

at.%

90

95

i 1 i id i h Al i h [19 3 l]

850

800

7507

00

Mg

Ti

Al

15.000.00

85.00

Mg

Ti

Al

0.001.50

98.50

Al

10 5

Mg, at.%

Al,

at.

%

90

95

Fig. 1: Al-Mg-Ti.

Liquidus contours in

the Al-rich corner

[1973Kol]

20

40

60

80

20 40 60 80

20

40

60

80

Mg Ti


Axes: at.%

(αTi)

α2

TiAl

TiAl2

TiAl3

(Al)

Mg2Al

3

(Mg)

+Mg 17

Al 12

+TiAl 3

(Mg)

τ

(αTi)+α2+(Mg)

α2+TiAl+(Mg)

TiAl 2+TiA

l 3+(M

g)

TiAl3+τ+(Al)

Mg2Al

3+TiAl

3+τ

Mg2Al

3+τ+(Al)

Mg17

Al12

Mg2Al

3+Mg

17Al

12+TiAl

3

Fig. 2: Al-Mg-Ti.


427°C

191


MSIT®

Al–Mg–Zn

Aluminium – Magnesium – Zinc

Dmitriy Petrov, Andy Watson, Joachim Gröbner, Peter Rogl, Jean-Claude Tedenac, Marina Bulanova,

Volodymyr Turkevich, updated by Hans Leo Lukas

Literature Data

The Al-Mg-Zn system has a relatively complex equilibrium diagram. The first investigation of the entire

system was carried out by [1913Ege]. By thermal analysis and metallography he determined the liquidus

surface and some invariant equilibria. He also detected the first ternary phase, having a large homogeneity

range between the binary phases Mg17Al12 and MgZn2. Many investigations were done in the following

time by thermal analysis and metallography [1926San, 1936Ham1, 1936Ham2, 1936Koe1, 1936Koe2,

1936Koe3, 1940Ura, 1943But, 1945But, 1945Mik, 1949Sal, 1961Cla, 1962Ale, 1985Kuz1, 1985Kuz2], by

electric conductivity measurements [1936Fin], by X-ray diffraction [1935Lav, 1936Fin, 1936Ham1,

1936Ham2, 1936Rie, 1957Ber, 1961Cla, 1985Kuz1, 1985Kuz2, 1995Tak, 1997Don, 2000Lee, 2000Sun]

or by TEM [1997Don, 2001Bou1].

Thermodynamic datasets of the ternary system were assessed by [1997Lia, 1998Lia2].

Binary Systems

The Al-Mg system is adopted from [2003Luk] which is based on the thermodynamic assessment of the

COST 507 action [1998Ans] modified in the central part by [1998Lia1] based on new experimental data.

The Al-Zn and Mg-Zn binary systems are accepted from the COST 507 action [1998Ans].

Solid Phases

Two ternary phases are known since long time [1913Ege, 1961Cla]. Recently a stable ternary

quasicrystalline phase and another stable crystalline phase were detected [1995Tak]. The first ternary phase

( 1 in Table 1) was called Mg7Zn6Al3 by [1913Ege], Mg30Zn25Al20 by [1929Ish, 1930Ish]. [1935Lav]

determined a cubic unit cell with a = 1416 pm at a composition Mg3Zn3Al2. [1936Rie] measured the lattice

parameter at different compositions and found 1429 to 1471 pm along the line from Mg3Zn3Al2 to Mg2Al3and 1429 to 1460 pm along the line from Mg3Zn3Al2 to Mg17Al12. [1952Ber, 1957Ber] determined the

complete crystal structure using single crystal X-ray diffraction. The ideal formula is Mg32(Zn,Al)49 with

162 atoms per unit cell. [2000Sun] and independently [2000Lee] refined the crystal structure stating the

position 2(a) of space group Im3 to be empty, where [1952Ber, 1957Ber] assumed Al occupation. For seven

of the eight remaining positions of [1952Ber, 1957Ber] both papers agree in having 100% Mg occupation

at the sites 16(f), 24(g), 12(e) and Al,Zn mixed occupation on a 48(h) and two different 24(g) positions.

[2000Sun], like [1952Ber, 1957Ber] assume also 100% Mg occupation at another 12(e) position, where

[2000Lee] assume mixed occupation of about 8 atoms Mg+Zn and about 4 sites to be empty. This agrees

with the experimental homogeneity range having significant extension also perpendicular to Al-Zn

exchange. The phase is closely related to the quasicrystalline phases and characterized as 1/1 crystalline

approximant of these phases.

[1959Cla, 1961Cla] established the existence of another ternary phase at 40Mg-40Zn-20Al (mass%)

(54.9Mg-20.4Zn-24.7Al (at.%)) and designated this phase ( in Table 1 and in [2001Bou1]). The phase

is in equilibrium with (Mg) at 335, 204°C and probably at room temperature. [1961Wri] also discovered the

presence of this ternary phase. From transmission electron diffraction data [1997Don] derived an

orthorhombic unit cell of this ternary phase and successfully indexed the X-ray powder diagram (except for

two reflections) with lattice parameters a = 897.9 pm, b = 1698.8 pm and c = 1934 pm. [1997Don] gave the

solubility range of the this phase as (53 to 55)Mg-(18 to 29)Al-(17 to 28)Zn (at.%), which are in good

agreement with [1961Cla]. [2001Bou1] derived a model for the crystal structure of this phase from electron

diffraction patterns obtained in a transmission electron microscope: Space group Pbcm, Mg84(Al,Zn)68

192


MSIT®

Al–Mg–Zn

(4 times Mg21(Al,Zn)17) on 23 different Wyckoff-positions, 13 of them occupied by Mg and 10 ones

occupied by Al+Zn. The lattice parameters were adopted from [1997Don].

Icosahedral quasicrystalline phases can be prepared by ultra rapid quenching of melts of compositions

inside the stability range of the 1 phase. [1995Tak] found the quasicrystalline phase to be stable even up to

the melting temperature of about 380°C, but only in a very narrow composition range at Mg44Zn41Al15. A

sample of this composition is still in the quasicrystalline state after 5h annealing at 360°C, whereas a sample

of composition Mg45Zn40Al15 shows small amounts of (Mg) precipitates in a quasicrystalline matrix after

1 h annealing at 360°C. The sample 1 mol% opposite, Mg43Zn42Al15, after the same annealing conditions

transforms into a cubic phase with a = 2291 pm, which also is stable in a very narrow composition range

and which is characterized as the 2/1 crystalline approximant of the icosahedral quasicrystalline phase. This

phase was confirmed and characterized as package of 8 Bergman atom clusters in the cubic unit cell

[2002Hir, 2002Sug], the same clusters, two of which are in the unit cell of 1. The space group is Pa3.

Lattice parameter (2310 pm) and composition (Mg46Zn37Al17) differ slightly from those given by

[1995Tak]. [2001Bou2] confirmed the appearance of a quasicrystalline phase in Mg cast alloys containing

4 mass% Al and 8 mass% Zn. [2000Bok] reported a reversible transformation from a quasicrystalline phase

to the 2/1 approximant in Mg44Zn41Al15 at 340°C and a subsequent irreversible decomposition of the 2/1

approximant above 420°C.

[1961Cla] indicated additionally the probability of the existence of a further ternary phase of undetermined

composition near the Mg-Zn boundary in the region of the phases MgZn, Mg2Zn3 and MgZn2. This may be

taken as an earlier hint to the two phases identified by [1995Tak].

Metastable precipitates are formed during low temperature annealing of supersaturated (Al) solid solutions,

which were formed by quenching from temperatures with higher solubility of Mg and Zn in (Al).

Guinier-Preston zones are formed at and slightly above room temperature. At somewhat higher

temperatures a metastable phase, ´, is formed coherently in the Al matrix. Several models for its structure

were proposed. [2001Wol] calculated the Gibbs energy of three models from first principles and proposed

the model of [1974Aul] to be the most likely one. At even higher temperatures (above about 200°C) also

the stable phases and 1 may precipitate.

The crystallographic data of the stable solid phases are summarized in Table 1.


Data are available in the literature on the following sections Al-MgZn2 [1926San], MgZn2-Mg2Al3[1936Ham1], MgZn2- 1 [1936Koe1], Al- 1 [1936Koe1] and 1-Mg17Al12 [1936Koe2]. Essentially, there

are three reports on sections from MgZn2 through the ternary 1 phase to the Al-Mg side, in two of these

works, [1913Ege, 1936Koe2], the section ends at the Al-Mg side in Mg17Al12 whereas that of [1936Ham1]

ends at Mg2Al3. In [1913Ege], only one of the phases, Mg17Al12, is shown in the Al-Mg system. It should

be noted that the composition of the 1 phase which is associated with the pseudobinary equilibria at p1

(L+MgZn2 1), e1 (L (Al)+ 1), e3 (L + 1) and e2 (L + 1) varies considerably. This variation in

composition reflects the very wide homogeneity range of the 1 phase.

The section Mg2Al3-MgZn2 ( -MgZn2) is nearly a true quasibinary section with the peritectic point p1, the

eutectic point e2 and the compositions of 1 associated with these two invariant equilibria lying very near

to the vertical plane of section [1936Ham1]. Fig. 1 shows a calculation of this section using the dataset of

[1998Lia2]. The (MgZn2) homogeneity range is very narrow along constant Mg content and thus extends

outside the plane of section. Therefore the corner of the three-phase field L+ + 1 in Fig. 1 markedly

deviates from true quasibinary behavior. However, if the plane of section near this area would be shifted a

few tenth of at.% towards more Mg, it would look truly quasibinary.

The section from Mg17Al12 to MgZn2 ( to MgZn2), although regarded as a quasibinary section by

[1913Ege], was shown by [1936Koe1, 1936Koe2] not to be a quasibinary. The tie line joining

Mg17Al12-e3- 1 at 450°C deviates significantly from the section Mg17Al12-MgZn2 [1936Koe3]. The

calculation of this section using the dataset of [1998Lia2] is shown in Fig. 2.

A special problem arises in connection with the Al-MgZn2 section. The section was first established in

[1913Ege] with the eutectic point at 63Zn-11.6Mg-25.4Al (mass%) and at 473°C. [1923San] observed the

193


MSIT®

Al–Mg–Zn

same section when studying the Al-Mg-Zn system for the presence of quasibinary sections. In an alloy with

approximately 84 mol% MgZn2 on the join Al-MgZn2 the authors observed a clearly pronounced purely

eutectic structure between Al and MgZn2 and thus confirmed the quasibinary nature of the system.

[1926San] gave a picture of this quasibinary system with the eutectic point at approximately

63Zn-12Mg-25Al (mass%) and 475°C. [1936Ham1], when summarizing the results of his extensive work

on the Al-Mg-Zn system, indicated that the Al-MgZn2 system can be treated as an independent binary

system. [1936Koe1] established by microscopic studies the presence of the four-phase transition reaction

L+ 1 (Al)+MgZn2. The quadrangle representing the four-phase region, however, is very nearly a triangle

with the liquid phase lying close to the (Al)-MgZn2 join. According to [1986Luk] this type of reaction may

be considered as a ternary degenerate equilibrium combining the four-phase equilibrium

L+ 1+(Al)+MgZn2 with the maximum of the three-phase equilibrium L+ 1+(Al). [1943Sie] carried out

metallographic and X-ray diffraction investigations and determined the solubility of MgZn2 in (Al). He

found a maximum solubility of 17 mass% MgZn2 at 475°C, thus rejecting the result of 30% MgZn2 obtained

by [1938Sal]. [1979Sti] gave results of measurements relating to the reaction L+ 1 (Al)+MgZn2. The

four-phase region was represented in this work, as in [1936Koe1], by a triangle with the liquid phase apex

on the side Al-MgZn2. The temperature of the reaction is the same, i.e. 475°C. (A misprint occurs in the text

of [1936Koe1] and the temperature is given as 375°C). The composition of (Al) in the reaction is shifted

substantially towards the Al side in disagreement with the results given in [1936Koe1]. In all the

publications mentioned, [1926San], [1938Sal, 1943Sie], except [1913Ege], the temperature of the

quasibinary reaction coincides with the temperature of the four-phase equilibrium L+ 1 (Al)+MgZn2

[1936Koe1, 1979Sti]. The calculation with the dataset of [1998Lia2] shows the liquid phase 2.1 at.%

outside the triangle 1-(Al)-MgZn2, therefore, contrary to [1993Pet], here this reaction is labelled as U type,

not as degenerate. The calculated section Al-MgZn2 is shown in Fig. 3. Except near the MgZn2 phase the

approximation of it as quasibinary is quite good. At lower temperatures down to about 410°C the

three-phase equilibrium 1+(Al)+MgZn2 crosses the section as its (Al) corner is slightly shifted from the

plane of section to the more Zn rich side. The ternary solid solution of MgZn2 extends outside the section

into the more Mg-rich side, therefore near MgZn2 the section cuts the three-phase equilibrium

MgZn2+(Al)+(Al,Zn), going from the (Al)-miscibility gap of the binary Al-Zn system to MgZn2.


Three reaction schemes can be found in the literature. They relate respectively to the Zn corner [1936Koe1],

the Mg corner [1936Koe3], and the Mg corner for compositions greater than 50 at.% Mg [1961Cla]. The

first and the third schemes served as the basis of the scheme of invariant equilibria in [1986Des] however,

the second scheme cannot be accepted because the later discovered ternary phase, [1961Cla], is missing

there. This phase over an appreciable range of composition prevents equilibrium between the ternary 1

phase and (Mg). [1988Ito] presented invariant equilibria based on isothermal sections in the Al-rich corner.

Since these equilibria used old binary data they were rejected in this assessment.

The thermodynamic assessments of two different groups, H. Liang et al. [1997Lia] and P. Liang et al.

[1998Lia2], enable the calculation of a complete comprehensive set of invariant equilibria. In both

assessments the stable quasicrystalline phase and their 2/1 approximant ( 2) detected by [1995Tak] are

missing. Below in the section “Thermodynamics” an attempt is described to incorporate these two phases

into the dataset of [1998Lia2]. The invariant equilibria calculated from this updated dataset are listed in

Table 2 and the corresponding reaction scheme is given in Figs. 4a and 4b. As the stability ranges of the

quasicrystalline phase and 2 are known approximative only the reaction scheme has to be taken as partially

tentative, indicated by dashed boxes for the reactions.

Liquidus Surface

[1913Ege] was the first to construct the liquidus surface for the entire field of the Al-Mg-Zn system. Besides

(Al), (Mg) and (Zn), he detected only three more solid phases in the system: (Mg17Al12), MgZn2, and a

ternary phase Mg7Zn6Al3. [1936Ham1] used thermal, microscopic and X-ray analysis and gave a

completely different ternary homogeneity range than [1913Ege], shown in the form of a relatively narrow

194


MSIT®

Al–Mg–Zn

band displaced generally towards the Al-Mg system from the field given by [1913Ege]. In [1936Ham2], the

MgZn2 field occupies an excessively wide region in the system. The ternary phase was confirmed in general

in later works [1934Fus, 1936Koe1, 1936Koe2, 1936Koe3, 1936Fin, 1937Fin, 1940Ura, 1943But,

1943Lit1, 1943Lit2, 1945Mik, 1949Sal], but the ternary homogeneity range has been widened noticeably

compared with [1913Ege, 1936Ham1, 1936Ham2] at the expense of the MgZn2 field.

[1959Cla, 1961Cla] established the existence of another ternary phase, , and connected a new invariant

transformation with this phase: L+ + 1 at 393°C; the nature of this transformation was determined by

[1961Cla] as peritectic. The author claimed this as the first report of a ternary peritectic reaction in literature

on alloy constitution.

[1962Ale] investigated in detail the liquidus of the system in the Zn corner within 10 mass% Al and 8

mass% Mg. [1961Yue] used the method of zone melting to determine eutectic compositions in complex

metallic systems and established the existence of a ternary eutectic of the composition 50Zn-3Al-47Mg

(mass%) at 338°C in the Al-Mg-Zn system. The authors repeated this composition in another work,

[1970Yue]. It should be noted, however, that the schematic ternary eutectic equilibrium L (Mg)+ +MgZn

suggested by [1970Yue] was earlier given by [1961Cla] as L (Mg)+MgZn+ 1.

[1973Wil], in a brief but rather detailed review, gave the liquidus surface of the whole Al-Mg-Zn system.

The data of the reviews by [1971Mon, 1976Mon] generally agree with those of [1973Wil]. The liquidus

surface of [1973Wil] has been adopted virtually without changes in [1986Des]. [1985Kuz1] investigated by

thermal, microscopic and X-ray methods the Al-Zn region of the system and corrected to a certain extent

the liquidus region of (Mg2Al3). It is narrowed somewhat at the Al-Mg side and spreads to 15 mass% Zn

as compared with the results of [1973Wil]. [1986Kuz] studied the portion of the liquidus surface in the Al

corner of the system up to 30 mass% Mg and 30 mass% Zn and repeated to a certain extent the results of

[1985Kuz1].

The liquidus surface calculated by [1998Lia2] differs in some details from that constructed in the review of

[1993Pet]: the line of double saturation of the liquid with 1 and MgZn2 is markedly curved towards the

MgZn2 phase, the equilibria between liquid, MgZn2, 1, Mg2Zn3 and MgZn are interchanged and due to the

updated Al-Mg system [1998Lia1] the Al-Mg side is simplified. Taking into account the stability of the

quasicrystalline phase q (Mg44Zn41Al15) as described below in section “Thermodynamics” also a small

field of primary solidification of q appears. The liquidus surface calculated from this updated dataset is

shown in Fig. 5.

Solidus and Solvus Surfaces

[1936Koe1, 1936Koe2, 1936Koe3] presented the solidus polytherm of the Al corner. The solidus polytherm

in [1952Han] is displaced sharply to substantially lower concentrations of Mg along its whole length as

compared with the polytherm in [1936Koe3]. [1945But] constructed a series of solidus isotherms in the

temperature interval 630 to 500°C. [1973Wil] gave a generalized solidus surface of the Al corner of the

system, which was adopted in the review of [1993Pet]. The solidus isotherms of [1979Sti, 1985Kuz1,

1985Kuz2] as well as the calculated ones of [1998Lia2] virtually coincide with those of [1973Wil]. Figure

6 shows the solidus isotherms of the (Al) phase calculated from the dataset of [1998Lia2].

The solubility of Zn and Mg in (Al) was studied by a number of authors [1933Boc, 1945But, 1947Str,

1955Zam, 1961Sal, 1971Mon]. Their results are compared and presented in the solvus surface of

[1973Wil], which was adopted in the review of [1993Pet]. The solvus of the (Al) phase, calculated from the

dataset of [1998Lia2] (Fig. 7) agrees very well with that of [1993Pet].

The dataset of [1998Lia2] enables also a calculation of solidus and solvus isotherms of the (Mg) phase. They

are shown in Fig. 8. To allow different scaling of the composition axes of Al and Zn, in this figure not Gibbs

triangular, but Cartesian rectangular coordinates are chosen. The existence of equilibria between the (Mg)

solid solution and the stable quasicrystalline phase is supported by [1995Tak] and [2001Bou2]. The

extension of the fields of these equilibria, however, must be taken as tentative.

195


MSIT®

Al–Mg–Zn

Isothermal Sections

[1936Fin, 1937Fin] used microscopic and X-ray methods and measurements of electrical resistance to

construct a series of isothermal sections in the Al corner in the range up to 20 mass% Mg and 20 mass% Zn

over the temperature interval 400 to 200°C, and an isothermal section adjoining the Al-Zn side of the system

in the range up to 40 mass% Mg. Similar work was carried out by [1943Lit1, 1943Lit2] who constructed a

series of isothermal sections in the Al corner in the range up to 12 mass% Mg and 12 mass% Zn. [1971Mon]

gave isothermal sections in the Al corner in the range up to 16 mass% Mg and 16 mass% Zn and compared

it with the results of [1936Fin, 1937Fin]. Of the highest interest is the work of [1961Cla] who constructed

the isothermal section of the whole ternary system at 335°C using metallographic and X-ray methods and

thermal analysis on heating and cooling. This temperature is only 3°C below the lowest solidus temperature

of the system. The ternary phase was discovered in the same study. [1961Cla] studied the nature of the

invariant equilibria associated with this ternary phase and constructed another two schematic isothermal

sections at 394 and 374°C which enabled him to establish the equilibrium L+ + 1 at 393°C. A number

of other equilibria involving the phase were investigated so that a reaction scheme could be constructed,

as mentioned in the section “Invariant Equilibria”. [1973Wil] in his review reproduced completely the

isothermal section at 335°C taken from [1961Cla]. [1986Des] proposed some corrections for the isothermal

section of [1961Cla]. [1998Lia2] investigated 34 samples annealed at 335°C mainly to fix the maximum

ternary solubilities of the binary phases. These alloys were prepared by induction melting in a graphite

crucible placed inside a copper crucible. The samples were re-melted twice to ensure homogeneity, and

precautions were taken to reduce Zn and Mg loss. The phase compositions as measured by EPMA are

largely in agreement with previous work [1986Des] apart from the Al solubilities in MgZn2, Mg2Zn3 and

MgZn, which are of the same magnitude, whereas in the review of [1993Pet] a much lower Al solubility

was assumed for Mg2Zn3. The 335°C isothermal section in Fig. 9 is based on the calculated one of

[1998Lia2] with corrections to include the phases q and 2.

Thermodynamics

Kim et al. [1997Kim] determined the enthalpies of mixing of the liquid phase over the temperature range

610-660°C using high temperature calorimetry. From the results thermodynamic functions of the liquid

were assessed using the associated solution model.

A thermodynamic description of the solidus and liquidus surfaces in the Al-rich corner has been evaluated

by [1990Kuz]. Thermodynamic datasets covering the whole ternary system were assessed by two groups,

[1997Lia, 1998Lia2]. The first group used more simplified models for the description of the phases. Ternary

solubilities in binary phases are considered only for liquid, (Al), (Mg), (Zn) and the Laves phase MgZn2,

is described as a stoichiometric phase and 1 as a line compound with constant Mg content of 39.5 at.%,

Mg32(Zn,Al)49. Apart from the restrictions due to these simplifications both assessments do not deviate

severely. [1998Lia2] used own measurements to adjust the Al solubilities in Mg2Zn11, MgZn2, Mg2Zn3 and

MgZn, as well as the Zn solubilities in the , and phases of the Al-Mg system. These phases, except

MgZn2 and are described as line compounds Mgx(Zn,Al)y, is also described as a line compound, 1 is

done by the model Mg26(Mg,Al)6(Mg,Zn,Al)48Al1, reproducing fairly well the experimental homogeneity

range.

In the present assessment an attempt was made to incorporate the stable quasicrystal (q) and the 2/1

approximant ( 2) into the dataset of [1998Lia2], both described as stoichiometric phases. The Gibbs

energies per mole of atoms are expressed as:

Gq - 0.15 GAlfcc - 0.44 GMg

hcp - 0.41 GZnhcp = -7760. -1. T (J mol-1)

and

G 2 - 0.15 GAlfcc - 0.43 GMg

hcp - 0.42 GZnhcp = -7900. -1. T (J mol-1)

Also the description of the phase was updated to satisfy the formula Mg21(Zn,Al)17 given by [2001Bou1]:

GMg:Al - 17 GAlfcc - 21 GMg

hcp = 38 (-1380. -1.5 T) (J mol-1)

GMg:Zn - 21 GMghcp - 17 GZn

hcp = 38 (-7150. +1.9 T) (J mol-1)

LMg:Zn,Al = 38 (-2100 +1. T) (J mol-1)

196


MSIT®

Al–Mg–Zn

This updated dataset was used to calculate all figures and tables of this assessment. With respect to the

phases q and 2 it has to be considered as tentative.

Miscellaneous

[1968Tha] investigated, by electron diffraction, the process of decomposition on long term ageing of

quenched Al-6Zn-2Mg (mass%) alloy and discovered for the first time the formation of an intermediate '

phase with a hexagonal cell and lattice parameters a = 267 pm and c = 490 pm. [1976Aul] investigated, by

X-ray diffraction, single crystals of the alloy Al-4.0Zn-2.9Mg (mass%) after quenching from 490°C

followed by ageing at 155°C for 24 h and also demonstrated the formation of an intermediate phase called

' with a hexagonal cell and the lattice parameters a = 1388 pm and c = 2752 pm. Long holding at 230°C

(for 200 d) of an Al-Mg-Zn alloy led to the formation of the exclusively equilibrium 1 phase. The influence

of predeformation on the precipitation and resulting mechanical properties were studied by [1997Des] for

the alloy Al-6.1Zn-2.35Mg (mass%).

The formation and decomposition of metastable quasicrystalline phases were investigated in various rapidly

solidified alloys [1986Cas, 1986Raj, 1986Sas, 1988Cha, 2000Bok, 2000Miz, 2000Tak, 2001Bou2]. The

resulting internal melting of several Al-Mg-Zn alloys was studied by rapid quenching in a salt bath

[1994Dro].

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MSIT®

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198


MSIT®

Al–Mg–Zn

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199


MSIT®

Al–Mg–Zn

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Assessment, 53)

[1994Dro] Droenen, P.-E., Ryum, N., “Local Melting in Al-Mg-Zn Alloys”, Metall. Mater. Trans. A,

25A, 521-530 (1994) (Equi. Diagram, Experimental, 23)

[1995Tak] Takeuchi, T., Mizutani, U., “Electronic Structure, Electron Transport Properties and

Relative Stability of Icosahedral Quasicrystals and their 1/1 and 2/1 Approximants in the

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Precipitation in an Al-Zn-Mg Alloy”, Z. Metallkd., 88(8), 601-606 (1997) (Experimental,

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[1997Don] Donnadieu, P., Quivy, A., Tarfa, T., Ochin, P., Dezellus, A., Harmelin, M.G., Liang, P.,

Lukas, H.L., Seifert, H.J., Aldinger, F., Effenberg, G., “On the Crystal Structure and

Solubility Range of the Ternary Phase in the Mg-Al-Zn System”, Z. Metallkd., 88(12),


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System”, Met. Mater. Trans., 28A, 1725-1734 (1997) (Equi. Diagram, Thermodyn.,

Assessment, 75)

[1997Kim] Kim, Y.B., Sommer, F., Predel, B., “Calorimetric Investigation of Liquid

Aluminium-Magnesium-Zinc Alloys”, J. Alloys Compd., 247, 43-51 (1997) (Thermodyn.,

Theory, Experimental, 20)

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Light Metal Alloys”, Ansara, I., Dinsdale, A.T., Rand, M.H. (Eds.), European

Communities, Luxembourg, Vol. 2, 329-332 (1998) (Equi. Diagram, Thermodyn.,

Assessment, 0)

[1998Lia1] Liang, P., Su, H.-L., Donnadieu, P., Harmelin, M.G., Quivy, A., Ochin, P., Effenberg, G.,



(1998) (Equi. Diagram, Thermodyn., Experimental, Assesssment, *, #, 33)

[1998Lia2] Liang, P., Tarfa, T., Robinson, J. A., Wagner, S., Ochin, P., Harmelin, M.G., Seifert, H.J.,

Lukas, H.L., Aldinger, F., “Experimental Investigation and Thermodynamic Calculation of

the Al-Mg-Zn System”, Thermochim. Acta, 314, 87-110 (1998) (Equi. Diagram,

Thermodyn., Experimental, Assessment, *, #, 69)

[2000Bok] Bokhonov, B.B., Ivanov, E.Y., Tolochko, B.P., Sharaphutdinov, M.P., “In Situ Study of

Structural Transformations of Mg44Al15Zn41 Quasicrystals under Heating”, Mater. Sci.

Eng. A, A278, 236-241 (2000) (Crys. Structure, Experimental, 8)

[2000Lee] Lee, C.-S., Miller, G.J., “Where are the Elements in Complex Aluminides? An

Experimental and Theoretical Investigation of the Quasicrystalline Approximant,

Mg2-y(ZnxAl1-x)3+y”, J. Am. Chem. Soc., 122, 4937-4947 (2000) (Crys. Structure,

Experimental, Theory, 72)

[2000Miz] Mizutani, U., “Electron Transport Mechanismin the Pseudogap System: Quasicrystals,

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(Crys. Structure, Experimental, Theory, 23)

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(Al,Zn)49Mg32-Type Phases by Single-Crystal X-Ray Diffraction”, Mater. Sci. Eng. A,

A294-296, 327-330 (2000) (Crys. Structure, Experimental, 9)

200


MSIT®

Al–Mg–Zn

[2000Tak] Takeuchi, T., Mizuno, T., Banno, E., Mizutani, U., “Magic Number of Electron

Concentration in the Icosahedral Cluster of AlxMg40X60-x (X = Zn, Cu, Ag, and Pd) 1/1

Cubic Approximants”, Mater. Sci. Eng. A, A294-296, 522-526 (2000) (Crys. Structure,


[2001Bou1] Bourgeois, L., Muddle, B.C., Nie, J.F., “The Crystal Structure of the Equilibrium Phase

in Mg-Zn-Al Casting Alloys”, Acta Mater., 49, 2701-2711 (2001) (Crys. Structure,


[2001Bou2] Bourgeois, L., Mendis, C.L., Muddle, B.C., Nie, J.F., “Characterization of Quasicrystalline

Primary Intermetallic Particles in Mg-8wt% Zn-4 wt% Al Casting Alloy”, Philos. Mag.

Lett., 81, 709-718 (2001) (Crys. Structure, Experimental, 33)

[2001Wol] Wolverton, C., “Crystal Structure and Stability of Complex Precipitate Phases in

Al-Cu-Mg-(Si) and Al-Zn-Mg Alloys”, Acta Mater., 49, 3129-3142 (2001) (Crys.

Structure, Theory, 64)

[2002Hir] Hiraga, K., Sugiyama, K., Ishi, Y., “Arrangement of Atomic Clusters in a 2/1 Cubic

Approximant in the Al-Zn-Mg Alloy System”, Philos. Mag. Lett., 82, 341-347 (2002)


[2002Sug] Sugiyama, K., Sun, W., Hiraga, K., “Crystal Structure of a Cubic Al17Zn37Mg46; a 2/1

Rational Approximant Structure for the Al-Zn-Mg Icosahedral Phase”, J. Alloys Compd.,

342, 139-142 (2002) (Crys. Structure, Experimental, 8)




Assessment, 49)


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

, (Al)

< 660.45

cF4

Fm3m

Cu

a = 404.96 at 25°C [Mas2]

(Mg)

< 650

hP2

P63/mmc

Mg

a = 320.94

c = 521.07

at 25°C [Mas2]

(Zn)

< 419.58

hP2

P63/mmc

Mg

a = 266.50

c = 494.70

at 25°C [Mas2]

, Mg2Al3 452

cF1168

Fd3m

Mg2Al3

a = 2816 to 2824 60-62 at.% Al [2003Luk]


[2003Luk]

, Mg17Al12

< 458

cI58

I43m

Mn

a = 1054.38 at 41.4 at.% Al [V-C2]

39.5 to 51.5 at.% Al [2003Luk]

, Mg23Al30

410 - 250

hR159

R3

Mn44Si9

a = 1282.54

c = 2174.78

54.5 to 56.5 at.% Al [2003Luk]

, Mg51Zn20

342 - 325

oI142

Immm

Mg51Zn20

a = 1408.3

b = 1448.6

c = 1402.5

[Mas2], called Mg7Zn3, lattice

parameters

for Mg72Zn28 [V-C]

201


MSIT®

Al–Mg–Zn


MgZn

< 347

- - [Mas2]

, Mg2Zn3

< 416

mC110

B2/m

- [Mas2]

, MgZn2

< 590

hP12

P63/mmc

MgZn2

a = 522.1

c = 856.7

[Mas2, V-C]

, Mg2Zn11

< 381

cP39

Pm3

Mg2Zn11

a = 855.2 [Mas2, V-C]

* 1, Mg32(Zn,Al)48 cI160

Im3

Mg32(Zn,Al)48

a = 1413 - 1471 [1957Ber] (gave cI162) [2000Lee,

2000Sun]

1/1 approximant of icosahedral phase

* , Mg21(Zn,Al)17 oP152

Pbcm

Mg21(Zn,Al)17

a = 897.9

b = 1698.8

c = 1934

[2001Bou1]

lattice parameters from [1997Don]

* 2, Mg43Zn42Al15 or

Mg46Zn37Al17

cP640 (?)

Pa3

a = 2291

a = 2310

[1995Tak]

[2002Hir, 2002Sug]

2/1 approximant of icosahedral phase

* q, Mg44Zn41Al15 quasicrystalline,

icosahedral

[1995Tak] stable quasicrystalline phase


Al Mg Zn

L + MgZn2 1 530 p1 (max) L

MgZn2

1

21.8

7.0

18.2

37.0

33.5

36.2

41.2

59.5

45.6

L (Al) + 1 480 e1 (max) L

(Al)

1

51.1

90.2

23.2

21.4

4.2

33.8

27.5

5.6

43.0

L + 1 (Al) + MgZn2 476 U1 L

1

(Al)

MgZn2

5.1

19.5

89.6

7.5

18.7

33.0

2.9

33.3

36.2

47.5

7.5

59.2

L + 1 451 e2 (max) L

1

55.4

56.6

48.1

39.2

38.9

41.0

5.4

4.5

10.9

L + 1 451 e3 (max) L

1

46.6

46.5

46.0

47.0

49.0

42.4

6.4

4.5

11.6

L + + 1 449 E1 L 52.0

59.5

50.6

43.1

38.9

47.6

4.9

1.6

1.8

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

202


MSIT®

Al–Mg–Zn

L (Al) + + 1 447 E2 L

(Al)

1

60.1

85.8

56.4

47.7

34.3

13.2

38.9

40.1

5.6

1.0

4.7

12.2

L + MgZn2 + 1 Mg2Zn3 434 P1 L

MgZn2

1

Mg2Zn3

5.6

3.6

10.7

5.0

60.6

33.8

40.3

40.0

33.8

62.6

49.0

55.0

L + + 1 388 P2 L

1

17.4

34.1

30.1

12.0

65.8

58.1

44.1

56.0

16.8

7.8

25.8

32.0

L + 1 q 380 p5 (max) L

q

7.2

15.8

15.0

66.7

41.8

44.0

26.1

42.4

41.0

q + 1 2 377 p6 (max) q

1

2

15.0

15.0

15.0

44.0

41.7

43.0

41.0

43.3

42.0

L + (Mg) + 368 U2 L

(Mg)

14.0

6.0

32.1

25.5

69.8

92.2

60.1

55.2

16.2

1.8

7.8

19.3

L + 1 q + 362 U3 L

1

8.7

19.7

19.9

68.4

42.7

55.3

22.9

37.6

24.8

L + 1 Mg2Zn3+ q 362 U4 L

1

Mg2Zn3

4.7

12.1

5.2

67.9

41.5

40.0

27.4

46.4

54.8

q + 1 Mg2Zn3 + 2 356 U5 1

Mg2Zn3

12.3

5.4

41.3

40.0

46.4

54.6

L + MgZn2 Mg2Zn11 + (Al) 355 U6 L

MgZn2

Mg2Zn11

(Al)

10.9

2.0

3.1

45.0

7.6

33.0

15.4

0.3

81.5

65.0

81.5

54.7

L + q + Mg2Zn3 MgZn 353 P3 L

Mg2Zn3MgZn

4.4

5.1

5.1

68.7

40.0

48.0

26.9

54.9

46.9

L + (Mg) + q 345 U7 L

(Mg)

6.5

18.3

3.0

70.1

55.3

94.6

23.4

26.4

2.4

L (Al,Zn) + Mg2Zn11 + (Zn) 344 E3 L

(Al,Zn)

Mg2Zn11

(Zn)

8.7

34.8

2.4

2.5

6.0

0.1

15.4

0.2

85.3

65.1

82.2

97.3


Al Mg Zn

203


MSIT®

Al–Mg–Zn

L + Mg51Zn20 (Mg)+MgZn 340 U8 L

(Mg)

MgZn

0.1

0.05

0.1

71.0

97.0

48.0

28.9

2.95

51.9

L q + (Mg) + MgZn 339 E4 L

(Mg)

MgZn

4.4

2.0

5.3

70.1

95.3

48.0

25.5

2.7

46.7

MgZn2 + (Al,Zn) (Al) + Mg2Zn11 331 U9 MgZn2

(Al,Zn)

(Al)

Mg2Zn11

1.8

52.3

76.7

2.9

33.1

0.2

0.1

15.4

65.1

47.5

23.2

81.7

(Al,Zn) (Al)+(Zn), Mg2Zn11 277 D1 (Al,Zn)

(Al)

(Zn)

Mg2Zn11

41.0

85.8

1.6

2.1

0.06

0.03

0.1

15.4

58.94

14.17

98.3

82.5

Mg2Zn3 + 1 MgZn2 + 2 214 U10 Mg2Zn3

1

MgZn2

2

6.1

15.8

3.8

15.0

40.0

39.0

33.4

43.0

53.9

45.2

62.8

42.0

q + (Mg) + MgZn 84 U11 q

(Mg)

MgZn

15.0

0.4

19.3

4.4

44.0

99.3

55.3

48.0

41.0

0.3

25.4

47.6


Al Mg Zn

10 20 30 40 50 60

200

300

400

500

600

Mg 38.00Zn 0.00Al 62.00

Mg 33.33Zn 66.67Al 0.00Zn, at.%

Te

mp

era

ture

, °C

LL+η

η

τ1+η

L+τ1+ηL+τ1

τ1+(Al) τ1τ1+β+(Al)

L+(Al)+β

(Al)+β

L+τ1+(Al)

L+τ1+β

η+ζτ1+η+ζ

η+q

τ1+η+q

η+ζ+q

590°C

L+β

Fig. 1: Al-Mg-Zn.

Calculated section

from (Mg38.5Al62.8)

to MgZn2

204


MSIT®

Al–Mg–Zn

10 20 30 40 50 60

200

300

400

500

600

Mg 58.62Zn 0.00Al 41.38

Mg 33.33Zn 66.67Al 0.00Zn, at.%

Te

mp

era

ture

, °C

590°C

L

L+τ1

L+η

L+τ1+η

ητ1+η

ζ+η

L+η+ζ

L+τ1+q τ1

τ1+τ2+ζ

τ1+η+ζ

τ1+ζτ1+τ2

τ1+φ+q

τ1+τ2+q

τ1+φ

γ+φ+τ1

γ+φ

γ+τ1

L+γ+τ1

γ

L+γ

η+ζ+τ2

τ1+τ2+ητ2+η

Fig. 2: Al-Mg-Zn.

Calculated section

from (Mg17Al12) to

MgZn2

20 40 60 80

200

300

400

500

600

700

Mg 33.33Zn 66.67Al 0.00

Al

Al, at.%

Te

mp

era

ture

, °C

L

(Al)+η

(Al)

L+(Al)

(Al)+τ1+η(Al)+τ1

L+(Al)+η

L+η

η

(Al)+θ+ηθ+η

(Al,Zn)+(Al)+η(Al,Zn)+θ+η

L+η+θ

L+(Al)+τ1

660°C

590°C

Fig. 3: Al-Mg-Zn.

Calculated section

from MgZn2 to Al

205


MSIT®

Al–Mg–Zn

Fig

. 4a:

Al-

Mg-Z

n.

Rea

ctio

n s

chem

e, p

art

1

Al-

Mg

Al-

Zn

A-B

-C

l (

Al)

+ β

45

0e 4

l (

Al)

+ (

Zn)

38

0e 8

L +

ητ 1

53

0p

1

Al-

Mg-Z

n

L +

τ1

(A

l) +

η4

76

U1

Mg-Z

n

l +

ηζ

41

6p

2

lβ

+ γ

44

9e 6

l (

Mg)

+ γ

43

6e 7

β +

γε

41

0p

3

l +

ηθ

38

1p

4

l (

Zn

) +

θ3

60

e 9

L (

Al)

+ τ

1

48

0e 1

Lβ

+ τ 1

45

1e 2

Lτ 1

+ γ

45

1e 3

Lβ

+ γ

45

0e 5

Lτ 1

+ β

+ γ

44

9E

1

L (

Al)

+ τ

1 +

β4

47

E2

L+

τ 1 +

γφ

38

8P

2

L +

τ1 +

ηζ

43

4P

1

L +

γ (

Mg)

+ φ

36

8U

2

L +

τ1

φ +

q3

62

U3

L +

τ1

q +

ζ3

62

U4

q +

τ1

τ 2 +

ζ3

56

U5

L +

τ1

q

38

0p

5

q +

τ1

τ 2

37

7p

6

L+

(Al)

+η

L+

γ+φ

L+

τ 1+

φ

τ 1+

q+

ζ

L+

q+

ζτ 1

+τ 2

+ζ

τ 1+

η+ζ

q+

τ 1+

τ 2

U1

0U

6

U6

E3

q+

τ 2+

ζ

τ 1+

φ+q

L+

q+

φ

L+

(Mg

)+φ

γ+(M

g)+

φ

τ 1+

γ+φ

(Al)

+τ 1

+β

τ 1+

β+γ

τ 1+

(Al)

+η

E3

L+

τ 1+

ζ

e 13U

7

P3

U1

0

U7

206


MSIT®

Al–Mg–Zn

Fig

. 4b

:A

l-M

g-Z

n.

Rea

ctio

n s

chem

e, p

art

2

Al-

Mg

Al-

Zn

A-B

-C

δ (

Mg)

+ M

gZ

n

32

5e 1

1

Al-

Mg-Z

n

L +

η (

Al)

+ θ

35

5U

6

Mg-Z

n

l +

ζ M

gZ

n

34

7p

7

(Al,

Zn)

(A

l)+

(Zn)

27

7e 1

2

l +

(Mg)

δ341.1

p8

l M

gZ

n +

δ3

41

e 10

L +

q +

ζ M

gZ

n3

53

P3

L +

φ (

Mg)

+ q

34

5U

7

Lθ

+ (

Zn)

+ (

Al,

Zn)

34

4E

3

L +

δ (

Mg)

+

MgZ

n3

40

U8

L q

+ (

Mg)

+ M

gZ

n3

39

E4

η +

(Al,

Zn)

θ +

(A

l)3

31

U9

(Al,

Zn

) (A

l) +

(Z

n),

θ2

77

D1

τ 1 +

ζη

+ τ 2

21

4U

10

q +

(M

g)

φ +

Mg

Zn

84

U1

1

εβ

+ γ

25

0e 1

3

e 9e 8

τ 1+

η+ζ

U3

p3

U2

U4

U5

p4

L+

(Al)

+η

L+

φ+q

L+

(Mg

)+φ

L+

q+

ζ

τ 1+

τ 2+

ζ

τ 1+

τ 2+

ζ

L+

(Al,

Zn

)+θ η+

θ+(A

l,Z

n)

L+

q+

MgZ

n

q+

ζ+M

gZ

n

L+

q+

(Mg)

q+

φ+(M

g)

(Zn

)+(A

l,Z

n)+

θ

L+

(Mg

)+M

gZ

n

q+

(Mg

)+M

gZ

n

(Al,

Zn

)+(A

l)+

θη+

(Al)

+θ

(Al)

+(Z

n)+

θ

τ 1+

η+τ 2

η+ζ+

τ 2

q+

φ+M

gZ

n(M

g)+

φ+M

gZ

n

η+(A

l,Z

n)+

(Al)

P1

U1

207


MSIT®

Al–Mg–Zn

20

40

60

80

20 40 60 80

20

40

60

80

Mg Zn


Axes: at.%

(Zn)

e8

U6

E3

p4

e9

(Al)

U1e

3

η

p1

τ1

(Mg)

U2

P2

γ

p8

p7 p2

q

ζ

U7

P1

e4

e6 β

e2

E2

E1

600

500

450

500

600

500

450

e7

θ

E4

P3

550

500

e1

550

400400

φ U3

U4

p5

U8

MgZne

10

10

10

90

Mg 15.00Zn 0.00Al 85.00

Mg 0.00Zn 15.00Al 85.00


Axes: at.%

E2

e1

U1

640°C

620

600

580

560

540

520

500

480

460

Fig. 5: Al-Mg-Zn.

Liquidus surface,

calculated using the

data of [1998Lia2]

Fig. 6: Al-Mg-Zn.

Calculated solidus

isotherms of the (Al)

phase

208


MSIT®

Al–Mg–Zn

10

10

90

Mg 15.00Zn 0.00Al 85.00

Mg 0.00Zn 15.00Al 85.00


Axes: at.%

E2

e1

U1

τ1

β

η

460°C

440

420

400380360340

0 2 4 6 8 1210

Al, at.%

Mg

0

1

2

3

Zn

,a

t.%

solidus isotherms

solvus isotherms

univariant equilibria

�

q

MgZn

350

300

250

200

150

600

550

500

450

400°C

U8

E4

U7

U2

e7

Fig. 7: Al-Mg-Zn.

Calculated solvus

isotherms of the (Al)

phase

Fig. 8: Al-Mg-Zn.

Calculated solidus

and solvus isotherms

of the (Mg) phase

209


MSIT®

Al–Mg–Zn

20

40

60

80

20 40 60 80

20

40

60

80

Mg Zn


Axes: at.%(Al)

(Al,Zn)

(Zn)

ηζMgZn

q(Mg)

γ

ε

β

τ1φ

θ

τ2

(Mg)+MgZn+q

(Mg)+q+φ(Mg)

+φ+γ

(Mg)+φ

(Mg)+MgZn

θ+(Zn)

+(Al,Zn)η+θ+(Al,Zn)

η+(Al)+(Al,Zn)

(Al)+η

(Al)+τ1

γ+τ1

φ+τ1

γ+φ+τ1

Fig. 9: Al-Mg-Zn.


335°C, calculated

from the data of

[1998Lia2]

210


MSIT®

Al–Mg–Zr

Aluminium – Magnesium – Zirconium

Natalia R. Bochvar, updated by Marina V. Bulanova

Literature Data

[1958Ich, 1959Ich, 1968Bab] investigated the solubility of Zr and Al in liquid Mg by the chemical analysis

of alloys quenched in water from 700 to 800°C. The solubility of zirconium decrease sharply on the addition

of even minor quantities of aluminium. It was also found that the ZrAl3 compound is in equilibrium with

liquid magnesium. [1960Sch] found, by very careful chemical analysis of the Mg rich liquid separated at

740°C from the solid equilibrium phases and by X-ray investigation of those, that ZrAl3 did not appear up

to at least 0.0506 mass% Al and 0.0372 mass% Zr. Below that Al content the precipitates were identified

as ZrAl2, Zr2Al3, Zr4Al3 and a more Zr-rich phase. [1964Cro] found by metallographic and X-ray

inspection that ZrAl2 and Mg17Al12 were present in cast Mg rich alloys with 3 to 10 mass% Al and 0.01 to

1.6 mass% Zr. ZrAl2 was present even in the sample with 0.01 mass% Zr, indicating a very low joint solid

solubility in (Mg). [1969Dri1, 1969Dri2] studied the aluminium corner of the system by metallographic and

differential thermal analyses and reported the existence of a ternary phase (ZryMgyAlx) in equilibrium with

(Al) though the composition and crystal structure of this phase have not been determined. In following

investigations of [1989Ker, 1992Fri], however, the existence of the ternary compound in the Al rich corner

of the system was not confirmed. [1989Ker] studied the interaction of Al-Mg alloys with Zr by the

diffusion-couple technique and constructed the isothermal section at 400°C in the whole concentration

range; transition zones were analyzed by the electron microprobe technique. Using metallographic analysis,

X-ray diffraction and electron microprobe techniques, [1992Fri] investigated the alloys annealed at 400°C

with the constant Mg content of 6 mass%. According to [1989Ker, 1992Fri], only the phases which belong

to the corresponding binaries exist in the system. [1977Asa1, 1977Asa2] investigated the magnesium corner

of the system by measurements of electrical conductivity, thermo-emf and calculations based on general

thermodynamic relations. They supposed the existence of two additional ternary compounds in the system,

“Zr3Mg8Al9” and “ZrMg6Al3”. The vertical sections, however, have been constructed by the authors with

violations of the phase rule. Their interpretation of the obtained results is considered unreliable and these

works will not be further discussed.

Binary Systems

The following binary systems are accepted: Al-Mg [2003Luk], Al-Zr [Mas2, 1992Per] in the Zr-rich part

and Mg-Zr [Mas2].

Solid Phases

The solid phases are given in Table 1. No ternary phase is accepted.


It may be supposed that in the Al rich corner of the system, a ternary eutectic (Al)+ , Mg2Al3+ZrAl3 exists

at 450°C. This supposition is based on the absence of a ternary phase [1989Ker, 1992Fri] and on the absence

of any arrests on the thermocurves except that at 450°C [1969Dri2].

Isothermal Sections

Figure 1 shows the isothermal section at 400°C [1989Ker].

Miscellaneous

Figure 2 shows the solubility limit at 740°C in the magnesium corner from 0.001 to 0.05 mass% Al

[1960Sch]. The most Al poor precipitate could not be identified. With increasing Al content Zr4Al3, Zr2Al3

211


MSIT®

Al–Mg–Zr

and ZrAl2 were found in this careful investigation [1960Sch]. The solubility data of 0.4 to 0.04 mass% Zr

and 0.05 to 0.015 mass% Al are much lower compared with earlier studies [1958Ich, 1959Ich]. Data from

those earlier studies are given in Fig. 3 as isotherms of solubility at 700 and 800°C in the magnesium corner

of the system according to [1958Ich, 1959Ich], where only ZrAl3 is mentioned as the precipitated phase.

[1984Kai] reported the possibility of obtaining a zirconium-supersaturated solid solution of aluminium in

an Al-2 Mg-0.16Zr (at.%) alloy by quenching at a high rate.

[1991Lav] studied the alloy Mg-8.4Al-0.2Zr (mass%) (Mg-7.64Al-0.04Zr (at.%)) after it was spray

atomized and the deposited. The research methods were SEM, EDAX, X-ray diffraction. The deposited

residual contained (Mg) and eutectic (Mg)+ , Mg17Al12. The deposit contained (Al) as well, that proves

nonequilibrium process.

References

[1958Ich] Ichikawa, R., “Solubility of Zr in Mg and its Alloys in the Liquid State. II. Alloys

Containing Al, Fe, Mn and Si” (in Japanese), Nagoya Kogo Daigaku Gakuho, 10, 197-203

(1958) (Equi. Diagram, Experimental, #, 8)

[1959Ich] Ichikawa, R., “Al, Fe, Mn and Si as Impurities in Mg-Zr Alloys. I. Solubility of Zr in Molten

Mg. II. Intermetallic Compounds Formed by Impurities and Zr” (in Japanese), Nippon

Kinzoku Gakkaishi, 23, 192-194 (1959) (Equi. Diagram, Experimental, #, 5)

[1960Sch] Schneider, A., Stendel, J., “Precipitation of Intermetallic Phases from a Liquid Solvent

Metal” (in German), Z. Anorg. Allg. Chem., 303, 227-246 (1960) (Experimental, 28)

[1964Cro] Crosby, R.L., Higley, L.W., “Intermetallic Compounds in Mg-Rich Mg-Al-Zr Alloys”, U.S.

Bur. Mines, Rep. Invest., 1-23 (1964) (Equi. Diagram, Experimental, 7)

[1968Bab] Babkin, V.M., “Solubility of Zr in Molten Mg and ML5 Alloy” (in Russian), Metalloved.

Term. Obrab. Met., 3, 61-64 (1968) (Equi. Diagram, Experimental, #, 4)

[1969Dri1] Drits, M.E., Kadaner, E.S., Kuzmina, V.I., “Phase Diagram of the Al-Mg-Zr System in

Al-Rich Region”, Russ. Metall., translated from Izv. Akad. Nauk SSSR, Met., 5, 170-173


[1969Dri2] Drits, M.E., Kadaner, E.S., Kuzmina, V.I., “Interaction of Components in Ternary Systems”

(in Russian), in a Collection of Papers “Aluminium Alloys”, Metallurgiya, Moscow, 6,

146-149 (1969) (Equi. Diagram, Experimental, #, 2)

[1977Asa1] Asanovich, V.Ya., Sryvalin, I.T., Korpachev, V.G., “An Electrometric Study of the

Aluminium-Magnesium Zirconium System” (in Russian), Nauchn. Tr. Kuban. Univ., 3,

42-48 (1977) (Equi. Diagram, Experimental, #, 6)

[1977Asa2] Asanovich, V.Ya., “Phase Diagram of the Aluminium- Magnesium- Zirconium System”,

Russ. Metall., (4), 169-171 (1977), translated from Izv. Akad. Nauk SSSR, Met., (4), 208-210


[1984Kai] Kaibyshev, O.A., Valiev, R.Z., Tsenev, N.K., “Influence of the Grain Boundary State on the

Superplastic Flow”, Sov. Phys. -Dokl., 29, 752-754 (1984), translated from Dokl. Akad.

Nauk SSSR, 278, 93-97 (1984) (Experimental, 12)

[1989Ker] Kerimov, K.M., Dunaev, S.F., Slusarenko, E.M., “Investigations on Phase Equilibria in

Aluminium - Magnesium- (Titanium, Zirconium, Hafnium) Systems” (in Russian), Vestn.

Mosk. Univ. Ser. 2: Khim., 30, 156-161 (1989) (Experimental, Equi. Diagram, 8)

[1991Lav] Lavernia, E.J., Baram, J., Gutierrez, E., “Precipitation and Excess Solid Solubility in

Mg-Al-Zr and Mg-Zn-Zr Processed by Spray Atomization and Deposition”, Mat. Sci.

Eng. A, A132, 119-133 (1991) (Experimental, Crys. Structure)

[1992Fri] Fridman, A.S., Dobatkina, T.V., Muratova, E.V., “Section of Isothermic Tetrahedron of the

Al-Rich Portion of the Al-Mg-Sc-Zr System at 500°C” (in Russian), Izv. Akad. Nauk SSSR,

Met., (1), 234-236 (1992) (Equi. Diagram, Experimental, 4)

[1992Per] Peruzzi, A., “Reinvestigation of the Zr-Rich End of the Zr-Al Equilibrium Phase Diagram”,

J. Nucl. Mater., 186, 89-99 (1992) (Equi. Diagram, Experimental, 17)

212


MSIT®

Al–Mg–Zr




Assessment, 49)


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Al)

< 660.45

cF4

Fm3m

Cu

a = 404.96 pure Al at 25°C [Mas2]

(Mg)

< 650

hP2

P63/mmc

Mg

a = 320.94

c = 521.05


( Zr)

1855 - 863

cI2

Im3m

W

a = 360.99 [Mas2]

( Zr)

< 863

hP2

P63/mmc

Mg

a = 323.12

c = 514.77

pure Zr at 25°C [Mas2]

, Mg2Al3 452

cF1168

Fd3m

Mg2Al3

a = 2816 to 2824 60-62 at.% Al [2003Luk]


[2003Luk]

, Mg17Al12

< 458

cI58

I43m

Mn

a = 1054.38 at 41.4 at.% Al [V-C2]

39.5 to 51.5 at.% Al [2003Luk]

, Mg23Al30

410 - 250

hR159

R3

Mn44Si9

a = 1282.54

c = 2174.78

54.5 to 56.5 at.% Al [2003Luk]

ZrAl3< 1580

tI16

I4/mmm

ZrAl3

a = 401.4

c = 1732.0

[V-C, Mas2]

ZrAl2< 1645

hP12

P63/mmc

Zn2Mg

a = 528.24

c = 874.82

[V-C, Mas2]

Zr2Al3< 1595

oF40

Fdd2

Zr2Al3

a = 960.1

b = 1390.6

c = 557.4

[V-C, Mas2]

ZrAl

< 1275

oC8

Cmcm

CrB

a = 335.3

b = 1086.6

c = 426.6

[V-C, Mas2]

Zr4Al3< 1030

hP7

P6

Zr4Al3

a = 543.3

c = 539.0

[V-C, Mas2]

213


MSIT®

Al–Mg–Zr

Zr3Al2< 1480

tP20

P42/mnm

Zr3Al2

a = 763.0

c = 699.8

[V-C, Mas2]

Zr5Al31395 to ~1000

tI32

I4/mcm

W5Si3

a = 1104.9

c = 539.6

[V-C, Mas2]

Zr2Al

< 1215

hP6

P63/mmc

Ni2In

a = 489.39

c = 592.83

[V-C, 1992Per]

Zr3Al

< 1019

cP4

Pm3m

AuCu3

a = 437.2 [V-C, 1992Per]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

20

40

60

80

20 40 60 80

20

40

60

80

Mg Zr


Axes: at.%(Al)

β

γ

(Mg)

αZr

ZrAl3

ZrAl2

Zr2Al

3

ZrAl

Zr4Al

3

Zr3Al

2

Zr2Al

Zr3Al

Fig. 1: Al-Mg-Zr.


400°C

214


MSIT®

Al–Mg–Zr

L+?

L+Zr Al4 3

L+Zr Al2 3

L+ZrAl2

L

Zr,

at.

%

10-0.5

10-2.5

10-2.0

10-1.5

10-1.0

10-3.0 10-2.510-2.0

10-1.5 10-1.0

Al, at.%

Fig. 2: Al-Mg-Zr.

Liquid solubility in

the Mg corner at

740°C (log scale)

[1960Sch]

Mg Mg 98.00Zr 2.00Al 0.00

Mg 98.00Zr 0.00Al 2.00 Data / Grid: at.%

Axes: at.%

L

700°C800°C

L+ZrAl3

Fig. 3: Al-Mg-Zr.

Liquidus solubility

limits in the Mg

corner at 700 and

800°C suggested by

[1958Ich, 1959Ich]

215


MSIT®

Al–Mn–Pd

Aluminium – Manganese – Palladium

Oleksandr Dovbenko, Tamara Velikanova, Sergiy Balanetskyy

Literature Data

The alloys of the Al-Mn-Pd system have been investigated in many works, however the equilibrium

diagram has not been determined for the whole composition range yet. In accordance with the assessment

by [1993Ran], based on the works of [1968Web, 1981Sol1, 1981Sol2], the MnPd2Al Heusler phase exists

in the system. The MnPd2Al alloy has bcc W type structure at high temperature, CsCl type below 1010°C

and the structure of MnCu2Al type below 710°C [1981Sol2]. Many reports with data on ternary

quasicrystalline and crystalline phases in Al-rich part of the phase diagram (up to 50 at.% Al) have been

published later.

[1968Web] determined the MnPd2Al structure as CsCl type in samples prepared by arc melting from pure

metals and annealed in vacuum at 800°C for 24 h, followed either by quenching or slow cooling over 36 h.

[1981Sol2] studied the solid state transition of the MnPd2Al alloy by X-ray diffraction. Components

(99.999% purity) were mixed and then melted in an induction furnace. The ingots were powdered and

pressed into tablets of 10 mm diameter. These were sealed in silica ampoules under vacuum of 0.1 to 0.01

Pa, homogenized at 800°C for 60 h, and quenched. The temperatures of phase transformations were

determined by diffractometry and high-temperature photography in the range of 30 to 1230°C with a heating

rate of 600K h-1. The order-disorder transition was also treated theoretically by [1981Sol1].

The icosahedral quasicrystalline phase in the composition range 5-25 at.% Pd and 10-20 at.% Mn was first

found in the rapidly solidified samples by [1990Tsa1] and in samples prepared by arc melting and then

annealed between 850 and 900°C (10h) in vacuum by [1990Tsa2]. The XRD, TEM [1990Tsa1, 1990Tsa2]

and DTA, DSC, SEM and optical microscopy (OM) [1990Tsa2] methods were used for investigation.

The structure of alloys and phase equilibria in the Al-rich part of the system were investigated by [1991Yok,

1991Tsa, 1992Was, 1992Yok, 1993Aud, 1995Goe1, 1995Goe2, 1998Sim, 1999Gru, 1999Yok, 2000Gru,

2000Kle1, 2001Goe, 2002Ant, 2002Yur]. [1991Yok, 1992Yok] studied the ternary alloys in the

composition range (14-26)Pd-(3-15)Mn (at.%), prepared from pure metals (Pd 99.996 mass% and Al, Mn

99.999 mass%) by arc melting under argon atmosphere. Ingots were annealed for 12 h at 850 and 870°C in

vacuum and quenched in water. In addition the pre-alloyed ingots, after heating in an argon atmosphere for

1 h at 1130°C, were examined by DTA at a cooling rate of 0.033 K s-1. EDX, XRD, SEM, SSM and OM

analysis were used, too. The temperature-composition section at Al80-xPd20Mnx, (with x = 0 to 20) and

equilibria of the icosahedral phase with liquid at 850 and 870°C were reported. The area of formation

(Pd+Mn=20-30 at.%) and the composition dependence of the lattice parameter (a) of the supercooled

icosahedral phase have been determined by XRD [1999Yok]. The samples were prepared from master

alloys by the zone melting process in ultrahigh vacuum using metals of the purity as in the works of

[1991Yok, 1992Yok]. Czochralski and Bridgman techniques with a flux were used to produce icosahedral

single phase samples. The composition range in which the icosahedral phase forms was also studied by

[1991Tsa, 1992Was, 2000Gru] and a partial liquidus projection in the Al corner was drafted by [1993Aud].

The authors investigated samples prepared from pure metals by induction melting under argon atmosphere.

The single quasicrystalline samples were prepared by Bridgman and Czochralski technique. SEM-XEDS,

DTA, EPMA, TEM, HREM, X-ray wavelength dispersive spectroscopy (XWDS) were used to analyze the

structure of the phases and for investigating the phase composition in the area of quasicrystalline phases.

The liquidus projection and the isothermal sections at 894, 875°C were constructed by [1995Goe1,

1995Goe2, 2001Goe]. The liquidus surface results from DTA at usual cooling and heating rates, the

isothermal sections are based on samples annealed at 894 and 875°C for 4 d, at 840°C for 6 d and 600°C for

18 d, in a range of 60-100 at.% Al. Similarly the vertical sections were constructed from thermal analysis

data. Approximately 70 samples of different compositions were investigated prepared from high purity

metals (Mn 99.985 %, Al and Pd better than 99.998 %) by induction melting. The components were placed

in a corundum crucible and this in turn was closed in a silica ampule, evacuated and filled with 650 hPa

216


MSIT®

Al–Mn–Pd

argon. Silica ampoules with extremely thin walls were used for DTA with heating and cooling rates of

2-10°C/min. Several alloys have been investigated by measuring magnetic susceptibility versus temperature

(MTA) and some samples were studied by EDX analysis.

Based on a limited number of samples annealed for 45 min [2000Kle1] determined liquid-solid equilibria

at 685, 730, 840, 920, 945 and 952°C and proposed a projection of the liquidus surface in the region 50-100

at.% Al. The alloys were prepared from 99.999 mass% Al, 99.9 mass% Pd and Mn by induction melting

and some of the ingots were re-melted in alumina crucibles under argon atmosphere in order to produce

large single grains by the Bridgman technique. The crystallization sequences were deduced from

SEM-XEDS, XWDS and TEM methods and from DTA samples of about 0.15 g, applying cooling and

heating rates of 5°C/min.

In the Al-rich part of the system isothermal sections at 880, 870, 850 and 790°C were reported by [1999Gru,

2002Yur] in the vicinity of the quasicrystalline phases. The 33 ingots were prepared by induction

(levitation) melting in a cold crucible under argon atmosphere. Part of the ingots were annealed at 880, 870

and 850°C for 65-70 h, at 790°C for 590 h and at 710°C for 1450 h and quenched in water. The samples

were examined by OM, SEM-EDX and XRD methods. The bulk composition of some single-phase samples

was measured by inductively coupled plasma optical emission spectroscopy, ICP-OES. Single-phase

samples whose compositions were measured by ICP-OES were used as standards for the correction of the

EDX data. Selected samples were studied by DTA at heating and cooling rates of 20°C/min and by TEM.

The TEM examinations were carried out on the powdered materials, which were spread on copper grids

with carbon film.

The composition of the liquid phase coexisting with icosahedral and compositions of adjacent ternary

crystalline phases were determined by [1998Sim]. Five ternary alloys (Mn7.2Pd20.7Al72.1, Mn5Pd18Al77,

Mn3.6Pd16.6Al79.8, Mn3.5Pd20Al76.5 and Mn1.7Pd19.2Al79.1) and one binary alloy (Pd19Al81) were

investigated, prepared from Al 99.999 mass%, Pd and Mn 99.9 mass% by induction melting in a cold

crucible under argon atmosphere. DTA, chemical composition analysis, structure analysis by X-ray or

electron diffraction and neutron scattering were applied to the as-cast samples. All investigated ternary solid

phases were found to melt incongruently.

Thin films were deposited by simultaneous evaporation of the metals from separate sources on carbon

substrates or on glass plates at temperatures up to 500°C and the structures and compositions of ternary

films were investigated by TEM, electron diffraction, and EDX [2002Ant]. Icosahedral order was observed

for aluminium contents above 75 at.% and a phase diagram for thin films at a deposition temperature of

475°C was constructed.

There is a number of experimental and theoretical works devoted to the crystallographic investigation of the

quasicrystalline and periodic crystalline ternary phases in the Al-rich part of this system [1991Bee,

1991Bou, 1991Don, 1992Bou, 1992Was, 1993Aud, 1993Bou, 1993Dau, 1993Hir1, 1993Ste, 1993Sun,

1993Tsa, 1993Was, 1994Aud, 1994Bee, 1994Li1, 1994Li2, 1995Bee1, 1995Bee2, 1995Boi, 1995Ish,

1996Bou, 1996Yam, 1997Ama, 1997Hae, 1997Kle, 1997Kra, 1997Mat, 1997Son, 1997Zur, 1998Ber,

1998Boi, 1998Mat, 1998Wan, 1999Aud, 1999Cap, 1999Fis, 2000Bee1, 2000Bee2, 2000Dun, 2000Fra,

2000Fre, 2000Gwo, 2000Hir, 2000Jan, 2000Jac, 2000Kaj, 2000Kle3, 2000Let, 2000Nic, 2000Qua,

2000Sch1, 2000Shr, 2000Sta, 2000Ste1, 2000Ste2, 2000Uch, 2000Yam, 2001Nau, 2002Hir, 2002Lei,

2002Shr, 2002Yam, 2002Yan, 2002Zha1, 2002Zha2]. [1992Bou, 1993Bou] examined single grain samples

of the icosahedral phase by XRD and neutron diffraction. [1993Sun] investigated Al70Pd20Mn10 samples

rapidly solidified and annealed at 800°C by TEM, XRD, HREM, EDXA. TEM and XRD methods were

used by [1991Bee, 1992Was, 1993Was, 1995Ish, 1998Boi, 1999Aud, 1994Aud]. The modifications of the

icosahedral phase have been studied by [2000Hir, 1998Boi, 1995Ish, 2000Let, 1999Aud, 2002Yam].

[1995Ish] investigated samples of composition MnxPd29-xAl71, (6.5 < x < 9.5), prepared from pure elements

(Al 99.999 mass%, Pd 99.95 mass%, Mn 99.99 mass%) in a plasma-jet furnace. The samples, put into a

graphite crucible and sealed in silica tubes, were annealed for 50 h at 803 4 and 48-400 h 602 2°C,

quenched into water and subsequently into liquid nitrogen. An alloy of Mn8Pd21Al71 composition annealed

at different temperatures was examined by powder X-ray diffraction using CuK radiation at room

temperature by [2000Hir], whereas [1998Boi] performed in-situ heating experiments using synchrotron

light source. [2000Let] examined ingots of composition Mn8.8Pd21.4Al69.8 by X-ray and in-situ neutron

217


MSIT®

Al–Mn–Pd

diffractions. [2002Yam] investigated a sample at Mn8Pd21Al71 by XRD with an IP-Weissenberg camera

and carried out a structure refinement of the icosahedral quasicrystal. A two-axis diffractometer was used

by [2000Shr] to measure the diffuse background for two different single grained samples: icosahedral

(Mn8.2Pd21.4Al70.4) and ´ (Mn4Pd22Al74) phases.

The T and R orthorhombic ternary phases in this system were found by [1993Aud, 1994Aud, 1997Kle,

1997Mat, 1998Ber]. [1997Kle] obtained single crystal T-Al3(Pd,Mn) phase by Bridgman technique from

an ingot of nominal composition Mn19Pd7.1Al73.9, prepared by induction melting in a cold crucible from

the pure elements (99.999 mass% Al, 99.9 mass% Pd, 99.9 mass% Mn). [1997Mat] performed XRD

analysis of the single crystal of nominal composition Mn25Pd5 Al70 prepared by arc melting pure metals

under Ar atmosphere. [1998Ber] examined the sample of the Mn3Pd9.3Al87.7 composition by XWDS, TEM

and HREM and observed the structural defects near the R- interfaces and within both phases.

Single crystals of the ternary ´ phase were obtained by [1996Bou] using the Bridgman technique from ingot

with starting nominal composition of Mn3.5Pd19Al77.5.

[2002Shr] observed this phase in samples of the composition around 4.5Mn-22.9Pd-72.6Al (at.%), (which

corresponds to the nominal composition of the ´ phase, prepared by the Bridgman crystal growth method.

Powder and single crystal XRD experiments were carried out.

The electronic structure and electronic densities of the decagonal phase Mn17Pd13Al70 and related

crystalline phases have been calculated by [1997Kra] and a structure-induced pseudogap in the Al band was

shown to exist in decagonal as well as in related ternary periodic phases of similar composition. The

stabilization by the Hume-Rothery-like band-gap was found to decrease in the sequence: crystalline

(Al6Mn) - icosahedral (Mn8Pd22Al70) - decagonal (Mn17Pd13Al70). The results of the photoemission

spectroscopy on the electronic structure of quasicrystals have been reviewed by [2000Sta] and the existence

of the theoretically predicted pseudogap at the Fermi level being confirmed.

Binary Systems

The Al-Pd and Al-Mn systems are accepted from [2003Bal] and [1997Oka, 2003Pis], respectively. Data

concerning the Mn-Pd system are from [Mas2].

Solid Phases

Crystallographic data on the known unary and binary phases as well as recently reported ternary ones are

listed in Table 1. A peculiar feature of the system is the formation the two stable quasicrystalline phases:

decagonal 2, usually labeled as “D”, and icosahedral 3, usually labeled as “I”. Another peculiarity is the

close crystallographic relationship between the solid solutions based on binary phases such as Mn4Al11 (h),

and , MnAl4 and other periodic phases, and the quasicrystalline phases. The mutual solid solubility of the

isostructural , MnAl and ( Mn) on the one hand, with the , PdAl (h) and , MnPd (h) phases on the other

are reported by [1995Goe1, 1995Goe2, 2000Kle1]. However, the authors did not give details about the

phase relationships in the range where disordered bcc , MnAl of the W type and the ordered CsCl type ,

PdAl do coexist.

At 1010°C the cubic W type phase transforms into the cubic CsCl type phase [1981Sol2] which at 710°C

transforms into the 1 MnPd2Al Heusler phase. The composition of the decagonal 2 quasicrystalline phase

is very close to Mn18.1Pd12.1Al69.8, its melting point is at 896°C according to [1995Goe1, 1995Goe2]. A

crystalline “pseudo-decagonal” phase of the same composition identified as DH, a high temperature

modification of D, with a B-centered orthorhombic cell is reported by [1995Bee1, 1995Bee2, 1995Goe2]

above 864°C.

The homogeneity range of the icosahedral phase 3 is Mn8-10.2Pd20.3-23.2Al68-69.5 according to [2000Kle1]

and in temperature range from 880 to 710°C it is Mn6-10Pd19.2-24.5Al69.5-70.8, according to [1999Gru,

2000Gru, 2002Yur]. Three modifications of the icosahedral phase were found by [1995Ish, 1998Boi,

2000Hir, 2000Let] at different temperatures and in close compositional vicinity, i.e. Mn9.2Pd22.0Al68.8,

Mn8.8Pd21.4Al69.8, and Mn8.7Pd22.0Al69.3, labeled as F, F2 and F2M, respectively. The high temperature F

phase has a 6 dimensional reciprocal primitive cubic cell with strong chemical order. The F2 phase is

considered as a superstructure of the F-phase and could be described as a P type 6D hypercubic lattice with

218


MSIT®

Al–Mn–Pd

parameter aP = aF (aP = 2088.1 pm), or a 6 dimensional diamond-type structure with aF2 = 2aF

(aF2 = 2581.0 pm) [1995Ish, 2000Hir]. According to [2000Let] the F2 phase is not stable and corresponds

to a transient state in the process of the transformation of the icosahedral F phase to the F2M phase. The F2M

phase has a domain structure with a cubic but non-periodic symmetry [1998Boi] and forms from high

temperature F phase at 715°C [2000Let].

According to [1993Aud, 1994Aud, 1996Sun, 1997Kle, 1998Ber, 1996Bou, 2000Shr, 2000Kle1, 2002Shr,

2002Tex] four additional orthorhombic ternary phases designated as R, T, ´ and exist. The R, T and ´

phases form from the liquid by peritectic reactions. According to [2000Kle1] the composition range of the

R and T phases on solidus overlap (3.5-6.6 at.% Pd, 16-25.5 at.% Mn), but separated in the liquidus surfaces

of the system. The T phase is reported to be isostructural to the MnAl3 phase where Pd substitutes Al only

in one position and this position shifts along the b axis. The R phase is reported to be isostructural to

Mn11Ni4Al60 [1994Aud, 2000Kle1, 1997Kle]. According to [1997Kle] the T phase might transform into an

R phase with aR = aT, bR = bT, cR = -1cT. From HREM data both the R and T phases are pointed to exhibit

different tilings in their ac plane but built from the same flattened hexagon. The orthorhombic ' phase has

a composition of Mn5.0-4.6Pd22.1-22.4Al73.3 at 730°C according to [2000Kle1], and exists in several

modifications [1996Kle, 2002Shr]. According to [1996Bou] the structure of this phase is very close to 6

the “PdAl3” phase in which the Mn atoms substitute only two Pd sites of high symmetry among 12 Pd sites.

The additional two modifications, called ´_3 and ´_5 by [2002Shr], have a periodicity defined as

a ´_n = a ´, b ´_n = b ´ and c ´_n = c ´´ (n+ ), where n equals 3 or 5, respectively and is the golden mean.

The structure of the ´_5 phase is different from the structure of the ´ phase, but not very much. Also, a

modification ´_2 with c ´_2 = c ´ (2+ ) may exist in this system as well [2002Shr]. According to

[2000Kle1] the T, R, and ´ phases are ternary compounds and have to be distinguished from the binary

phases. One of the reasons for this conclusion was that in the samples Mn1.7Pd19.2Al79.1 (obtained by

Bridgman) and Mn11.4Pd1.3Al87.3 (not annealed sample after DTA experiment) interfaces between the

ternary and binary phases have been observed, i.e. between “PdAl3” and ´ phases, or Mn4Al11 and T

phases, respectively. These, however, also could result from segregation and nonequilibrium conditions of

the samples. The reported composition ranges of the R, T and ´ phases are close to the binary MnAl4,

Mn4Al11(h), 6 and 28 phases and their crystal structures are very similar; R phase is isostructural to

metastable -MnAl4 phase, T phase to Mn4Al11 (h) or MnAl3, and ´ is isostructural to the 6 phase.

According to [1995Goe1, 1995Goe2, 1999Gru, 2002Yur] the H solid solution, based on the Mn4Al11 (h)

binary phase, and “ ” solid solution exist in the composition range reported for the T and ´ phases,

respectively, where “ ” is considered to be the ternary extension of the 6 and 28 binary phases.

According to [2002Yur] both 6 and 28 binary phases were observed in the ternary overall “ ” field and, in

addition, 22 and 34 ternary phases were observed by TEM. However the SEM/EDX analysis did not reveal

any compositional inhomogeneities typical for such a multiphase sample.

The 6, 22, 28, 34 phases can be correlated with ´, ´_2, ´_3 and ´_4 respectively, according to

designation by [2002Shr]. The coexistence of the 6, 22, 28, 34 phases in the “ ” continuous range of the

ternary Al-Mn-Pd system is unclear. Their stability at temperatures under investigation is consistent with

long-term annealing and the “ -phase” is stable in the ternary system up to the melting temperature of 845°C

according to DTA data [2002Yur]. Thus the question wether l ( ´), R and T phases are solid solutions based

on binary phases or individual ternary phases is open and needs additional investigations. All data for the

R, T and ´ phases are given in Table 1 together with data for , MnAl4, Mn4Al11(h) and 22, 34 phases,

respectively.

The stability range of the cubic and orthorhombic phases reported by [1992Was, 1993Aud, 1994Aud,

1994Li2, 1998Ber, 1999Hip, 2000Kle1], based on XRD and TEM data, is not clear. An orthorhombic

phase, labeled as 2-R, with space group Amm2 and lattice parameters a = 1243, b = 2030, c = 6250 pm, or

a = bR, b = 2cR, c = 2aR, according to [2000Kle1] is reported by [1993Aud, 1994Aud]. The periodic cubic

approximant (2/1) with lattice parameter a = 2030 pm was observed by [1992Was] in the Mn4Pd26Al70

sample after heat treatment at 750°C and by [1994Li2] in the Mn10Pd20Al70 sample annealed at 800°C for

3 d. The authors of [1994Li2] also found in this sample the cubic phase (1/1) with lattice a parameter

a = 1240 pm. Four orthorhombic phases: the (1/1,1/1) with lattice parameters a = 1260, b= 1240, c = 1480

pm; (2/1,5/3) with a = 1920, b = 1240 c = 6140 pm; (5/3,3/2) with a = 5050, b = 1240, c = 3780 pm and

219


MSIT®

Al–Mn–Pd

the (8/5,5/3) with a = 8400, b = 1240, c = 6200 pm were found by [1994Li2] in a Mn15Pd15Al70 sample

which was annealed for 3 d at 800°C and quenched in water.


A partial reaction scheme is presented in Fig. 1a for the Al-rich part, according to [1995Goe1] with some

corrections made according to the accepted binary diagrams and data of [1999Gru, 2002Yur] on additional

invariant ternary transition equilibrium 3+ 2+ (U4) at 860°C. The thermal DTA effects obtained by

[1995Goe1, 1995Goe2] at 864°C are shown as dotted lines in Figs. 9-11, 14. These lines are very close to

the above-mentioned temperature reported by [1999Gru, 2002Yur], and interpreted by [1995Goe2] as

DH D transition temperature. It is not clear wether these temperatures really correspond to different

processes. The temperature of the invariant equilibrium L+ 3+ (U3) that is 867°C according to

[1995Goe1] and 870 < T < 880°C according to [1999Gru, 2002Yur], was accepted from the last work

(~875°C). The temperature of the reaction P3 (L+ + 3 ) is accepted from [2002Yur].

According to [1995Goe1, 1995Goe2] the decagonal quasicrystalline phase can form in two ways.

According to the equilibrium diagram, the decagonal phase crystallizes from the liquid by peritectic reaction

P1: L+Mn4Al11+ 2 at 896°C (Figs. 1a, 2a) before the icosahedral phase. In a metastable state, for

example after rapid solidification, 2 phase is formed in the solid state at ~770°C [1995Goe1] by the

reaction H+ + 3 2 (Figs. 1b, 2b). The icosahedral phase is formed from the liquid by peritectic reaction

at approximately 893°C [1995Goe1, 1995Goe2, 1999Gru, 2002Yur].

As mentioned in the chapter Solid Phases, according to [2000Kle1] the T, ´ and R phases are ternary

compounds and form from the liquid by peritectic reactions. The authors [2000Kle1] however noted

nonequilibrium state of the samples therefore this data were not used in the discussion of the reaction

scheme and this question needs additional investigations. The ternary eutectic reaction L (Al)+ ´+R at a

temperature 617 5°C is reported by [1998Sim]. According to [1995Goe1] the equilibrium U9 and binary

eutectic e4 occur at the same temperature. Thus, the nature of this transition remains undefined.

Liquidus and Solidus Surfaces

The projection of the liquidus surface in the Al-rich part of the system in Fig. 2a is given after the data by

[1995Goe1, 1995Goe2, 2001Goe] and exhibits surfaces of the binary phases and only two stable ternary

compounds (decagonal and icosahedral phases) firmly determined. The surface with the stable liquid phases

and the primary solidification of 2 and the corresponding nonvariant equilibria were extrapolated in

[1995Goe1] from the long term heating curves of the DTA and MTA plots and from the heat treatment

experiments on alloys of 18 at.% Pd and 10-18 at.% Mn.

A continuous ternary (MnPdAl) liquidus surface, extending from Al-Mn (55.1 to 71.7 at.% Al) to Al-Pd

(38.3 to 70.5 at.% Al) binary boundary liquidus, is proposed by [1995Goe1, 1995Goe2, 2000Kle1]

notwithstanding that the different crystal structures (W and CsCl types) of the binary compounds form

corresponding mutual ternary solid solutions. It can not be excluded that the peculiar shape of the (MnPdAl)

liquidus reflects this fact. But mutual transformations of cubic phases - W and CsCl type are not

investigated. The doted lines in Fig. 2 correspond to the expected phase transition of cubic (MnPdAl) solid

solution and to the monovariant line l+ 28+ 6 that starts from peritectic reaction p6 in the Al-Pd system.

The metastable liquidus surface is projected in Fig. 2b for the Al-rich part of the phase diagram after

[1995Goe1, 1995Goe2]. A different partial metastable liquidus surface projection was proposed by

[2000Kle1]. The extensions of the liquidus phase fields was determined from the solidification sequences

of DTA samples and the liquid-solid equilibria in samples annealed for 45 min at 952, 945, 920, 840, 730,

685°C and subsequently quenched. The three additional ternary compounds (T, R, ´) on the solidus

supposed by [2000Kle1] and the separate liquidus fields corresponding to these ternary phases are accepted

to exist in addition to the one quasicrystalline phase (icosahedral): for T and R phases in the H phase

liquidus surface (after [1995Goe1]) and for the ´ phase in “ ” region. The temperature limits of the

monovariant reactions obtained by [2000Kle1] are in good agreement with the data of [1995Goe1,

1995Goe2] if one supposes that the phase fields H and “ ” [1995Goe1, 1995Goe2] correspond to the phase

fields R, T and ´ [2000Kle1], respectively.

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The partial solidus surface projection corresponding to the liquidus presented in Fig. 2 is given in Fig. 3

according to data [1995Goe1].

Isothermal Sections

The partial isothermal section (in the Al-rich part of the system) at 894°C is given in Fig. 4 after work

[1995Goe1]. The sections at 875, 840, 710 and 600°C presented in Figs. 5-8 result from evaluating all the

data presented by [1995Goe1, 1999Gru, 2002Yur]. Mutual transformations of cubic phases - W and CsCl

type, which are to take place in the ternary system, are not investigated. In isothermal sections above 840°C

which is the temperature of eutectoid decomposition of the bcc phase in the Al-Mn system, the phase

fields of (W) and (CsCl type) are separated in the same way as in the solidus projection in Fig. 3.

According to [1995Goe1] at 894°C (Fig. 4) the decagonal 2 phase coexists in equilibrium with liquid and

solid solutions based on the binary cubic phase and H phase, which is a ternary extension of the binary

Mn4Al11 (h) phase. According to [1995Goe1, 1999Gru, 2002Yur] the last phase is stabilized by Pd

additions in the ternary alloys down to temperatures that are lower then its decomposition temperature in

the binary system (stable down to 600°C [1995Goe1]).

Two ternary compounds 2 (decagonal) and 3 (icosahedral) in the isothermal sections at 875, 840, 710 and

600°C are shown. At 840°C (Fig. 6) the overall “ ” phase field is a ternary extension (solid solution) of the

binary phases, which includes 6, 22, 34 and 28 phases stabilized by Mn to higher temperature than in

the binary Al-Pd system and at 600°C this phase field joins the edge boundary system.

There is the nonvariant four-phase equilibrium L+ 3+ (transition type) at 867°C according to

[1995Goe1], but according to more precise data of [1999Gru, 2002Yur] it should be in the temperature

range 880-870°C (Fig. 5) (accepted here ~875°C).

The authors [1995Goe1] reported the equilibrium between 2 phase and solid solution at 875, 840 and

600°C. However, according to the detailed investigation of [1999Gru, 2002Yur] at 850 and 710°C the 2

phase is in equilibrium with the binary phase as a result of the transition type reaction 3+ 2+ which

was determined at 860°C (from the DTA data for Al-20Pd-12Mn alloy annealed at 850°C). The isothermal

sections at 840, 710 and 600°C are shown in Figs. 6-8 taking into account the above mentioned data of

[1999Gru, 2002Yur].

The phase equilibria at 790°C are similar to those for 710°C, [1999Gru, 2002Yur].

At 875°C a broad band of liquidus extends from the Al-Pd binary system to 10 at.% Mn and the 3 phase

was found to be homogeneous from 7.5 to 10 at.% Mn, [1995Goe1].

The homogeneity range of the 3 phase at 880°C extends from 70.2 to 71.2 at.% Al and from ~8.2 to

10.4 at.% Mn [2002Yur]. The homogeneity range of the 3 phase at 870°C spans from 69.6 to 71.6 at.% Al

and about 8 to 10.5 at.% Mn [2000Gru, 2002Yur]. The composition of the H phase in equilibrium with 2

and 3 is 71-73 at.% Al and 6-7 at.% Mn. The solubility of the Mn in the solid solution based on the phase

reaches up to 2.0 at.% at 850°C and 1.6 at.% at 710°C. The homogeneity range of the icosahedral phase

extends from 70.0 to 71.6 at.% Al and 6.7 to 10 at.% Mn at 850°C and from 70.0 to 71.0 at.% Al. and 5.6

to 8.5 at.% Mn. The homogeneity range of the 2 phase at 710°C spans from 69.4 to 70.2 at.% Al and from

14.8 to 17.3 at.% Mn.

The partial isothermal sections at 850 and 870°C reported by [1992Yok] expose the liquid-solid equilibria

of the 3 phase. These data are in satisfactory agreement with the above given data.


The temperature - composition sections across 10 at.% Pd, 20 at.% Pd, 70 at.% Al and 6 at.% Mn are given

in Figs. 9-12 according to [1995Goe1]. The sections Mn32.8Al67.2 - Pd27Al73, Mn31.2Al68.8 - Pd29Al71 and

Pd3Al97 - Mn13Pd30Al57 (Figs. 13-15) are given according to [1995Goe2] taking into account the above

mentioned corrections from the isothermal sections. The latter section was obviously mislabeled as

Pd3Al97 - Mn22.8Pd30Al47.2 by [1995Goe2].

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Thermodynamics

The heat capacity of a Mn10Pd20Al70 alloy was measured by [1997Ina] in the 1-350 K temperature range

with an adiabatic calorimeter (above 5 K) and an isoperibolic one installed in a helium-3 cryostat (below 7

K). The samples were obtained by arc melting pure metals and from their melts single grains of the

icosahedral phase were grown at 1200°C at a growth rate of 1mm/h from melt (Bridgman method). The

Cp/T vs T2 plot from these data is presented in Fig. 16.

The specific heat of the icosahedral phase at constant pressure (Cp) and at constant volume (Cv) was

investigated by [2000Eda]. Below 427°C (Fig. 17a) Cp shows the typical behavior of ordinary solids, it

appears to approach 3 kB asymptotically with increasing temperature, obeying the Dulong-Petit’s law.

Above 427°C, however, it increases dramatically and reaches approximately 5 kB at 807°C. The

temperature dependence of Cv has been calculated in the same work (Fig. 17b). The sample Mn9Pd20Al71

was prepared from elemental constituents by arc melting under argon atmosphere and annealed at 750°C

for 72 h. XRD, DSC measurements with heating rate 10°C/min were applied in the temperature range of

277-807°C and XRD measurements at higher temperatures (7-527°C).

For the same phase the specific heat Cp(T) was measured by [1998Wae] in the 1.6-280 K temperature range

using two different techniques (Fig. 18). A standard relaxation technique was adopted at 1.6-30 K, and

adiabatic continuous-heating calorimeter was used in the 15-300 K range. [1998Wae] grew single grain

with final composition Mn9Pd22.8Al68.2 by Czochralski method. The polygrain Mn9Pd21Al70 was

synthesized using Al (99.997 mass%), Pd (99.9 mass%) and Mn (99.94 mass%). For homogeneity the

sample was arc melted several times and subsequently quenched into water from 800°C.

[2000Kaj] concluded the thermal expansion of the Mn9Pd20Al71 alloy in the temperature range between

-263 and 427°C from X-ray diffractometry. The alloy was prepared by arc melting and annealed at 750°C

for 72 h. The linear thermal expansion coefficient (T) of the icosahedral phase is about half of that for the

pure aluminium phase at room temperature and does not show negative thermal expansion at low

temperature. For the 3 phase, the same author estimated Cv(T) above 77°C by the Debye approximation

using the reported Debye temperature (a limiting value as T of 188°C).

Generalized vibrational densities of states (GVDOS) at T = 23, 327, 527, 727°C have been measured by

[2002Sch]. The sample with composition Mn4Pd22Al74 ( ´ phase) was prepared from pure elements

(Al 99.999%, Pd 99.95% and Mn 99.99%) arc melted and annealed for 6 d at 590°C, 2 d at 825°C.

Subsequently it was cooled down to room temperature with a rate of 5°C/min. Neutron scattering

experiments under vacuum have been made and the temperature dependence of the heat capacity was

calculated from GVDOS data between 2 and 32 K. The specific heat Cp (Fig. 19) was measured in the 2-80

K temperature range by relaxation-type method and the sound carrying velocities was calculated.


The diffusion coefficient of 63Ni in Mn9Pd21Al70 alloy [2000Zum] and those of 65Zn and 114In in single

icosahedral quasicrystals of undefined composition [2000Gal] were measured by radioactive tracers. The

activation enthalpies of diffusion Q in icosahedral phase was found to be 209.0 kJ mol-1 for Ni, 121.3 1.3

kJ mol-1 for Zn and 165.9 5 kJ mol-1 for In. The diffusion of 103Pd and 195Au in icosahedral quasicrystal

(Mn8.5Pd21.3Al70.2) under proton irradiation was investigated by [2000Blu].

Paramagnetic Curie (-215 30 K) and Néel (240 K) temperatures were determined by [1968Web] for the

MnPd2Al Heusler phase. The magnetic properties of the quasicrystalline and related crystalline phases have

been studied in several works [1998Sim, 1999Fis, 1999Hip, 1999Yok, 2000Lai, 2000Sch2, 2000Sim,

2002Miz, 2002Mot]. The neutron scattering experiments on several Al-Mn-Pd liquid alloys with Mn

content between 3.5 and 7.2 at.% were carried out by [1998Sim]. The comparison between polarized

neutron scattering experiments and magnetic susceptibility suggested that the magnetic moments are

present in the liquid state but not in the solid. Temperature dependence of the magnetic susceptibility has

been measured in the liquid state and during the solid-liquid transformation. [1998Sim] found that all

investigated samples are diamagnetic at room temperature. It is found that the appearance of a magnetic

moment on a Mn atom strongly depends on its position in the crystalline lattice [2000Lai]. [1999Hip]

investigated the magnetic properties of the Mn6Pd24Al70 alloy, produced by planar flow casting and

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annealed at 815°C for 2 h, and three orthorhombic phases: ´-Mn4.6Pd23.5Al71.9, as single crystal obtained

by Czochralski growth, T-Mn21.7Pd5.2Al73.1 and T-Mn16.6Pd4.9Al78.5 obtained by Bridgman growth. No

magnetic Mn atoms are presented in the Mn6Pd24Al70 alloy, none in the ´ phase nor in the Mn-poor T

phase. However, a fraction of the Mn atoms carries magnetic moments in the Mn-rich T phase in which a

spin-glass transition occurs at -259°C. The temperature dependence of the dc magnetic susceptibility of

these phases has been studied. A strong decrease of the total magnetic moment after annealing was observed

by [2000Sch2] using a vibrating sample magnetometer to study an annealed Mn9Pd20Al71 sample which

was produced from the melt in a vacuum furnace and cast in a cold Cu mould, annealed at 825°C for 3 h

and quenched in water; annealed again at 625°C for 0.5 h and cooled down to room temperature with a rate

of 1 K min-1. The number and the magnitude of the magnetic moments were determined by measuring the

field and temperature dependence of the total magnetic moment. [2002Mot] measured the temperature

dependence of the magnetic susceptibility of the F2M phase (Mn8Pd21Al71 sample) annealed at different

temperatures in the range of -271 to 327°C and found that the susceptibility decreases with increasing

temperature from -273 < T < 73°C, it increases with increasing temperature from 73 < T < 327°C. The data

could be fitted to the Curie-Weiss law with an additional term proportional to the square of the temperature.

The additional term indicates a pseudogap in the electronic density of states at the Fermi energy. The

magnetic properties of Mn8.2Pd20.3Al71.5 as-cast samples and those of samples annealed at 727°C for 50 h,

have been investigated by [1999Yok] from -268.8 to RT when an external magnetic fields up to 10 kOe is

applied. The magnetic order at low temperature was found to be typical for the canonical spin glass

phenomenon. The value of diamagnetic component susceptibility ( 0) was estimated to be about

3.3-4 10-7 emu/gOe. The results of electrical properties measurements can be found in [1997Son, 1999Fis,

1999Yok, 2000Tho, 2002Bil, 2002Dem, 2002Miz, 2003Ban, 2003Cap]. In particular, [2002Bil] measured

the thermoelectric properties of polygrained icosahedral quasicrystal Mn8.5Pd19.5Al72. They used a kind of

self-flux technique, where the ternary melt is first slowly cooled, and then the remaining melt is decanted

in the temperature range from -263°C to RT. The electrical resistivity is RT = 1.2 m cm-1 at RT and

increases with decreasing temperature showing a maximum at -153°C. The thermoelectric power is positive

in the whole temperature range and at room temperature S = 70 V K-1. Thermal conductivity at RT is K =

3.4 W/mK and shows a maximum at -243°C and a broad minimum around 148°C. According to [1999Yok]

the electrical resistivity in the -268 to -3°C temperature range shows anisotropic dependence on the different

symmetrical axes (2-, 3-, 5-fold directions) in the Mn10Pd20Al70 as-grown sample prepared by Czochralski

method. After annealing at 627°C for 50 h differences among the electric resistivities along the various

directions become smaller than that in the as-grown state. The electric resistivity of Mn10Pd20Al70 thin

films was measured from -185 to 12°C by [1997Son] and showed strong negative temperature dependence.

The thin films were prepared by laser ablation on fused silica at different deposition temperatures (-196,

RT, 165, 350°C) and examined by XRD. Several works were dedicated to the investigation of the

mechanical properties of the icosahedral alloys [1996Tan, 1999Yok, 2000Bar, 2000Bru1, 2000Bru2,

2000Feu, 2000Kaj, 2000Mes, 2000Sch3, 2000Sch4, 2002Duq, 2002Lei, 2002Kab, 2002Tak, 2002Tex].

[1996Tan] measured the elastic constants for a Mn6Pd24Al70 alloy (density 5150 kg m-3) over a temperature

range from -269 to 800°C by the rectangular parallelepiped resonance method. At room temperature the

results were: Lamé constants (c12) = 74.9, (c44) = 72.4, Young modulus 182 GPa, bulk modulus 123 GPa,

Poisson ratio 0.254. The temperature dependence of the above mentioned properties have been determined.

The Vickers hardness of the Al70Pd20Mn10 sample presented in Fig. 20 is noticeably different along the

2-fold, 3-fold and 5-fold directions [1999Yok]. The plastic deformation of icosahedral Al-Mn-Pd single

quasicrystals and ´ phase has been investigated by [2000Bar, 2000Bru2, 2000Feu, 2000Kle2, 2000Mes,

2000Wan, 2002Kab, 2002Tex].

The oxidation of the Mn9Pd20Al71 icosahedral quasicrystals (sample made by hot isostatically pressing) and

Mn8.5Pd21Al70.5 (obtained by arc melting) at 800°C is strongly influenced by the evaporation of Mn,

according to [2000Weh]. The reflectivity of the icosahedral phase (Mn8.4Pd21.2Al70.4 sample prepared by

induction melting) was found to be very high at low frequency in the far-infrared range, and then it

decreases suddenly [2002Dem]. [2000Lan] examined the phase structure and evolution of the

quasicrystalline coatings, thermal diffusivity, hardness, and friction coefficient. They found that the

quasicrystalline coatings with composition Mn10Pd20Al70 (prepared by a plasma spray process from gases

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of atomized powders) show cracking after heat treatment at 700 and 800°C. In the as-sprayed condition the

measured thermal diffusivity of the coatings is low, but heat treatment increases the thermal diffusivity with

increasing heat treatment temperature. The thermal diffusivity of the coatings thus increases as the volume

fraction of icosahedral phase decreases and the decagonal and crystalline phases increase. The coefficients

of friction of the coatings are reasonably low in the as sprayed condition. Little change is observed after heat

treatment up to 600°C, but after heat treatment at 700 and 800°C the values are considerably reduced.

[2000Bel] reviewed the main results on transport properties of quasicrystals and introduced the necessary

mathematical background. Some questions about practical applications of the quasicrystals are discussed in

[2000Cyr, 2000Dub].

Miscellaneous

The concentration regime and lattice parameter changes for the supercooled icosahedral phase are reported

by [1999Yok]. The temperature dependence of the (quasi-)lattice constants measured by high temperature

X-ray diffraction experiments are given in Fig. 21 according to [2000Kaj].

The high-temperature solution growth of large single-grain crystals and quasicrystals of the Al-Mn-Pd

system are discussed in [1999Fis, 2000Fis, 2001Can]. Such kind of methods as low-energy electron

diffraction (LEED), X-ray photoemission spectroscopy (XPS), scanning tunneling microscopy (STM),

ultraviolet-photoemission spectroscopy (UPS), X-ray photoelectron diffraction (XPD), secondary-electron

imaging and Auger electron spectroscopy (AES) are used to investigate the atomic and electronic structure,

properties, decomposition at elevated temperature, the surface structural phase transitions, voids in the

as-grown and annealing single quasicrystalline icosahedral crystal, etc. The results can be found in the

publications [1997Gie, 1997She, 1998Bol, 1998Gie, 1998Was, 2000Bol, 2000Cap1, 2000Cap2, 2000Klu,

2000Led, 2000Nau, 2000Ros, 2000Sch5, 2000Sch1, 2002Klu, 2000Klu, 2002Pap, 2003Ebe]. The

structural perfection of icosahedral phase has been studied by mechanical spectroscopy [2000Dam] as well

as by combined synchrotron X-ray diffractometry and imaging technique [2000Man] and by means of

positron annihilation spectroscopy and time-differential dilatometry [2000Bai].

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Al70Mn17Pd13 Quasicrystal”, Philos. Mag. B, B70(2), 215-240 (1994) (Crys. Structure,

Experimental, 33)

[1994Li1] Li, X.Z., Dubois, J.M., “Structural Sub-Units of the Al-Mn-Pd Decagonal Quasicrystal

Derived from the Structure of the T3 Al-Mn-Zn Phase”, J. Phys.: Condens. Matter, 6,

1653-1662 (1994) (Crys. Structure, Theory, 22)

[1994Li2] Li, H.L., Kuo, K.H., “Some New Crystaline Approximantas of Al-Pd-Mn Quasicrystals”,

Philos. Mag. Lett., 70(1), 55-62 (1994) (Crys. Structure, Experimental, 32)

[1994Shi] Shi, N.C., Li, X.Z., Ma, Z.S., Kuo, K.H., “Crystalline Phases Related to a Decagonal

Quasicrystal. I. A Single-Crystal X-Ray Diffraction Study of the Orthorhombic Al3Mn

Phase”, Acta Crystallogr., Sect. B: Struct. Crystallogr. Crys. Chem., B50, 22-30 (1993)


[1995Bee1] Beeli, C., Stadelmann, P., Gödecke, T., Lück, R., “The Decagonal Al-Mn-Pd Phase and its

Modification”, Proc. Intern. Conf. on Aperiodic Crystals (Aperiodic 94), Chapius, G.,

Paciorek, W., (Eds.), World Scientific Publ., Singapore 1995, 361-365 (1995) (Crys.


[1995Bee2] Beeli, C., Stadelmann, P., Lueck, R., Gödecke, T., “Decagonal Al-Mn-Pd Quasicrystals

Free of Linear Phason Strain”, Proc. 5th Intern. Conf. on Quasicrystals, Janot, C.,

Mosseri, R. (Eds.), World Scientific. Publ.; Singapore 1995, 680-683 (1995) (Crys.


[1995Boi] de Boissieu, M., Boudard, M., Hennion, B., Bellisent, R., Kycia, S., Goldman, A.I.,

Janot, C., Audier, M., “Diffuse Scattering and Phason Elasticity in the AlPdMn Icosahedral

Phase”, Phys. Rev. Lett., 75(1), 89-92 (1995) (Crys. Structure, Experimental, 24)

[1995Goe1] Goedecke, T., Lueck, R., “The Aluminium-Palladium-Manganese System in the Range

from 60 to 100 % Al”, Z. Metallkd., 86(2), 109-121 (1995) (Equi. Diagram, Experimental,

#, *, 30)

[1995Goe2] Goedecke, T., Lueck, R., Beeli, C., “The Formation of Quasicrystalline Alloys from the

Melt in the Aluminium-Palladium-Manganese System”, Proc. 5th Int. Conf. Quasicryst.,


[1995Ish] Ishimasa, T., “Superlattice Ordering in the Lowe-Temperature Icosahedral Phase of

Al-Pd-Mn”, Philos. Mag. Lett., 71(1), 65-73 (1995) (Crys. Structure, *, 14)

[1996Bou] Boudard, M., Klein, H., de Boissieu, M., Audier, M., “Structure of Quasicyrstalline

Approximant Phase in the Al-Pd-Mn System”, Philos. Mag. A, 74(4), 939-956 (1996) (Crys.

Structure, *, 31)

[1996Kle] Klein, H., Audier, M., Boudard, M., De Boissieu, M., “Phason Defects in Al-Pd-Mn

Approximant Phases”, Philos. Mag. A, 73(2), 309-331 (1996) (Crys. Structure, 31)

[1996Sun] Sun, W., Hiraga, K., “High-Resolution Transmission Electron Microscopy of the Al-Pd-Mn

Decagonal Quasicrystal with 1-6nm Periodicity and its Crystalline Approximants”, Philos.

Mag. A, 73(4), 951-971 (1996) (Crys. Structure, Experimental, 19)

[1996Tan] Tanaka, K., Mitarai, Y., Koiwa, M., “Elastic Constants of Al-Based Icosahedral

Quasicrystals”, Philos. Mag. A, 73(6), 1715-1723 (1996) Crys. Structure, Mechan. Prop.,

Experimental, 18)

[1996Yam] Yamamoto, A., “Crystallography of Quasiperiodic Crystals”, Acta Crystallogr., Sect. A:

Found. Crystallogr., 52, 509-560 (1996) (Calculation, Crys. Structure, Review, 211)

[1997Ama] Amazit, Y., Perrin, B., Fischer, M., Itie, J.P., Polian, A., “X-Ray Diffraction Measurements

in Icosahedral Al-Pd-Mn up to 40 GPa”, Philos. Mag. A, 75(6), 1677-1688 (1997) (Crys.


226


MSIT®

Al–Mn–Pd

[1997Gie] Gierer, M., van Hove, M.A., Goldman, A.I., Shen, Z., Chang, S.-L., Jenks, C.J.,

Zhang, C.-M., Thiel, P.A., “Structural Analysis of the Fivefold Symmetric Surface of the

Al70Pd21Mn9 Quasicrystal by Low Energy Electron Diffraction”, Phys. Rev. Lett., 78(3),


[1997Hae] Häussler, D., Beeli, C., Nissen, H.-U., “One-Dimensionally Modulated Quasicrystal Phase

Related to Icosahedral Al-Mn-Pd”, Philos. Mag. Lett., 75(2), 117-124 (1997) (Crys.


[1997Ina] Inaba, A., Tsai, A.P., Shibata, K., “Vibrational Properties of Quasicrystals of Al-Cu-Ru,

Al-Pd-Re and Al-Pd-Mn Deduced from Heat Capacities”, Proc. of the 6th International

Conf. on Quasicrystals, Takeuchi, S., Fujiwara, T. (Eds.) (ICQ6), World Scientific,

Singapore, 1997, p.443-450 (Crys. Structure, Experimental, 10)

[1997Kle] Klein, H., Boudard, M., Audier, M., de Boissieu, M., Vincent, H., Beraha, L., Duneau, M.,

“The T-Al3(Mn, Pd) Quasicrystalline Approximant: Chemical Oreder and Phason Defect”,

Philos. Mag. A, 75(4), 197-208 (1997) (Crys. Structure, Experimental, *, 21)

[1997Kra] Krajci, M., Hafner, J., Mihalkovic, M., “Atomic and Electronic Structure of Decagonal

Al-Pd-Mn Alloys and Approximant Phases”, Phys. Rev. B, 55(2), 843-855 (1997) (Crys.


[1997Mat] Matsuo, Y., Kaneko, M., Yamanoi, T., Kaji, N., Sugiyama, K., Hiraga, K., “The Structure

of an Al3Mn-Type Al3(Mn, Pd) Crystal Studied by Single-Crystal X-Ray Diffraction

Analysis”, Philos. Mag. Lett., 76(5), 357-362 (1997) (Crys. Structure, Experimental, *, 9)

[1997Oka] Okamoto, H., “Al-Mn (Aluminum-Manganese)”, J. Phase Equilib., 18(4), 398-399 (1997)

(Crys. Structure, Equi. Diagram, Review, 8)

[1997She] Shen, Z., Jenks, C.J., Anderegg, J., Delaney, D.V., Kograsso, T.A., Thiel, P.A., Goldman,

A.I., “Structure and Stability of the Twofold Surface of Icosahedral Al-Pd-Mn by

Low-Energy Electron Diffraction and X-ray Photoemission Spectroscopy”, Phys. Rev.

Lett., 78(6), 1050-1053 (1997) (Crys. Structure, Experimental, 22)

[1997Son] Sonsky, J., Jelinek, M., Jastrabik, L., Studnicka, V., Chvostova, D., Bryknar, Z., “Study of

Quasicrystalline Thin Films Based on Al-Pd-Mn and Al-Cu-Fe Prepared by PLD”,

Czechoslov. J. Phys., 47(10), 1019-1024 (1997) (Crys. Structure, Electr. Prop.,

Experimental, 16)

[1997Zur] Zurkirch, M., Crescenzi, M.D., Erbudak, M., Hochstrasser, M., “Comparison of the

Structure of AlPd and Al70Pd20Mn10”, Phys. Rev. B, 55(14), 8808-8811 (1997) (Crys.


[1998Ber] Beraha, L., Duneau, M., Klein, H., Audier, M., “Phason Defects in Al-Pd-Mn Approximant

Phases: Another Example”, Philos. Mag. A, 78(2), 345-372 (1998) (Crys. Structure,


[1998Boi] de Boissieu, M., Boudard, M., Ishimasa, T., Elkaim, E., Lauriat, J.P., Letoublon, A.,

Audier, M., Duneau, M., Davroski, A., “Reversible Transformation Between an Icosahedral

Al-Pd-Mn Phase and a Modulated Structure of Cubic Symmetry”, Philos. Mag. A, 78(2),

305-326 (1998) (Crys. Structure, Experimental, *,39)

[1998Bol] Bolliger, B., Erbudak, M., Vvedensky, D.D., Zurkirch, M., “Surface Strcutural Transitions

on the Icosahedral Quasicrystal Al70Pd20Mn10”, Phys. Rev. Lett., 80(24), 5369-5372 (1998)


[1998Gie] Gierer, M., Van Hove, M.A., Goldman, A.I., Shen, Z., Chang, S.-L., Pinhero, P.J.,

Jenks, C.J., Anderegg, J.W., Zhang, C.-M., Thiel, P.A., “Fivefold Surface of

Quasicrystalline AlPdMn: Strcuture Determination Using Low-Energy-Electron

Diffraction”, Phys. Rev. B: Condens. Matter, 57(13), 7628-7641 (1998) (Crys. Structure,

Experimental, 55)

[1998Mat] Matsuo, Y., Yamamoto, Y., Ishii, Y., “Investigation of Phason Strains in Decagonal

Al-Pd-Mn Single Qasicryustals by Means of X-ray Diffraction”, J. Phys.: Condens. Matter.,


227


MSIT®

Al–Mn–Pd

[1998Sim] Simonet, V., Hippert, F., Klein, H., Audier, M., Bellissent, R., Fischer, H., Murani, A.P.,

Boursier, D., “Local Order and Magnetism in Liquid Al-Pd-Mn Alloys”, Phys. Rev. B,

58(10), 6273-6286 (1998) (Crys. Structure, Experimental, Magn. Prop., 63)

[1998Wae] Waelti, C., Felder, E., Chernikov, M.A., Ott, H.R., de Boissieu, M., Janot, C., “Lattice

Excitations in Icosahedral Al-Mn-Pd and Al-Re-Pd”, Phys. Rev. B: Condens. Matter,

57(17), 10504-10511 (1998) (Experimental, Thermodyn., 46)

[1998Wan] Wang, R., Feuerbacher, M., Yang, W., Urban, K., “Stacking Faults in

High-Temperature-Deformed Al-Pd-Mn Icosahedral Quasicrystals”, Philos. Mag. A, 78(2),

273-284 (1998) (Experimental, 19)

[1998Was] Waseda, Y., Suzuki, S., Urban, K., “Novel Morphology of Voids in Single-Quasicrystalline

Icosahedral Al70.5Pd21.0Mn8.5”, Z. Naturforsch. A, 53A, 679-683 (1998) (Crys. Structure,

Experimental, 16)

[1999Aud] Audier, M., Duneau, M., de Boissieu, M., Boudard, M., Letoublon, A., “Superlattice

Ordering of Cubic Symmetry in an Icosahedral Al-Pd-Mn Phase”, Philos. Mag. A, 79(2),

255-270 (1999) (Crys. Structure, Experimental, *, 9)

[1999Cap] Capitan, M.J., Calvayrac, Y., Quivy, A., Joulaud, J.L., Lefebvre, S., Gratias, D. “X-Ray

Diffuse Scattering from Icosahedral Al-Pd-Mn Quasicrystals”, Phys. Rev. B, 60(9),


[1999Fis] Fisher, I.R., Kramer, M.J., Wiener, T.A., Islam, Z., Ross, A.R., Lograsso, T.A., Kracher, A.,

Goldman, A.I., Canfield, P.C., “On the Growth of Icosahedral Al-Pd-Mn Quasicrystals

from the Ternary Melt”, Philos. Mag. B, 79(10), 1673-1684 (1999) (Experimental, Phys.

Prop., 20)

[1999Gru] Grushko, B., Yurechko, M., Tamura, N., “A Contribution to the Al-Pd-Mn Phase Diagram”,

J. Alloys Compd., 290, 164-171 (1999) (Equi. Diagram, Experimental, #, *, 24)

[1999Hip] Hippert, F., Simonet, V., Trambly de Laissardiere, G., Audier, M., Calvayarac, Y.,

“Magnetic Properties of AlPdMn Appximant Phases”, J. Phys.: Condens. Matter, 11,

10419-10435 (1999) (Crys. Structure, Experimental, Magn. Prop., 48)

[1999Yok] Yokoyama, Y., Yamada, Y., Fukaura, K., Sunada, H., Note, R., Inoue, A., Sugiyama, K.,

Hiraga, K., “Strain Affected Properties of Icosahedral Al-Pd-Mn Single Ingot”, Jpn. J. Appl.

Phys., 38(1)(3A), 1495-1499 (1999) (Crys. Structure, Equi. Diagram, Electr. Prop., Magn.

Prop., Experimental, *, 16)

[2000Bai] Baier, F., Mueller, M.A., Grushko, B., Schaefer, H.-E., “Atomic Defects in Quasicrystals:

an Approach with Positron Annihilation Spectroscopy and Time-Differential Dilatometry”,

Mater. Sci. Eng. A, 294-296, 650-653 (2000) (Crys. Structure, Experimental, 13)

[2000Bar] Bartsch, M., Geyer, B., Haeussler, D., Feuerbacher, M., Urban, K., Masserschmidt, U.,

“Plastic Properties of Icosahedral Al-Pd-Mn Single Quasicrystals”, Mater. Sci. Eng. A,

294-296, 761-764 (2000) (Crys. Structure, Experimental, Phys. Prop., 14)

[2000Bee1] Beeli, C., Soltmann, C., Poon, S.J., “Relationship of Phason Strain and Electronic Properties

in Icosahedral Al-Pd-(Re,Mn) and Al-Cu-Os”, Mater. Sci. Eng. A, 294-296, 531-534 (2000)


[2000Bee2] Beeli, C., “High-Resolution Electron Microscopy of Quasicrystals”, Mater. Sci. Eng. A,

294-296, 23-28 (2000) (Crys. Structure, Experimental, 40)

[2000Bel] Bellissard, J., “Anomalous Transport: Results, Conjectures and Applications to

Quasicrystals”, Mater. Sci. Eng. A, 294-296, 450-457 (2000) (Crys. Structure, Phys. Prop.,

Review, 56)

[2000Blu] Blueher, R., Frank, W., Grushko, B., “Diffusion of 103Pd and 195Au in Icosahedral

Al70.2Pd21.3Mn8.5 under Proton Irradiation”, Mater. Sci. Eng. A, 294-296, 689-692 (2000)

(Crys. Structure, Experimental, Phys. Prop., 10)

[2000Bol] Bolliger, B., Erbudak, M., Hensch, A., Vvedensky, D.D., “Surface Structural Phase

Transitions on Icosahedral Al-Pd-Mn”, Mater. Sci. Eng. A, 294-296, 859-862 (2000) (Crys.


228


MSIT®

Al–Mn–Pd

[2000Bru1] Brunet, P., Zhang, L.M., Sordelet, D.J., Besser, M., Dubois, J-M., “Comparative Study of

Microstructural and Tribological Properties of Sintered, Bulk Icosahedral Samples”, Mater.

Sci. Eng. A, 294-296, 74-78 (2000) (Mechan. Prop., Experimental, 6)

[2000Bru2] Brunner, D., Plachke, D., Carstanjen, H.D., “The Strain-Softering Phenomenon of

Icosahedral Al-Pd-Mn Single Quasicrystals”, Mater. Sci. Eng. A, 294-296, 773-776 (2000)

(Mechan. Prop., Experimental, 19)

[2000Cap1] Cappello, G., Schmithuesen, F., Chevrier, J., Comin, F., Stierle, A., Formoso, V., Boissieu,

M., Boudard, M., Lograsso, T.A., Jenks, C., Delaney, D., “Characterization of Surface

Morphologies at the Al-Pd-Mn Fivefold Surface”, Mater. Sci. Eng. A, 294-296, 822-825


[2000Cap2] Cappello, G., Dechelette, A., Schmithuesen, F., Decossas, S., Chevrier, J., Comin, F.,

Formoso, V., Boissieu, M., Jach, T., Colella, R., Lograsso, T.A., Jenks, C., Delaney, D.,

“Bulk and Surface Evidence for the Long-Range Spatial Modulation of X-Ray Absorption

in the Al-Pd-Mn Quasicrystal at Bragg Incidence”, Mater. Sci. Eng. A, 294-296, 863-866


[2000Cyr] Cyron-Lackmann, F., “Quasicrystals as Potential for Thermoelectric Materials”, Mater. Sci.

Eng. A, 294-296, 611-612 (2000) (Calculation, Crys. Structure, Electr. Prop., Phys. Prop.,

Thermal Conduct., 15)

[2000Dam] Damson, B., Weller, M., Feuerbacher, M., Grushko, B., Urban, K., “Mechanical

Spectroscopy of i-Al-Pd-Mn and d-Al-Ni-Co”, Mater. Sci. Eng. A, 294-296, 806-809 (2000)


[2000Dub] Dubois, J-M., “New Prospects from Potential Applications of Quasicrystalline Materials”,

Mater. Sci. Eng. A, 294-296, 4-9 (2000) (Crys. Structure, Experimental, Phys. Prop.,

Review, 38)

[2000Dun] Duneau, M., “Covering Clusters in the Katz-Gratias Model of Icosahedral Quasicrystals”,

Mater. Sci. Eng. A, 294-296, 192-198 (2000) (Calculation, Crys. Structure,

Experimental, 34)

[2000Eda] Edagawa, K., Kajiyama, K., “High Temperature Specific Heat of Al-Pd-Mn and Al-Cu-Co

Quasicrystals”, Mater. Sci. Eng. A, 294-296, 646-649 (2000) (Crys. Structure, Thermodyn.,


[2000Feu] Feuerbacher, M., Klein, H., Bartsch, M., Messerschmidt, U., Urban, K., “A Comparative

Study of the Plastic Behavior of Icosahedral and pri-Al-Pd-Mn”, Mater. Sci. Eng. A,

294-296, 736-741 (2000) (Crys. Structure, Experimental, Phys. Prop., 23)

[2000Fis] Fisher, I.R., Kramer, M.J., Islam, Z., Wiener, T.A., Kracher, A., Ross, A.R., Lograsso, T.A.,

Goldman, A.I., Canfield, P.C., “Growth of Large Single-Grain Quasicrystals from

High-Temperature Metallic Solutions”, Mater. Sci. Eng. A, 294-296, 10-16 (2000) (Crys.


[2000Fra] Fradkin, M.A., “Finite-Resolution Correction to the Diffraction Intensity in Icosahedral

Quasicrystals”, Mater. Sci. Eng. A, 294-296, 319-322 (2000) (Calculation, Crys. Structure,

Experimental, 4)

[2000Fre] Frey, F., “Disorder Diffuse Scattering of Decagonal Quasicrystals”, Mater. Sci. Eng. A,

294-296, 178-185 (2000) (Crys. Structure, Experimental, 15)

[2000Gal] Galler, R., Mehrer, H., “Diffusion in Icosahedral Al-Pd-Mn Quasicrystals: Temperature and

Pressure Dependence”, Mater. Sci. Eng. A, 294-296, 693-696 (2000) (Crys. Structure,

Experimental, Phys. Prop., 17)

[2000Gru] Grushko, B., “Composition and Presipitation Behavior of Icosahedral Al-Pd-Mn

Quasicrystals”, Mater. Sci. Eng. A, 294-296, 45-48 (2000) (Crys. Structure, Experimental,

*, #, 16)

[2000Gwo] Gwozdz, J., Grushko, B., Surowiec, M., “Mosaic Structure of Single Al-Pd-Mn Icosahedral

Quasi-Crystals”, Mater. Sci. Eng. A, 294-296, 49-52 (2000) (Crys. Structure,

Experimental, 7)

229


MSIT®

Al–Mn–Pd

[2000Hir] Hirai, I., Ishimasa, T., Letoublon, A., Boudard, M., Boissieu, M., “Formation Conditions of

Two Quasiperiodic Modifications of Al-Pd-Mn Icosahedral Phase Studied by Annealing

Method”, Mater. Sci. Eng. A, 294-296, 33-36 (2000) (Crys. Structure, Experimental, *, 9)

[2000Jac] Jach, T., “Quasicrystal Element Correlations from X-Ray Standing Waves”, Mater. Sci.

Eng. A, 294-296, 315-318 (2000) (Calculation, Crys. Structure, Theory, 13)

[2000Jan] Janot, C., Loreto, L., Farinato, R., “Clusters in Quasicrystals: Tiling Versus Covering and

Porosity”, Mater. Sci. Eng. A, 294-296, 405-408 (2000) (Assessment, Crys. Structure, 19)

[2000Kaj] Kajiyama, K., Edagawa, K., Suzuki, T., Takeuchi, S., “Thermal Expansion of Icosahedral

Al-Pd-Mn and Decagonal Al-Cu-Co Quasicrystals”, Philos. Mag. Lett., 80(1), 49-56 (2000)

(Crys. Structure, Experimental, Mechan. Prop., 19)

[2000Kle1] Klein, H., Durand-Charre, M., Audier, M., “Liquid-Solid Equilibria in the Quasicrystalline

Regions of the Al-Pd-Mn Phase Diagram”, J. Alloys Compd., 296, 128-137 (2000) (Equi.


[2000Kle2] Klein, H., Feuerbacher, M., Schall, P., Urban, K., “Bending Experiments on the

´-(Al-Pd-Mn) Quasicrystal Approximant”, Philos. Mag. Lett., 80(1), 11-18 (2000) (Crys.


[2000Kle3] Klein, H., Feuerbacher, M., Urban, K., “Dislocation in Al-Pd-Mn Approximants: a High

Resolution Electron Microscopy Study”, Mater. Sci. Eng. A, 294-296, 769-772 (2000)


[2000Klu] Kluge, F., Ebert, P., Grushko, B., Urban, K., “Influence of Grown-in Voids on the Structure

of Cleaved Icosahedral Al-Pd-Mn Quasicrystal Surfaces”, Mater. Sci. Eng. A, 294-296,


[2000Lai] Laissardiere, G.T., Mayou, D., “Conditions on the Occuerence of Magnetic Moments in

Quasicrystals and Related Phases”, Mater. Sci. Eng. A, 294-296, 621-624 (2000)

(Calculation, Crys. Structure, Magn. Prop., 18)

[2000Lan] Lang, C.I., Sordelet, D.J., Besser, M.F., Shechtman, D., Biancaniello, F.S., Gonzales, E.J.,

“Quasicrystalline Coatings: Thermal Evolution of Structure and Properties”, J. Mater. Res.,

15(9), 1894-1904 (2000) (Experimental, Mechan. Prop., Phys. Prop., 41)

[2000Led] Ledieu, J., Muryn, C.A., Thornton, G., Cappello, G., Chevrier, J., Diehl, R.D.,

Lograsso, T.A., Delaney, D., McGrath, R., “Decomposition of the Five-Fold Surface of

Al70Pd21Mn9 at Elevated Temperature”, Mater. Sci. Eng. A, 294-296, 871-873 (2000)


[2000Let] Letoublon, A., Ishimasa, T., de Boissieu, M., Boudard, M., Hennion, B., Mori, M.,

“Stability of the F2-(Al-Pd-Mn) Phase”, Philos. Mag. Lett., 80(4), 205-213 (2000) (Crys.

Structure, Equi. Diagram, Experimental, *, 14)

[2000Man] Mancini, L., Letoublon, A., Agliozzo, S., Wang, J., Gastaldi, J., Boissieu, M., Haertwig, J.,

Klein, H., “Effect of Annealing on the Structural Perfection of Al-Pd-Mn Icosahedral

Quasicrystal Grains”, Mater. Sci. Eng. A, 294-296, 57-60 (2000) (Crys. Structure,

Experimental, 20)

[2000Mes] Messerschmidt, U., Haeussler, D., Bartsch, M., Geyer, B., Feuerbacher, M., Urban, K.,

“Microprocesses of the Plastic Deformation of Icosahedral Al-Pd-Mn Single

Quasicrystals”, Mater. Sci. Eng. A, 294-296, 757-760 (2000) (Crys. Structure,

Experimental, 15)

[2000Nau] Naumovic, D., Aebi, P., Schlapbach, L., Beeli, C., “Atomic and Electronic Structure of

Five-Fold i-Al-Pd-Mn Surfaces”, Mater. Sci. Eng. A, 294-296, 882-885 (2000) (Crys.

Structure, Experimental, Phys. Prop., 28)

[2000Nic] Nicula, R., Jianu, A., Grigoriu, C., Barfels, T., Burkel, E., “Laser Ablation Synthesis of

Al-Based Icosahedral Powders”, Mater. Sci. Eng. A, 294-296, 86-89 (2000) (Crys.

Structure, Experimental, Mechan. Prop., 12)

[2000Qua] Quandt, A., Elser, V., Kresse, G., Hafner, J., “An Ab Initio Based Structure Model of

i(Al-Pd-Mn)”, Mater. Sci. Eng. A, 294-296, 351-354 (2000) (Calculation, Crys.

Structure, 19)

230


MSIT®

Al–Mn–Pd

[2000Ros] Ross, A.R., Wiener, T.A., Fisher, I.R., Canfield, P.C., Lograsso, T.A., “Formation and

Morphological Development of Porosity in Icosahedral Al-Pd-Mn Alloys”, Mater. Sci.

Eng. A, 294-296, 53-56 (2000) (Crys. Structure, Experimental, 11)

[2000Sch1] Schaub, T., Delahaye, J., Berger, C., Grenet, T., Guyot, H., Belkhou, R., Taleb-Ibrahimi, A.,

Prejean, J.J., Calvayrac, Y., “High Resolution Experiment on the Electronic Density of

States in Icosahedral-Al-Pd-Mn”, Mater. Sci. Eng. A, 294-296, 512-515 (2000) (Crys.


[2000Sch2] Scheffer, M., Suck, J.-B., “Influence of Vacancies on the Magnetic Properties of

Icosahedral Al7.10Pd20.0Mn9.0”, Mater. Sci. Eng. A, 294-296, 629-632 (2000) (Crys.

Structure, Experimental, Magn. Prop., 14)

[2000Sch3] Schurack, F., Eckert, J., Schultz, L., “Quasicrystalline Al-Alloys with High Strength and

Good Ductility”, Mater. Sci. Eng. A, 294-296, 164-167 (2000) (Crys. Structure,

Experimental, Mechan. Prop., 6)

[2000Sch4] Schall, P., Feuerbacher, M., Bartsch, M., Messerschmidt, U., Urban, K., “Dislocation

Arrangement and Density in Deformed Al-Pd-Mn Single-Quasicrystals”, Mater. Sci.

Eng. A, 294-296, 765-768 (2000) (Calculation, Crys. Structure, Experimental, 9)

[2000Sch5] Schmithuesen, F., Boissieu, M., Boudard, M., Chevrier, J., Comin, F., “Electron Energy

Loss Spectroscopy Investigation of Volume and Surface Plasmonts at the Al-Pd-Mn

Fivefold Surface”, Mater. Sci. Eng. A, 294-296, 867-870 (2000) (Crys. Structure,

Experimental, 15)

[2000Shr] Shramchenko, N., Klein, H., Caudron, R., Bellissent, R., “Comparison of Local Order in

Icosahedral Al-Pd-Mn Quasicrystal and in Approximant Phase by Thermal Neutron

Scattering”, Mater. Sci. Eng. A, 294-296, 335-339 (2000) (Crys. Structure,

Experimental, 10)

[2000Sim] Simonet, V., Hippert, F., Audier, M., Calvayras, Y., “Magnetism of Approximants in the

Al-Mn and Al-Pd-Mn Systems”, Mater. Sci. Eng. A, 294-296, 625-628 (2000) (Crys.

Structure, Experimental, Magn. Prop., 24)

[2000Sta] Stadnic, Z.M., “Photoemission Studies of Qusicrystals”, Mater. Sci. Eng. A, 294-296,

470-474 (2000) (Crys. Structure, Electr. Prop., Experimental, 26)

[2000Ste1] Steurer, W., “The Quasicrystal-to-Crystal Transformation. I. Geometrical Principles”,

Z. Kristallogr., 215, 323-334 (2000) (Calculation, Crys. Structure, 44)

[2000Ste2] Steurer, W., “Geometry of Quasicrystal-to-Crystal Transformations”, Mater. Sci. Eng. A,

294-296, 268-271 (2000) (Assessment, Crys. Structure, 12)

[2000Tho] Thompson, E., Vu, P.D., Pohl, R.O., “Glasslike Lattice Vibrations in the Quasicrystal

Al72.1Pd20.7Mn7.2”, Phys. Rev. B, 62(17), 11437-11443 (2000) (Crys. Structure,

Experimental, Phys. Prop., Thermal Conduct., 49)

[2000Uch] Uchiyama, H., Takahashi, Y., Sato, K., Kanazawa, I., Kimura, K., Komori, F., Suzuki, R.,

Ohdaira, T., Tamura, R., Takeuchi, S., “Stable Quasicrystals Studied by Means of the Slow

Positron Beam”, Nucl. Instrum. Methods Phys. Res./B, B171, 245-250 (2000) (Crys.


[2000Wan] Wang, R., Yang, W., Gui, J., Urban, K., “Dislocation Mechanism of High-Temperature

Plastic Deformation of Al-Cu-Fe and Al-Pd-Mn Icosahedral Quasicrystals”, Mater. Sci.

Eng. A, 294-296, 742-747 (2000) (Crys. Structure, Experimental, 18)

[2000Weh] Wehner, B.I., Koster, U., Rudiger, A., Pieper, A., Sordelet, D.J., “Oxidation of Al-Cu-Fe

and Al-Pd-Mn Quasicrystals”, Mater. Sci. Eng. A, 294-296, 830-833 (2000) (Crys.


[2000Yam] Yamamoto, A., Hiraga, K., “Six-Dimensional Model of an i-Al-Pd-Mn Quasicrystal

Compatible with its 2/1 Approximant”, Mater. Sci. Eng. A, 294-296, 228-231 (2000) (Crys.

Structure, Experimental, Review, 3)

[2000Zum] Zumkley, T., Guo, J.Q., Tsai, A.P., Nakajima, H., “Diffusion in Quasicrystalline Al-Ni-Co

and Al-Pd-Mn”, Mater. Sci. Eng. A, 294-296, 702-705 (2000) (Crys. Structure,


231


MSIT®

Al–Mn–Pd

[2001Can] Canfield, P.C., Fisher, R., “Hihg-Temperature Solution Growth of Intermetallic Single

Crystal and Quasicrystal”, J. Cryst. Growth, 225, 155-161 (2001) (Crys. Structure,


[2001Goe] Gödecke, T., “Ableitung des Kristallizationpfades in Ternaeren Gusslegierungen” (in

German), Z. Metallkd., 92(8), 966-978 (2001) (Equi. Diagram, Experimental, *, 37)

[2001Nau] Naumovic, D., Aebi, P., Schlapbach, L., Beeli, C., Kunze, K., Lograsso, T.A.,

Delaney, D.W., “Formation of Stable Decagonal Quasicrystalline Al-Pd-Mn Surface

Layer”, Phys. Rev. Lett., 87(19), 195506-1-195506-4 (2001) (Crys. Structure,

Experimental, 35)

[2002Ant] Anton, R., Kreutzer, P., “Growth and Electrical and Optical Properties of Al(PdMn) Alloy

Thin Films Prodused by Simultaneous Vapor Deposition of the Components”, J. Alloys

Compd., 342(1-2), 464-468 (2002) (Crys. Structure, Equi. Diagram, Experimental 10)

[2002Bil] Bilusic, A., Bodrovic, Z., Smontara, A., Dolinsek, J., Canfield, P.C., Fisher, I.R., “Transport

Properties of Icosahedral Quasicrystal Al72Pd19.5Mn8.5”, J. Alloys Compd., 342(1-2),

413-415 (2002) (Electr. Prop., Experimental, Phys. Prop., Thermal Conduct., 23)

[2002Dem] Demange, V., Milandri, A., Weerd, M.C., Machizaud, F., Jeandel, G., Dubois, J.M.,

“Optical Conductivity of Al-Cr-Fe Approximant Compounds”, Phys. Rev. B, 65, 144205-1

-144205-11 (2002) (Calculation, Crys. Structure, Experimental, Optical Prop., 39)

[2002Duq] Duquesne, J.-Y., Perrin, B., “Elastic Wave Interaction in Icosahedral AlPdMn”, Physica B,

316-317, 317-320 (2002) (Experimental, Mechan. Prop., 9)

[2002Hir] Hiraga, K., “The Structure of Quasicrystals Studied by Atomic-Scale Observations of

Transmission Electron Microscopy”, Adv. Imag. Electr. Phys., 122, 1-86 (2002)

(Assessment, Crys. Structure, 99)

[2002Kab] Kabutoya, E., Edagawa, K., Tamura, R., Takeuchi, S., Guo, J.Q., Tsai, A.-P., “Plastic

Deformation of Icosahedral Al-Pd-Mn Single Qusicrystals to large Strains I. Experiments”,

Philos. Mag. A, 82(2), 369-377 (2002) (Experimental, Mechan. Prop., 15)

[2002Klu] Kluge, F., Yurechko, M., Urban, K., Ebert, Ph., “Influence of Growth Kinetics and

Chemical Composition on the Shape of Voids in Quasi-Crystal”, Surf. Sci., 519, 33-39


[2002Lei] Lei, J., Wang, R., Yin, J., Duan, X., “Diffuse Electron Scattering Determination of Elastic

Constants of Al-Pd-Mn Icosahedral Quasicrystal”, J. Alloys Compd., 342(1-2), 326-329

(2002) (Experimental, Mechan. Prop., 8)

[2002Miz] Mizutani, T., Nakano, H., Kashimoto, S., Takatani, Y., Mori, M., Ishimasa, T., Matsuo, S.,

“Ten-Fold-Like Magnetic Anisotropy in Electrical Conductivity of Al-Pd-Mn Icosahedral

Quasicrystal”, J. Alloys Compd., 342(1-2), 360-364 (2002) (Crys. Structure, Electr. Prop.,


[2002Mot] Motomura, S., Ishimasa, T., Hirai, I., Kashimoto, S., Nakano, H., Matsuo, S., “Magnetic

Properties of F2M-Type Al-Pd-Mn Quasicrystals”, J. Alloys Compd., 342(1-2), 393-396

(2002) (Crys. Structure, Experimental, Magn. Prop., 14)

[2002Pap] Papadopolos, Z., Kasner, G., Ledieu, J., Cox, E. J., Richardson, N.V., Chen, E. J.,

Diehl, R.D., Lograsso, T. A., Ross, A. R., McGrath, R., “Bulk Termination of the

Quasicrystalline Fivefold Surface of Al70Pd21Mn9”, Phys. Rev. B, 66(18),

184207-1-184207-13 (2002) (Crys. Structure, Experimental, 47)

[2002Sch] Scheffer, M., Suck, J-B., “Inelastic Neutron Scattreing Study of the Dynamics of

Al74Pd22Mn4 ( `)”, J. Alloys Compd., 342, 310-313 (2002) (Calculation, Experimental,

Thermodyn., 13)

[2002Shr] Shramchenko, N., Denoyer, F., “The Al-Pd-Mn Quasicrystalline Approximant ( )`-Phase

Revisited”, Eur. Phys. J. B, 29(1), 51-59 (2002) (Crys. Structure, Experimental, *, 27)

[2002Tak] Takeuchi, S., Tamura, R., Kabutoya, E., Edagawa, K., “Plastic Deformation of Icosahedral

Al-Pd-Mn Single Quasicrystals to Large Strains II. Deformation Mechanism”, Philos.

Mag. A, 82(2), 379-385 (2002) (Experimental, Mechan. Prop., 9)

232


MSIT®

Al–Mn–Pd

[2002Tex] Texier, M., Proult, A., Bonneville, J., Rabier, J., Baluc, N., Cordiers, P., “Microstructure of

Icosahedral Al-Pd-Mn Quasicrystals Deformed at Room Temperature in an Anisotropic

Confining Medium”, Philos. Mag. Lett., 82(12), 659-669 (2002) (Crys. Structure,


[2002Yam] Yamamoto, A., Takakura, H., Tsai, A.P., “Structure Refinement of i-Al-Pd-Mn

Quasicrystals by IP-Weissenberg Camera Data”, J. Alloys Compd., 342, 159-163 (2002)


[2002Yan] Yang, W., Feuerbacher, M., Urban, K., “Cluster Structure and Low-Energy Planes in

Icosahedral Al-Pd-Mn Quasicrystals”, J. Alloys Compd., 342(1-2), 164-168 (2002) (Crys.


[2002Yur] Yurechko, M., “Phase Equilibria in Ternary Systems on the Aluminum Basis, which

Contain Quasiperiodic and Related Periodic Phases”, PhD Thesis, The Taras Shevchenko

Kiev National University, Kiev (2002) (Equi. Diagram, Experimental, #, *, 160)

[2002Zha1] Zhang, Y., Colella, R., Kycia, S., Goldman, A.I., “Absolute Structure-Factor Measurements

of an Al-Pd-Mn Quasicrystal”, Acta Crystallogr., Sect. A: Found. Crystallogr., 58, 385-390


[2002Zha2] Zhang, Y., Ehrlich, S.N., Colella, R., Kopecky, M., Widom, M., “X-Ray Diffuse Scattering

in the Icosahedral Quasicrystal Al-Pd-Mn”, Phys. Rev. B, 66, 104202-1-104202-7 (2002)

(Crys. Structure, Experimental, Theory, 18)

[2003Ban] Banerjee, G.N., Banerjee, S., Goswami, R., “Point Contact Spectroscopy of Al70Pd30-xMnx

Quasicrystals”, J. Phys.: Condens. Matter, 15(14), 2317-26 (2003) (Crys. Structure,

Experimental, Electr. Prop., 22)

[2003Bal] Balanetskyy, S., Grushko, B., “Al-Pd (Aluminium - Palladium)”, MSIT Evaluation


Services GmbH, Stuttgart, to be published, (2003) (Equi. Diagram, Crys. Structure,

Assessment, 26)

[2003Cap] Capitan, M.J., Alvarez, J., Prejean, J.J., Berger, C., “Conductivity and Superlattice Ordering

in an Icosahedral Al-Pd-Mn Phase”, Phys. Rev. B, 68(6), 064203-1-9 (2003), (Crys.

Structure, Experimental, Electr. Prop., 32)

[2003Ebe] Ebert, Ph., Yurechko, M., Kluge, F., Cai, T., Grushko, B., Thiel, P.A., Urban K., “Surface

Structure of Al-Pd-Mn Quasicrystals: Existence of Supersaturated Bulk Vacancy

Concentrations”, Phys. Rev. B: Condens. Matter, 67(2), 24208-1-8 (2003) (Crys. Structure,


[2003Pis] Pisch, A., “Al-Mn (Aluminium-Manganese)”, MSIT Evaluation Program, in MSIT


Stuttgart, to be published, (2003) (Equi. Diagram, Crys. Structure, Assessment, 40)


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

( Al)

< 660.452

cF4

Fm3m

Cu

a = 404.96 at 25°C [Mas2]

dissolves ~0.2 at.% Pd [2003Bal]

and ~0.62 at.% Mn [1997Oka]

( Mn)

1246 - 1140

cI2

Im3m

W

a = 308.0 [Mas2]

dissolves ~39.5 at.% Al [1997Oka]

and ~4 at.% Pd [Mas2]

233


MSIT®

Al–Mn–Pd

( Mn) cF4

Fm3m

Cu

a = 386.0 [Mas2]

dissolves ~9.1 at.% Al at 1073°C

[1997Oka]

and ~22.6 at.% Pd at 1147°C [Mas2]

( Mn)

1100 - 727

cP20

P4132

Mn

a = 631.52 [Mas2]

dissolves ~40.5 at.% Al [1997Oka]


( Mn)

< 727

cI58

I43m

Mn

a = 891.26 at 25°C [Mas2]

dissolves ~2 at.% Al [1997Oka]


(Pd)

< 1555

cF4

Fm3m

Cu

a = 389.0 [Mas2]

dissolves ~20 at.% Al at 1055°C

[2003Pis]

and ~30.5 at.% Mn at 1350°C [Mas2]

, (MnPdAl)

1400 - 840

, MnAl

1177 - 840

cI2

Im3m

W

a = 308.3 (?)

a = 306.3 0.3

MnPd2Al at >1010°C [1981Sol2]

34.5 to 51.3 at.% Mn [1997Oka]

Mn45Al55 at 957°C [1990Ell]

1, MnxAl1-x

1048 - 957

30 to 38.2 at.% Mn [1997Oka]

2, Mn5Al8< 991

hR26

R3m

Cr5Al8

a = 1273.9

c = 1586.1

31.4 to 50.01 at.% Mn [2003Pis]

at 42 at.% Mn [V-C2]

, MnxAl1-x

1275 - 870

hP2

P63/mmc

Mg

a = 270.5 - 270.5

c = 436.1 - 438

53.2 to 60 at.% Mn [1997Oka]

usually called

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

234


MSIT®

Al–Mn–Pd

H, (MnPdAl)4(MnPdAl)11

Mn4Al11 (h)

1002 - 895 oP160

Pnma

MnAl3a = 1479

b = 1242

c = 1259

a = 1483

b = 1243

c = 1251

a = 1483.7 0.4

b = 1245.7 0.2

c = 1250.5 0.2

a = 1471.7

b = 1251.0

c = 1259.4

a = 1476

b = 1243

c = 1256

a = 1417

b = 1251

c = 1259.4

a = 1472.7 0.3

b = 1250.9 0.3

c = 1260.0 0.3

Labeled as “H” by [1995Goe1,

1995Goe2, 1999Gru, 2002Yur],

6 to 7 at.% Pd, 71 to 73 at.% Al at

880-870°C

7 at.% Pd, 70.6 to 71.6 at.% Al at

850°C

~6.5 at.% Pd, 73.5 to 76.8 at.% Al at

710°C [2002Yur]

Labeled as “T” by [1993Aud,

1994Aud, 1997Kle, 1997Mat,

2000Kle1, 2002Tex]

3.5 to 6.6 at.% Pd, 16 to 25.5 at.% Mn

[2000Kle1]

25 to 28.7 at.% Mn [1997Oka]

MnAl3 [1992Li]

MnAl3, as-cast [1993Hir2]

MnAl3 (Pn21a?) [1994Shi]

Mn24.5Pd3.2Al72.3 single crystal

obtained by Bridgman technique

[1997Kle].

Mn23Pd6Al71, obtained by Bridgman or

Czochralski technique [1993Aud,

1994Aud]

in Mn11.85Pd21.88Al66.27 rapid

quenched sample, together with 3 and

(Al-Pd) [2002Tex]

Mn20.9Pd4.4Al74.7 single crystal,

as-cast [1997Mat]

Mn4Al11 (l)

< 916

aP30

P1

Mn4Al11

a = 509.5 0.4

b = 887.9 0.8

c = 505.1 0.4

= 89.35 0.04°

= 100.47 0.05°

= 105.08 0.06°

25 to 28.7 at.% Mn [1997Oka]

[V-C2]

, MnAl4< 923

hP574

P63/mmc

MnAl4

a = 1998

c = 2467.3

19 to 20.8 at.% Mn [1997Oka]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

235


MSIT®

Al–Mn–Pd

, MnAl4

< 693

hP586

P63/m a = 2838.2

c = 1238.9

16.8 to19 at.% Mn [Mas2]

[2003Pis]

space group does not fit 100%,

probably P63

,

(Mn,Pd,Al)1(Mn,Pd,Al)4

, MnAl4

orthorhombic

Cmcm

Mn11Ni4Al60

a = 2360

b = 1240

c = 770

a = 2388

b = 1243

c = 776

a = 2388

b = 1243

c = 778

Phase of composition 3.5-6.6 at.% Pd,

16-25.5 at.% Mn [2000Kle1], labeled

as “R” by [1993Aud, 1994Aud,

1998Ber, 2000Kle1]

in rapidly solidified MnAl4 alloy after

heating to 600°C [1992Li], metastable

Mn15.6Pd5.7Al78.6 (EPMA), in samples

obtained by induction melting and

Bridgman technique [1993Aud,

1994Aud]

Mn15.7Pd5.7Al78.6 (EPMA) [1998Ber]

MnAl6< 705

oC28

Cmcm

MnAl6

a = 755.51

b = 649.94

c = 887.24

a = 754.5 0.2

b = 649.0 0.3

c = 868.1 0.2

14.2 at.% Mn [1997Oka]

[V-C2]

[2000Yam]

MnAl12

< 500

cI26

Im3

WAl12

a = 747

7.7 at.% Mn [1997Oka],

[V-C2]

, Mn3Al10

< 860

hP26

P63/mmc

Co2Al5

a = 754.6 0.3

c = 289.5 0.2

[2003Pis],

metastable

MnxAl1-x tP2

P4/mmm

CuAu

a = 278 - 279

c = 356 - 357

55.8 to 55.1 at.% Al,

metastable [2003Pis]

i-MnAl icosahedral

m35

~20 at.% Mn [2003Pis],

quasicrystal, metastable

d-MnAl decagonal

D3

a = 1240.0 ~22 at.% Mn, [2003Pis]

quasicrystal, metastable

PdAl4< 604 P6322

PtAl4

a = 1308.6

c = 963.1

20 at.% Pd [2003Bal],

usually labeled as “ ”

Pd8Al21

640

tI116

I41/a

Pt8Al21

a = 1299.8

c = 1072.9

27.5 at.% Pd [2003Bal],


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

236


MSIT®

Al–Mn–Pd

“ -phase”

6

782 - 579

28

792 - ?

22

34

x

6*

orthorhombic

Pnma

Pnma

Pnma or P21ma

orthorhombic

B2mm

-

orthorhombic

orthorhombic

orthorhombic

B2221 ?

orthorhombic

a = 2350.0

b = 1680.0

c = 1230.0

a = 2354.1

b = 1656.6

c = 1233.9

a = 2354.3 0.5

b = 1664.3 0.5

c = 1237.4 0.4

a = 2350.0

b = 1680.0

c = 5700.0

a = 2357.4

b = 1661.0

c = 5712.0

a = 2350

b = 1680

c = 4490

a = 2350

b = 1680

c = 7010

a = 2353.1

b = 1658.0

c = 8183.2

a = 2354.7 0.5

b = 1670 3

c = 8261 37

a = (4700.0)

b = 3360.0

c = 2460.0

Overall ternary extension of the binary

phases,

0 to 5 at.% Mn at 710°C [2002Yur]

~25.4 to 26.9 at.% Pd [2003Bal]

phase labeled as ´ by [1996Bou,

2000Kle1, 2002Shr], composition

Mn5.0-4.6Pd22.1-22.4Al73.3 at 730°C

[2000Kle1]

Mn4.1Pd22.4Al73.5, single-crystal

obtained by Bridgman method, XRD

[1996Bou]

Mn4.5Pd22.9Al72.6 sample obtained by

Bridgman method [2002Shr]

28.1 to < 26.9 at.% Pd [2003Bal]


Bridgman method, labeled as ´_3

[2002Shr]

3.1 to 4.6 at.% Mn at ~750-~800°C

[2002Yur],

according to [2002Shr] could be named

´_2

1.2 to 3.1 at.% Mn at 710-790°C

[2002Yur], according to [2002Shr]

could be named ´_4

Mn4.5Pd22.9Al72.6, labeled as ´_5, c =

(c ´ (5+ )) [2002Shr]

single crystal? [2002Shr]

ordered 6 [2003Bal]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

237


MSIT®

Al–Mn–Pd

orthorhombic

Cmcm

Bmmb

a = 2032

b = 1650

c = 1476

a = 2032

b = 1657

c = 1475

a = c ´

b = b ´

c = -1 a ´

a = 2032

b = 1650

c = 1476

Mn4Pd23Al73, Mn4Pd21.6Al74 (EMPA)

[1993Aud]

Mn4Pd23Al73 [1998Ber]

[2000Kle1]

in as-cast Mn5Pd20Al75 alloy

[1996Sun]

,

(Mn,Pd,Al)2(Mn,Pd,Al)3

, Pd2Al3 < 952

hP5

P3m1

Ni2Al3

a = 422.7

c = 516.7

0 to 2 at.% Mn at 850°C, 0 to 1.6 at.%

Mn at 710°C [2002Yur]

38 to ~42.2 at.% Pd [2003Bal]

at 40 at.% Pd [2003Bal]

, (MnPdAl)

, PdAl (h)

1645 - 545

, MnPd

1515 - 540

cP8

Pm3m

CsCl

a = 308.3

a = 303.6

a = 273.3

~5 to 12 at.% Mn, ~57 at.% Al at

880-870°C in equilibrium with 3

[2002Yur]

MnPd2Al, at 1010-710°C [1968Web]

43.5 to ~56 at.% Pd [2003Bal]

34 to ~62 at.% Mn [Mas2]

HT [V-C2]

', PdAl (l)

< 850

hR78

R3

a = 1565.9

c = 525.1

~48.5 to ~52.2 at.% Pd [2003Bal]

< 740

CP8

P213

FeSi

a = 486.2 ~48 to ~49 at.% Pd [2003Bal]

, Pd5Al31315 - 615

oP16

Pbam

Rh5Ge3

a = 1047.1

b = 537.3

c = 503.5

62.5 at.% Pd [2003Bal]

, Pd2Al

< 1418

oP12

Pnma

Co2Si

a = 541.0

b = 405.5

c = 776.0

~65 to ~76 at.% Pd [2003Bal]

at 66.1 at.% Pd [2003Bal]

Pd5Al2< 980

oP28

Pnma

Pd5Ga2

a = 540.0

b = 403.4

c = 1840.5

~70.7 to ~71.7 at.% Pd [2003Bal]

Usually labeled as “ ”

Pd3Al

< 775

orthorhombic

P21ma

-

a = 540.7

b = 403.2

c = 1580.2

75 at.% Pd [2003Bal],


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

238


MSIT®

Al–Mn–Pd

Pd5Al

660

-

Pnma

-

a = 1070.0

b = 400.0

c = 807.4

83 at.% Pd [2003Bal],

sometimes called “ ”

1, MnPd

1200

tP4

P4/mmm

AuCu

a = 406.9

c = 358.5

31 to ~53 at.% Mn [Mas2]

low T [V-C2]

MnPd3

< 750

tI16

I4mm

Au3Cd

a = 391.3

b = 1549.6

32.5 to ~35.5 at.% Mn [Mas2]

[V-C2]

Mn3Pd5

< 500

oC16

Cmmm

Ga3Pt5

a = 807.2

b = 727.9

c = 404.4

~37.5 at.% Mn [Mas2]

[V-C2]

Mn11Pd21

< 197

tP32

P4/mmm

Mn11Pd21

a = 806.1

c = 733.0

~34.5 at.% Mn

[V-C2]

* 1, MnPd2Al

< 710

cF16

MnCu2Al

a = 618.2 RT, Mn25.2Pd49.7Al25.1 [1981Sol2]

* 2 (l)

< 864

decagonal

D3

decagonal

P105/mmc

a = 1200

a5 = 1255.7 0.1

a = 1240

a = 1250

quasicrystal, usually labeled as “D”

Mn17.5-17.9Pd13.5-12.1Al69-70 at

880-710°C [2002Yur]

Mn17Pd13Al70 annealed at 800°C 4d

[1993Hir1]

Mn16.5Pd13Al70.5 single crystal

V5D = 1030.1 10-10 pm5 [1993Ste]

Mn16.5Pd13Al70.5 (SEM-XMA) in

sample Mn2Pd2Al9 annealed at 780°C

2.5d and quenched in water [1991Bee]

very close to Mn18.1Pd12.1Al69.8,

samples annealed at 855 (5d), 830

(14d), 600°C (60d) [1995Bee1,

1995Bee2]

* 2 (h)

896 - 864

B-centered

orthorhombic

a = 2030

b = 1250

c = 6250

very close to Mn18.1Pd12.1Al69.8,

labeled as DH by [1995Bee1,

1995Bee2, 1995Goe2]

in samples annealed above 865°C

[1995Bee1, 1995Bee2, 1995Goe2]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

239


MSIT®

Al–Mn–Pd

* 3

< 893

< 715 [2000Let]

icosahedral

6D face-centered

hypercubic

lattice

Pm35

superstructure of

the F-phase

P-type

“diamond” type

cubic symmetry

approximately

m3,

superstructure of

the F2-phase

a6D =

645.99 0.03

a6D = 645.1

aF = 1290.1

aF = 1290.1

(a6D)F2 = 2 a6D

aP = 2091.1

aP = 2088.1

aP = aF

aF2 = 2581.0

aF2 = 2aF

aF = 1292.6

aP = 646.3 0.5

quasicrystal, usually labeled as “I”

Mn6-10Pd24.5-19.2Al69.5-70.8 [2002Yur]

Mn8-10.2Pd20.3-23.2Al68-69.5 [2000Kle1]

sometims labeled as “F”

Mn10Pd19Al71, spinning and

subsequent annealing, neutron

diffraction [1991Bou]


Bridgman method (EPMA, XRD).

Density 5.1 0.2 g cm-3.

Mn9Pd21Al70, obtained by plasma jet

melting and annealed at 800°C 3d

subsequently quenched in liquid

nitrogen. Neutron diffraction

[1992Bou];

Mn8.5Pd21Al70.5 single grains, grown

using conventional casting procedures,

X-ray and neutron diffraction

[1993Bou]

in Mn8Pd21Al71 sample annealed at

different temperatures (XRD)

[2000Hir]

in Mn7.5Pd21.5Al71 sample obtained by

plasma jet melting and annealed at

803 4°C 50h, quenched into water

and subsequently into liquid nitrogen

(XRD) [1995Ish]

sometimes labeled as F2,

not stable, corresponds to a transient

state in the process of the

transformation F to F2M [2000Let]

[2000Hir]

[2003Cap]

Mn7.5Pd21.5Al71 alloy obtained by

plasma jet melting and annealed at

602 2°C 48-400 h, quenched into

water and subsequently into liquid

nitrogen. XRD [1995Ish]

Mn8.8Pd21.4Al69.8 (EPMA) at RT

[2000Let]

sometims labeled as F2M

at RT, Mn8.7Pd22.0Al69.3 (EPMA,

XRD) [1998Boi]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

240


MSIT®

Al–Mn–Pd



Al Mn Pd

L + 1 H + 952 U1 L 66 12 22

1 H + + 2 923 E1 1 - - -

L+ H + 2 896 P1 L 70.5 19 10.5

L + 2 + 3 893 P2 L 71 19.5 9.5

L + 2 H + 3 887 U2 L 72 18.5 9.5

L + 3 + ~875 U3 L 73.5 21 5.5

3 + 2 + 860 U4 - - - -

L + + 3 845 P3 L 76 20.5 3.5

L + 3 H + 832 U5 L 73 15.5 6.5

H + 2 + 2 755 U6 - - - -

L + H + MnAl6 647 U7 L 96 2 2

L + MnAl6 (Al) + H 626 U8 L 95 4 1

L + H + (Al) 618 U9 L 92.5 7 0.5

241


MSIT®

Al–Mn–Pd

l + γ γ1

1048 p1

L + γ1

β + H952 U1

β+γ1+Η l + β δ

952 p3

l+γ1 Mn4Al

11(h)

1002 p2

γ1 + γ γ

2

991 e1

γ1

Mn4Al

11(h)+γ

957 e2

l+Mn4Al

11(h) µ

923 p4

l + δ ε28

792 p5

l + ε28

ε6

782 p6

l (Al) + ε6

616 e4

l + µ MnAl6

705 p7

l (Al)+MnAl6

658.5 e3

γ1

β + γ2 + H923 E

1

γ β

γ β

γ β

L + H + β τ2

896 P1

L + τ2 + β τ

3893 P

2

L + τ2

H + τ3

887 U2 L + β τ

3 + δ~875 U

3

τ3 + β τ

2 + δ860 U

4

L + δ + τ3

ε845 P3

L + τ3

H + ε832 U5

H + β τ2 + γ

2755 U

6

L + µ H + MnAl6

647 U7

L+MnAl6

(Al)+H626 U8

L + H (Al) + ε618 U9

L+ε+(Al)

L+H+β

L+β+τ2

γ+β2+H

τ2+H+β

L+H+τ2

L+τ2+τ

3

τ2+β+τ

3

L+β+τ3

β+τ3+δ

L+τ2+δ

β+τ2+δ

τ3+τ

2+δ

L+τ3+ε

τ2+H+τ

3

L+H+τ3

L+δ+εδ+τ3+ε

τ3+H+ε

δ+τ3+γ H+τ

2+γ

2

L+H+ε

L+H+MnAl6

L+(Al)+H

µ+H+MnAl6

MnAl6+(Al)+H

H+ε+(Al)

Fig. 1a: Al-Mn-Pd. Reaction scheme in the Al-rich part of the system

Al-Mn Al-Mn-Pd Al-Pd

242


MSIT®

Al–Mn–Pd

10

20

30

40

10 20 30 40

60

70

80

90

Mn 50.00Pd 0.00Al 50.00

Mn 0.00Pd 50.00Al 50.00


Axes: at.%

e4,616°C

U9

U8

U7

γ(MnPdAl)p,1177°C

U5 P

3

U3

τ3

P2

U2

P1

U1

µ

(αAl)

H

ε

β((MnPdAl)

130012001100

11001050

1025

1000

950

γ1

950

925900

875850

800750

1000 δ

p6,782°C

p5,792°C

p3,952°C

e3,658.5°C

p7,705°C

p4,923°C

p2,1002°C

p1,1048°C

MnxAl

1-x

MnAl6

a)

τ2

b)

δU

3P

1

U5

P3

p6p

5

τ2

l + γ γ1

1048 p1

L + γ1

H + β952 U1

l + β δ952 p

3

Fig. 1b: Al-Mn-Pd. Metastable reaction scheme in the Al-Mn-Pd system according to [1995Geo1]

Al-Mn Al-Mn-Pd Al-Pd

γ1 + γ γ

2

991 e

l+γ1

Mn4Al

11(h)

1002 p2

γ1

Mn4Al

11(h)+γ

957 e

γ1

β + γ2

+ H923 E1

H + τ3

+ β τ2

770 P2

L + H + β τ3

894,876 P1

γ β

γ β

γ β

β+γ1+H

L+H+β

τ3+H+β

β+γ2+H

U6

U3

L+H+τ3

U5

L+β+τ3

U3

H+β+τ2

U6

H+τ2+τ

3τ

2+β+τ

3

Fig. 2: Al-Mn-Pd.

Partial liquidus

surface projection:

a) - stable,

b) - metastable

243


MSIT®

Al–Mn–Pd

10

20

30

40

10 20 30 40

60

70

80

90

Mn 50.00Pd 0.00Al 50.00

Mn 0.00Pd 50.00Al 50.00


Axes: at.%

τ2

τ3

ε

δ

β(MnPdAl)

γ(MnPdAl)

MnAl6

µ

H

γ1

H+γ1

(αAl)

(αAl)+H+ε

H+τ3

β+τ3

β+δ

H+γ1+β

τ2+τ

3+β

618°C

705

626

647

923

1002

1048

952

792

616

832887

952

896

875

845

893

10

20

30

40

10 20 30 40

60

70

80

90

Mn 50.00Pd 0.00Al 50.00

Mn 0.00Pd 50.00Al 50.00


Axes: at.%

L

β(MnPdAl)

δ

H

γ2

τ2

Mn4Al

11(l)

µL+H

H+β

γ2+β

L+β

L+δ

γ2+β+H

τ2+β+L

γ(MnPdAl)

Fig. 3: Al-Mn-Pd.

Solidus surface

projection

Fig. 4: Al-Mn-Pd.

Isothermal section in

the Al-rich part of the

system at 894°C

244


MSIT®

Al–Mn–Pd

10

20

30

40

10 20 30 40

60

70

80

90

Mn 50.00Pd 0.00Al 50.00

Mn 0.00Pd 50.00Al 50.00


Axes: at.%

L

β(MnPdAl)

τ3

δ

L+δ

γ2

γ2+β

H

H+β

H+L

τ2 τ

2+τ

3+β

γ2+H+β

µ

Mn4Al

11(l)

τ3+β

γ(MnPdAl)

β+δ

10

20

30

40

10 20 30 40

60

70

80

90

Mn 50.00Pd 0.00Al 50.00

Mn 0.00Pd 50.00Al 50.00


Axes: at.%

H

Mn4Al

11(l)

µ

L

H+L

β(MnPdAl)

δ

τ3

ε

δ+β

L+H+τ3

γ2+β

H+τ3

H+β τ2+βγ

2

γ2+H+β

L+ε

τ2+τ

3+δτ

2

Fig. 5: Al-Mn-Pd.



system at 875°C

Fig. 6: Al-Mn-Pd.



system at 840°C

245


MSIT®

Al–Mn–Pd

10

20

30

40

10 20 30 40

60

70

80

90

Mn 50.00Pd 0.00Al 50.00

Mn 0.00Pd 50.00Al 50.00


Axes: at.%

L

ετ

3

δ

β(MnPdAl)

β+δ

γ2

γ2+β

τ2

γ2+τ

2+β

HMn

4Al

11(l)

µ

τ2+β

L+H+εL+ε

γ2+H

γ2+τ

2+H

H+τ3

µ+L+H

ε6

ε28

τ2+δ+τ

3

ε+δ

10

20

30

40

50

10 20 30 40 50

50

60

70

80

90

Mn 60.00Pd 0.00Al 40.00

Mn 0.00Pd 60.00Al 40.00


Axes: at.%

γ2

τ2

MnAl6

µ

Mn4Al

11(l)

(αAl)

(αAl)+H+ε

PdAl4

ε

τ3

Pd8Al

21

δ

µ

β´

β(MnPdAl)

τ2+γ

2+β

γ2+β

τ2+β

H

(βMn)(βMn)+β

γ2+H+τ

2

Fig. 7: Al-Mg-Pd.



system at 710°C

Fig. 8: Al-Mn-Pd.



system at 600°C

246


MSIT®

Al–Mn–Pd

10 20 30

600

700

800

900

1000

1100

Mn 0.00Pd 10.00Al 90.00

Mn 32.00Pd 10.00Al 58.00Mn, at.%

Te

mp

era

ture

, °C

L

L+H

L+H+ε

(Al)+H+ε

β+γ2

L+β

τ2+β+γ2

H+τ2+γ2

H+β+γ2

H+τ2+β

H+τ3

β

H+εH+τ3+ε

L+γ1+βL+γ1

L+γ1+H

L+H+β

γ1+β

β+γ1+γ2

τ2+γ2

H+β864

H+τ2

896887

L+H+τ3

H+β+γ1

618

(Al)+ε(Al)+ε+L

ε+L

~755

832

952

923

(Al)+PdAl4+ε

10 20

600

700

800

900

1000

Mn 0.00Pd 20.00Al 80.00

Mn 28.00Pd 20.00Al 52.00Mn, at.%

Te

mp

era

ture

, °C

L

L+β

L+ε

L+H+ε

H+τ3+εH+ε

~755

832

864

860

τ2+τ3+δ

τ2+β

τ2+H+β

618(Al)+ε

(Al)+ε+H

(Al)+PdAl4+ε

(Al)+ε+L

896

893

τ3+β

L+τ3

L+H+τ3

L+β+τ3

L+ε+τ3

τ3

H+τ3

τ2+τ3+H

τ2+τ3

τ2+β+γ2

H+γ2+β

β+γ2

β923

952 γ1+β

L+H+β

H+β

τ2+δ

τ2+τ3+β

τ2+δ+β

Fig. 9: Al-Mn-Pd.

Partial vertical section

at 10 at.% Pd

Fig. 10: Al-Mn-Pd.


at 20 at.% Pd

247


MSIT®

Al–Mn–Pd

10 20 30

600

700

800

900

1000

Mn 6.00Pd 0.00Al 94.00

Mn 6.00Pd 40.00Al 54.00Pd, at.%

Te

mp

era

ture

, °C

L+β

β

τ2+β+δ

L+β+δ

τ2+τ3+δ

τ3+δ

L+τ3+δ

L+H+ε

L+H

L+(Al)+H

L+µ

L+MnAl6

647

626618

832

845 860

~875τ3+β+δ

L+τ3+εL+τ3

L+τ3+β

β+δ

τ3+δ+ε

τ3+ε

H+εH+τ3+ε

L

(Al)+MnAl6

L+H+µ

L+H+MnAl6

705°C

τ2+δ658.8°C

L+µ+MnAl6

(Al)+H

H+(Al)+MnAl6

(Al)+H+ε

10 20

600

700

800

900

1000

Mn 30.00Pd 0.00Al 70.00

Mn 0.00Pd 30.00Al 70.00Pd, at.%

Te

mp

era

ture

, °C

952

923

L+γ1

H+β

γ1+HL+γ1+H

L+βL+γ1+β

L

γ2+H

H+β+γ2

H+τ2+β

γ2+Mn4Al11(l)

γ2+Mn4Al11(l)+H

H+τ2+γ2

H+τ2+τ3

τ2+H

H+τ3

864 887 L+τ2+τ3

L+H L+τ2 L+τ2+β893

896

~875

845

L+τ3+δ

L+βL+τ3+β

τ3

L+τ3

τ3+β

τ3+δ+ε

δ+ε

L+δ+ε

τ3+δ

952°C

792°C

τ2+τ3

τ2+τ3+δPd8Al21+δ

Pd8Al21+δ+ε

~755

895°C

975°C

1002°C

1048°C

640°C

L+δ

Fig. 11: Al-Mn-Pd.


at 70 at.% Al

Fig. 12: Al-Mn-Pd.


at 6 at.% Mn

248


MSIT®

Al–Mn–Pd

10 20

900

Mn 32.80Pd 0.00Al 67.20

Mn 0.00Pd 27.00Al 73.00Pd, at.%

Te

mp

era

ture

, °C

880

H+β

H+β+τ2

P1,896

P2,893

U2,887

L+H+β L

L+δ

L+τ3+δ

τ3+L

L+β+τ3

τ3+β+L

τ3+β

τ2+τ3+β

τ2

864

L+βL+β+δ

H+τ2+L

τ2+L τ2+τ3+L

γ2

H+γ2

H+β+γ2

τ2(h?)

U3,~875

920

860

L+H

10 20

800

900

1000

Mn 31.20Pd 0.00Al 68.80

Mn 0.00Pd 29.00Al 71.00Pd, at.%

Te

mp

era

ture

, °C

L

L+δ

L+τ3+δ

L+β

τ2+τ3H+τ2+β

τ2

τ2(h?)

L+H+β

L+H

L+τ2+τ3

L+τ3+βL+τ3

τ3

δ+ε

L+ε+δ

τ3+δ

P3,845

U3,~875

L+γ1

L+γ1+H

γ1

γ1+H

H+γ2

γ1+γ2+H

β+γ1+γ2L+τ2+β

H+β

H+β+γ2

Mn4Al11(l)+H+γ2

Mn4Al11(l)+γ2

τ3+δ+ε

U1,952

E1,923

P1,896

P2,893

L+δ+β895°C

957°C

U2,887

Fig. 13: Al-Mn-Pd.


from Mn32.8Al67.2 to

Pd27Al73

Fig. 14: Al-Mn-Pd.


from Mn31.2Al68.8 to

Pd29Al71

249


MSIT®

Al–Mn–Pd

10 20

600

700

800

900

1000

Mn 0.00Pd 3.00Al 97.00

Mn 13.00Pd 30.00Al 57.00Pd, at.%

Te

mp

era

ture

, °C

L

L+β

τ2+τ3+δ

τ2+β

L+H+ε

(Al)+ε+H

L+H

βL+τ3

U5,832

H+τ3+ετ3

H+ε

H+τ3

τ2+τ3

H+τ2+τ3

τ3+δ

L+τ3+β τ3+βτ2+τ3+β

U9,618

L+(Al)

L+(Al)+εL+(Al)+H

(Al)+ε

τ2+δ+β

U4,860

τ2+δ

Fig. 15: Al-Mn-Pd.


from Pd3Al98 to

Mn13Pd30Al57

Temperature , K22

Cp/T

,m

J/m

olK�

2

0 50 100 150

2

4

6

8Fig. 16: Al-Mn-Pd.

Low-temperature heat

capacities plotted in

the form of Cp/T vs T2

for the Mn10Pd20Al70

quasicrystal

[1997Ina]

250


MSIT®

Al–Mn–Pd

Temperature, K

Cp

B/K

0

0

200 400 600 800 1000

1

2

3

4

5

6

[1997Ina]

a)

b)

Temperature, K

Cp,

Jg

-ato

mK

-1-1

··

1

10-4

10 100

10-3

10-2

10-1

10

101

Temperature, K

0 300200100

0.5

1.0

1.5

2.0

(-

)/,

%C

CC

vp

v

Fig. 17: Al-Mn-Pd.

Temperature

dependence of

specific heat per atom

at constant pressure

(a) and at constant

volume (b) measured

for Mn9Pd20Al71

quasicrystal

[2000Eda]

Fig. 18: Al-Mn-Pd.

Low temperature

specific heat Cp(T) of

icosahedral

Al68.2Mn9Pd22.8 as a

function of

temperature between

1.6 and 280 K. The

Cp(T) and Cv(T) due

to thermal-expansion

effects, as a function

of temperature

[1998Wae]

251


MSIT®

Al–Mn–Pd

Temperature, K

Cp,

Jm

ol

K-1

-1·

·

1

0.001

10 100

0.01

0.1

1

10

Temperature, K

Ha

rdn

ess,

H/k

gm

mV

2·

0

300 400 500 600 700 800

200

400

600

900

2-fold3-fold5-fold

800

Fig. 20: Al-Mn-Pd.

Vickers hardness as a

function of

temperature for

several symmetrical

atomic surfaces in

as-grown single

icosahedral

Mn10Pd20Al70 ingots

with 3-fold growth

direction by the

Czochralski method

[1999Yok]

Fig. 19: Al-Mn-Pd.

Temperature

dependence of the

specific heat Cp of

Mn4Pd22Al74,

measured with a heat

relaxation system

[2002Sch]

252


MSIT®

Al–Mn–Pd

Temperature, K

aa/

(10

)0

3

0

0

100 200 400 500 600 700300

1

2

4

5

6

7

8

9

3

Fig. 21: Al-Mn-Pd.

Temperature

dependence of the

quasilattice constant

of the icosahedral

phase [2000Kaj]:

a = a(T)-a0, where

a0 = a(0 K)

253


MSIT®

Al–Mn–Ti

Aluminium – Manganese – Titanium

Andy Watson

Literature Data

A number of phase diagram studies have been made of the ternary system. [1954Dom, 1955Dom]

investigated some 100 ternary alloys ranging from pure titanium to the 60 at.% Al and 40 at.% Mn binary

limits. Metallographic and incipient melting techniques were employed. Samples prepared from iodide

titanium, high purity aluminium and electrolytically-refined manganese were homogenized by repeated arc

melting and annealing at 1000°C for 24 h. Phase equilibria were determined by the metallographic

examination of specimens quenched after annealing under argon or in vacuum at temperatures between 700

and 1200°C. Corresponding annealing times ranged from 17 to 4 d; no X-ray examination or tests for

equilibrium were reported. Six vertical sections and isothermal sections at 750, 800, 900, 1000, 1100 and

1200°C were constructed. The major variation in form occurs between 900 and 800°C, with a ( Ti)+TiAl

high temperature equilibrium being superseded at lower temperatures by an ( Ti)+TiMn2 phase field as a

result of a four-phase invariant reaction, ( Ti)+TiAl ( Ti)+TiMn2, which is reported to occur at 865°C. A

second four-phase reaction ( Ti)+TiMn2 ( Ti)+ TiMn was inferred to occur between 700 and 550°C. The

results of [1954Dom, 1955Dom] are summarized in [1974Zwi]. The isothermal sections reported by

[1954Dom, 1955Dom] are not completely consistent with more recent binary data in that they fail to take

into account the existence of the phases Ti3Al and TiMn and they assume that the Laves phase TiMn2 is

a purely binary phase of invariant composition. Electrical resistivity measurements and metallographic

observations on samples annealed at 600, 800, 1000 and 1200°C were used by [1960Sat] to construct a

partial phase diagram for alloys containing up to 49 at.% Al and 52 at.% Mn. Since this investigation

assumes the intervention of a high-temperature phase Ti3Al2 and associated reactions, L+( Ti) Ti3Al2(1620°C), L+Ti3Al2 TiAl (1460°C) and ( Ti)+Ti3Al2 Ti3Al (1400 to 1300°C), their proposed ternary

equilibria are generally incompatible with other studies. As in [1954Dom, 1955Dom], an invariant reaction

was proposed at 860 10°C, but [1960Sat] suggests that it takes the form ( Ti)+Ti3Al ( Ti)+ TiMn. A

limited study of ternary equilibria around the TiAl phase has been made by [1988Has] using optical

microscopy, electron probe microanalysis and X-ray diffraction. Extension of the Al-rich boundary of the

TiAl phase field to approximately 4.5Mn-46.7Al (at.%) at 1000°C was reported, but it should be noted that

in a second publication of these results [1989Tsu] the authors emphasize that equilibrium as not achieved

by the heat treatment used (7 d).

The manganese-rich regions of the system (0 to 40 at.% Al, 0 to 40 at.% Ti) have been investigated by

[1977Cha]. Samples were argon-arc melted from electrolytic or high purity grade materials. After annealing

under argon at 1000°C for 7 d the alloys were quenched and the phases present were determined using

optical metallography and X-ray powder diffraction techniques. Observations on the binary Al-Mn alloys

conflict with generally accepted information since they indicate the ( Mn) field as extending to only 39 at.%

Al. The Laves phase was shown to exhibit an extensive ternary homogeneity range projecting along lines

of approximately constant Ti content. Its limits were not completely established by [1977Cha] but X-ray

studies by [1974Dwi] and [1978Jac] show that it extends at least to the equiatomic ternary composition. The

presence of the hexagonal phase at the composition 37.5Ti-25Mn (at.%) has been reported by [1988Has].

Aluminium-rich alloys containing up to 1.2 at.% Ti and 2.3 at.% Mn have been investigated by [1958Mal].

Using differential thermal analysis, liquidus surfaces and the corresponding lines of the secondary reactions

L+TiAl3 (Al), L (Al)+MnAl6 and L TiAl3+MnAl6 were determined and an invariant four-phase reaction

point, L+TiAl3 (Al)+MnAl6 was located at 0.07Ti-0.94Mn (at.%) and 663°C. Solid state isothermal

equilibria were investigated by means of metallographic and X-ray examination of alloys annealed at 650°C

(121 h), 600°C (136 h), 500°C (228 h) and 400°C (168 h).

The ternary sections given by [1954Dom] and [1955Dom] did not include the Al rich portion of the diagram

and thus the ternary L12 phase was not seen. A number of studies have focussed on equilibria involving this

phase. The extent of the L12 phase field was investigated by [1991Nic], who arc melted Al, Mn and Ti

254


MSIT®

Al–Mn–Ti

(>99.9 mass%) under argon followed by a HIP treatment at 1200°C and 172 MPa for 2 h. Microstructures

were examined optically and by SEM. EDS and WDS analyses were also carried out. Later, [1993Nak]

studied the Al rich part of the diagram using arc melted samples (material purities: Al >99.999 mass%, Ti

>99.9 mass%, Mn >99.99 mass%) which had been homogenized for 2 d at 1000°C, and powdered metals

(purities: Al 99.9 mass%, Mn 99.9 mass%, Ti 99.5 mass%) which were sintered at 1000°C. They found that

the ternary phase field centered at Ti25Mn9Al66. [1996Mab, 1998Mab] used similar techniques and material

purities in their investigations at 1150°C. Also, they discovered that the ternary phase melts incongruently.

Equilibria involving the ( Ti)(hcp), ( Ti)(bcc) and the TiAl (L10) phases were studied experimentally by

[2000Kai] over the temperature range 1000-1300°C. Arc melted samples using Al, Mn and Ti of purity

99.99, 99.9 and 99.7 mass% respectively, were wrapped in Mo foil and sealed in Ar filled silica capsules

for annealing. Heat treatment times of 504 h, 168 h and 24 h were used for temperatures of 1000, 1200 and

1300°C, respectively. Microstructural and EPMA analysis revealed equilibria that were qualitatively in

agreement with [1954Dom] and [1955Dom], but the composition of the ( Ti) phase was substantially

different. [1996Che] studied arc-melted and annealed samples by SEM, EDS and EPMA. An alloy

composition of Ti-42Al-10Mn was chosen, and annealing temperatures of 1000 and 800°C were used. At

the higher temperature, an equilibrium between the ( Ti), TiAl and the TiMn2 Laves phase containing a

considerable amount of Al, given as Ti(AlMn)2 was found. The appearance of this Laves phase contradicts

the work of [1988Has] who suggested the presence of a ternary compound with the formula Ti3Mn2Al3.

The data for the equilibrium between the ( Ti) and the TiAl phase are consistent with the data of [2000Kai].

At the lower temperature, a three phase equilibrium between ( Ti), TiAl and the Laves phase was found

suggesting a four phase invariant just above 800°C, as was presented by [1955Dom].

While [1955Dom] and [1960Sat] both indicated the presence of a four-phase solid transition reaction at

approximately 860 10°C, they disagree on the phases involved. Considering the extent of observed

stability ranges of ( Ti) [1955Dom, 2000Kai] and Ti(Mn,Al)2 [1974Dwi, 1977Cha, 1978Jac, 1996Che], the

reaction suggested by [1960Sat] involving ( Ti) and TiMn perhaps appears least probable. Taking into

account currently accepted binary Al-Ti equilibria suggests, however, that TiAl3 rather than ( Ti) is likely

to be a product in the reaction proposed by [1955Dom]. However, [1996Che] suggest that the nature of the

reaction is eutectoid, (Ti) (TiAl)+Ti(Mn,Al)2, but unfortunately, this refers to unpublished work.

In a series of articles, [1997But, 1998But, 1999But] investigated the solidification behavior and phase

transition sequences in alloys with a Ti:Al ratio of 1:14, plus 5, 10, 20 and 30 at.% Mn (for [1999But] the

alloy used was Ti-20Mn-37Al (at.%). DTA measurements and microstructural observation were used and

the results were compared with thermodynamic calculations. Only partial agreement was obtained, but the

thermodynamic data used for the calculations were only extrapolations from the edge binary systems. Some

agreement between the experimental results and those of [2000Kai] were found at 1300°C for the 5 at.%

Mn composition.

CVM calculations of the locations of the (Ti)/Ti3Al and Ti3Al/TiAl phase boundaries were conducted by

[2001Kan]. The results agreed well with the experimental data of [1988Has].

The crystallography of the ternary L12 phase has been discussed by a number of authors [1990Kum,

1991Dur, 1991Mae, 1992Mor, 1996Mab, 2001Mil]. Most authors suggest that the ternary compound is

based on the addition of Mn to the binary TiAl3 compound. However, [1991Dur] argues that following

geometrical considerations the compound should be based on Ti5Al11. The site selectivity of Mn in TiAl

has been studied by [1991Bab]. They conclude that at x < 1.85 in Ti50-xMnxAl50, Mn substitutes for Ti. At

higher Mn concentrations, some Al sites are also occupied by Mn. [1993Ers] conducted LMTO-ASA

calculations to study the effects of Mn substitution on Al and Ti sites in TiAl. The calculations were made

for compositions of TiMnAl2 and Ti2MnAl but they argue that it is possible to extrapolate their results to

small percentage additions of Mn to the binary compound. They found that Mn substitutes for Al which is

in broad agreement with the earlier work of [1991Bab]. This work was confirmed by [1999Hao], who used

atom location channelling enhanced microanalysis. They determined the substitutional sites by Mn

(1-5 at.%) in TiAl and (1-2 at.%) in Ti3Al. They found Mn substitution for Ti in Ti3Al.

255


MSIT®

Al–Mn–Ti

Binary Systems

The Al-Ti and Al-Mn binary systems produced by the MSIT binary evaluation programme have been

accepted [2003Pis, 2003Sch]. The Mn-Ti phase diagram is taken from [Mas2].

Solid Phases

The stable phases are given in Table 1. There is only one ternary phase which occurs at a composition of

Ti25Mn9Al66. It was found that the range of homogeneity increased with increasing temperature

[1996Mab]. Allusions had been made to two more ternary compounds (TiMnAl [1974Dwi] and Ti3Mn2Al3[1988Has]) but these are more likely to be the Ti(Mn,Al)2 Laves phase owing to the extensive solubility of

Al in the binary TiMn2 compound [1996Che].


The only ternary invariant reaction established with certainty is the U type liquid transition reaction reported

in aluminium-rich alloys at 663°C. A partial reaction scheme is given in Fig. 1. The temperature and

composition of the liquid and ( Al) solid phases given in Table 2 were read from diagrams given in

[1958Mal]; negligible ternary solution in TiAl3 and MnAl6 was indicated. The presence of a four phase

reaction has been suggested at approximately 860 10°C by [1955Dom, 1960Sat, 1996Che]. Considering

the stability ranges of the phases involved, it is most likely to be the transition reaction,

( Ti)+TiAl Ti3Al+Ti(Mn,Al)2. However, owing to uncertainties in the locations of the phase boundaries

of the phases involved, it is not possible to give their compositions.


Investigations of ternary liquidus surfaces are confined to those in the extremely aluminium-rich corner of

the system. Figure 2 shows the isotherms of the surfaces of primary crystallization of ( Al), TiAl3 and

MnAl6, the secondary reaction lines, the tertiary point and the solid limits of the tertiary reaction plane as

reported by [1958Mal]. It was necessary to make some adjustments to the location of the isotherms to ensure

agreement with the accepted Al-Ti binary phase diagram.

Isothermal Sections

Figures 3 and 4 show partial isothermal sections for equilibria involving the ( Ti), ( Ti) and TiAl phases

at 1300 and 1200°C taken from [2000Kai]. Very minor adjustments have been made to ensure consistency

with the phase boundaries of the accepted Al-Ti phase diagram. Figure 5 presents the equilibria surrounding

the L12 ternary phase based on the work of [1998Mab]. However, the information in the original article

referring to equilibria between the ternary phase and the Mn8Al5 phase have been ignored here. This binary

phase is unstable above 991°C, and moreover, the liquid phase is stable at Al contents up to around 60 at.%

in the binary system at this temperature. No equilibria involving the liquid phase are given in the original

article. Phase boundaries drawn in Fig. 5 involving the Al-Mn binary should therefore be taken as very

tentative. Figure 6 shows the isothermal section for 1000°C. It was constructed by combining the results of

[1958Mal, 1974Dwi, 1993Nak, 1996Che] and [2000Kai]. Modifications to the original phase boundaries

were made to maintain consistency with the accepted binary phase diagrams. This was particularly

important with respect to equilibria between the ternary phase and the Al-Mn binary edge as in Fig. 5. The

work of [1993Nak] claimed equilibrium between the ternary phase and the Mn5Al8 ( 2) phase at 1000°C.

This is unlikely as this phase is unstable at this temperature in the binary system as indicated above.

Therefore, tentative equilibria between the ternary compound and the and 1 phases of the binary Al-Mn

system have been added to maintain consistency with the accepted binary diagram. The composition range

of the 1 phase increases with increasing temperature which also results in a slight shift of the phase field

in Fig. 6 with respect to Fig. 5. The work of [1996Che] and [2000Kai] suggest a higher solubility of Al and

Mn in the ( Ti) phase than the earlier work. Figure 7 shows the partial isothermal section for 800°C given

by [1955Dom] with minor alterations to allow consistency with the binary edges. Figure 8 is a composite

256


MSIT®

Al–Mn–Ti

diagram of partial isothermal sections for Al-rich alloys for temperatures 650-400°C taken from [1958Mal]

with some minor adjustments to the phase boundaries coincident with the accepted binary phase diagrams.


Both TiAl and TiAl3 have been identified as possible materials for aerospace applications owing to their

high temperature stability and their low density. However, they suffer from poor room temperature ductility

and poor workability, even at high temperatures. Hence there has been much interest in studying the

mechanical properties of materials consisting of TiAl or TiAl3 alloyed with a third component [1991Kum,

1991Mae, 1991Nic, 1992Win, 1992Mor, 1993Has, 1996Mab, 1999Has, 2000Jin, 2001Mil]. It was found

that the addition of Mn to either of the binary compounds improved ductility at room temperature. The

increased ductility correlates with a lowering of the antiphase boundary energy allowing formation of

partial super dislocations [1992Mor].

Studies of the hydrogen absorption/desorption of Ti3Al found that absorption properties were improved

with the addition of Mn to the compound [2001Ish]. Mn substituted Ti3Al showed a reduction in the

hydrogen desorption temperature.

Nanocrystallites of the L12 phase have been prepared by ball-milling [1998Var]. A Ti25.6Mn9.4Al65 alloy,

homogenized for 100 h at 1000°C was ball milled for 386 h. It was found that the crystallite size of

~20-30 m produced was unchanged after 200 h of milling.

References

[1954Dom] Domagala, R.F., Rostoker, W., “The System Titanium - Aluminium - Manganese”, Trans.

Am. Soc. Met., Reprint No. 4 (1954) (Equi. Diagram, Experimental, #, *, 6)

[1955Dom] Domagala, R.F., Rostoker, W., “The System Titanium - Aluminium - Manganese”, Trans.

Am. Soc. Met., 47, 565-577 (1955) (Equi. Diagram, Experimental, #, *, 6)

[1958Mal] Mal`tsev, M.V., Van Bok, Y., “Investigation of the Equilibrium Diagram of the

Aluminium-Manganese-Titanium System” (in Russian), Izv. Vyss. Uchebn. Zaved., Tsvetn.

Metall., (2), 130-142 (1958) (Experimental, Equi. Diagram, #, *, 3)

[1960Sat] Sato, T., Huang, Y.-C., Kondo, Y, “A Metallographic Study on Titanium - Aluminium -

Manganese Alloys”, Sumitomo Light-Metal Technical Reports, 1, 36-44 (1960) (Equi.


[1974Dwi] Dwight, A.E., “Alloying Behaviour of Zirconium, Hafnium and the Actinides in Several

Series of Isostructural Compounds”, J. Less-Common Met., 34, 279-284 (1974)

(Experimental, Crys. Structure, 6)

[1974Zwi] Zwicker, U., “Titanium and Titanium Alloys” in “Pure and Applied Metallurgy in

Individual Descriptions” (in German), 21, 576-585 (1974) (Equi. Diagram, Review, 22)

[1977Cha] Chakrabarti, D.J., “Phase Stability in Ternary Systems of Transition Elements with

Aluminium”, Metall. Trans. B., 8B, 121-123 (1977) (Experimental, Equi. Diagram, #, *, 13)

[1978Jac] Jacob, I., Shaltiev, D., “A Note on the Influence of Aluminium on the Hydrogen Sorption

Properties of Ti(AlxB1-x)2 (B = Cr, Mn, Fe, Co)”, Mater. Res. Bull., 13, 1193-1198 (1978)


[1988Has] Hashimoto, K., Doi, H., Kasahara, K., Tsujimoto, T., Suzuki, T., “Effects of Third Elements

on the Structures of TiAl-Based Phases”, J. Jpn. Inst. Met., 52, 816-825 (1988)

(Experimental, Equi. Diagram, Crys. Structure, 31)

[1989Tsu] Tsujimoto, T., Hashimoto, K., “Structure and Properties of TiAl-Base Alloys Containing

Manganese”, Mater. Res. Soc. Symp. Proc., 133, 391-396 (1989) (Experimental, Crys.

Structure, 9)


Cr, Mn, Fe, Co, Ni, Cu, Zn)”, Int. Mater. Rev., 35(6), 293-327 (1990) (Crys. Structure, Equi.



Z. Metallkd., 81(6), 389-396 (1990) (Crys. Structure, Equi. Diagram, Experimental,

Review, #, 33)

257


MSIT®

Al–Mn–Ti

[1991Bab] Babu, S.V., Seehra, M.S., “Site Selectivity of Mn Atoms in -TiAl Alloys Determined by

X-Ray Scattering”, J. Mater. Res., 6(2), 339-342 (1991) (Crys. Structure, Experimental, 10)

[1991Dur] Durlu, N., Inal, O.U., Yost, F. G., “L1(2)-Type Ternary Titanium Aluminides of the

Composition Ti25X8Al67: TiAl3-Based or TiAl2-Based?”, Scr. Metall. Mater., 25(11),

2475-2479 (1991) (Crys. Structure, Review, 30)

[1991Kum] Kumar, K.S., Brown, S.A., Whittenberger, J.D., “Compression, Bend and Tension Studies

on Forged Al67Ti25Cr8 and Al66Ti25Mn9 L1(2) Compounds”, Mater. Res. Soc. Symp.

Proc.: High-Temp. Ordered Intermetallic Alloys IV, 213, 481-486 (1991) (Mechan. Prop.,

Experimental, 11)

[1991Mae] Maeda, T., Okada, M., Shida, Y., “Ductility and Strength in Mo Modified TiAl”, Mater.

Res. Soc. Symp. Proc.: High-Temp. Ordered Intermetallic Alloys IV, 213, 555-560 (1991)

(Experimental, Phys. Prop., 15)

[1991Nic] Nic, J.P., Zhang, S., Mikkola, D.E., “Alloying of Al3Ti with Mn and Cr to Form Cubic

L1(2) Phases”, Mater. Res. Soc. Symp. Proc.: High-Temp. Ordered Intermetallic Alloys IV,

213, 697-702 (1991) (Crys. Structure, Equi. Diagram, Experimental, 12)

[1992Kat] Kattner, U.R., Lin, J.-C., Chang, Y.A., “Thermodynamic Assessment and Calculation of the

Ti-Al System”, Metall. Trans. A, 23(8), 2081-2090 (1992) (Assessment, Calculation, Equi.

Diagram, Thermodyn., #, *, 51)

[1992Mor] Morris, D.G., Gunter, S., “Ordering Ternary Atom Location and Ageing in Ll2 Trialuminide

Alloys”, Acta Metall. Mat., 40(11), 3065-3073 (1992) (Crys. Structure, Mechan. Prop.,

Experimental, 23)

[1992Win] Winnicka, M.B., Varin, R.A., “Microstructure and Ordering of L12 Titanium

Trialuminides”, Metall. Trans. A, 23A(11), 2963-2972 (1992) (Crys. Structure, Mechan.


[1993Ers] Erschbaumer, H., Podloucky, R., Rogl, P., Temnitschka, G., Wagner, R., “Atomic

Modelling of Nb, V, Cr and Mn Substitutions in -TiAl. I: c/a Ratio and Site Preference”,

Intermetallics, 1, 99-106 (1993) (Calculation, Crys. Structure, 31)

[1993Has] Hashimoto, K., Masao, K., “Effects of Third Element Addition on Mechanical Properties of

TiAl”, Struct. Intermet.: 1st Int. Symp. Struct. Intermetallics, Champion Pa. Sept., 309-318

(1993) (Equi. Diagram, Experimental, Mechan. Prop., 18)

[1993Nak] Nakayama, Y., Mabuchi, H., “Formation of Ternary L1(2) Compounds in Al3Ti-Base

Alloys”, Intermetallics, 1, 41-48 (1993) (Crys. Structure, Equi. Diagram, Experimental, 40)

[1996Che] Chen, Z., Jones, I.P., Small, C.J., “Laves Phase in Ti-42Al-10Mn Alloy”, Scr. Mater., 35(1),

23-27 (1996) (Equi. Diagram, Experimental, *, 14)

[1996Mab] Mabuchi, H., Kito, A., Nakamoto, M., Tsuda, H., Nakayama, Y., “Effects of Manganese on

the L12 Compound Formation n Al3Ti-Based Alloys”, Intermetallics, 4, S193-S199 (1996)

(Experimental, Equi. Diagram, 34)

[1997But] Butler, C.J., McCartney, D.G., Small, C.J., Horrocks, F.J., Saunders, N., “Solidification

Microstructures and Calculated Phase Equilibria in the Ti-Al-Mn System”, Acta Mater.,

45(7), 2931-2947 (1997) (Calculation, Equi. Diagram, Experimental, 26)

[1997Sah] Sahu, P.Ch., Chandra Shekar, N.V., Yousuf, M., Govinda Rajan, K., “Implications of a

Pressure Induced Phase Transition in the Search for Cubic Ti3Al”, Phys. Rev. Lett., 78(6),


[1998But] Butler, C.J., McCartney, D.G., “An Experimental Study of Phase Transformations and a

Comparison with Calculated Phase Equilibria in Ti-Al-Mn Alloys”, Acta Mater., 46(6),

1875-1886 (1998) (Calculation, Crys. Structure, Equi. Diagram, Experimental, 31)

[1998Mab] Mabuchi, H., Tsuda, H., Tateno, T., Morii, K., “Phase Equilibrium and the Formation of a

Graded Diffusion Layer by Bonding L10- and L12- Alloys in the Ti-Al-Mn System” (in

Japanese), J. Jpn. Inst. Met., 62(11), 999-1005 (1998) (Equi. Diagram, Experimental, 13)

[1998Var] Varin, R.A., Wexler, D., Calka, A., Tbroniek, L., “Formation of Nanocrystalline Cubic

(L1(2)) Titanium Trialuminide by Controlled Ball Milling”, Intermetallics, 6, 547-557

(1998) (Calculation, Crys. Structure, Experimental, Mechan. Prop., 26)

[1999But] Butler, C.J., McCartney, D.G., “Phase Transformations and Phase Equilibria in a Ti-37%

Al-20% Mn Alloy”, Intermetallics, 7, 663-669 (1999) (Calculation, Equi. Diagram,

Experimental, 14)

258


MSIT®

Al–Mn–Ti

[1999Hao] Hao, Y.L., Xu, D.S., Cui, Y.Y., Yang, R., Li, D., “The Site Occupancies of Alloying

Elements in TiAl and Ti3Al Alloys”, Acta Mater., 47(4), 1129-1139 (1999) (Crys.


[1999Has] Hashimoto, K., Yamamoto, Y., Kimura, T., Nobuki, M., “Effect of Vanadium on Residual

Strain in L12-Type (AlMn)3Ti(V) Alloy Powders and Bend Ductility of Pre-Milling

Alloys”, Mater. Trans., JIM, 40, 400-403 (1999) (Crys. Structure, Experimental, Phys.

Prop., 14)

[2000Dub] Dubrovinskaia, N., Dubrovinsky, L., Vennstrom, M., Anderson, Y., Abrikosov, I.,

Eriksson, O., “High-Pressure, High-Temperature In-Situ Study of Alloys: Ti3Al”, Proc.

Disc. Meet. Thermodyn. Alloys, 23 (2000) (Thermodyn.)

[2000Jin] Jinxu, Z., Gengxiang, H., Jiansheng, W., “Electron Structure and Bonding Characteristics

of Al3Ti Intermetallic Alloys”, J. Mater. Sci. Lett., 19(18), 1685-1686 (2000) (Equi.

Diagram, Experimental, Phys. Prop., 7)

[2000Kai] Kainuma, R., Fujita, Y., Mitsui, H., Ishida, K., “Phase Equilibria Among Alfa (hcp), (bcc)

and (L1(0)) Phases in Ti-Al Base Ternary Alloys”, Intermetallics, 8, 855-867 (2000)

(Crys. Structure, Equi. Diagram, Experimental, #, *, 29)


Binary System Ti-Al”, Metall. Mater. Trans. A, 32A, 1037-1048 (2001) (Crys. Structure,

Equi. Diagram, Experimental, #, *, 34)

[2001Ish] Ishikawa, K., Hashi, K., Suzuki, K., Aoki, K., “Effect of Substitutional Elements on the

Hydrogen Absorption-Desorption Properties of Ti3Al Compounds”, J. Alloys Compd., 314,

257-261 (2001) (Phys. Prop., Experimental, 9)

[2001Kan] Kang, S.Y., Onodera, H., “Analyses of HCP/D019 and D019/L10 Phase Boundaries in

Ti-Al-X (X = V, Mn, Nb, Cr, Mo, Ni and Co) Systems by the Cluster Variation Method”,

J. Phase Equilib., 22, 424-430 (2001) (Calculation, 15)

[2001Mil] Milman, Yu.V., Miracle, D.B., Chugunova, S.I., Voskoboinik, I.V., Korzhova, N.P.,

Legkaya, T.N., Podrezov, Yu.N., “Mechanical Behaviour of Al3Ti Intermetallic and

L1sub/2/ Phases on Its Basis”, Intermetallics, 9, 839-845 (2001) (Crys. Structure,


[2003Pis] Pisch, A., “Al-Mn (Aluminium-Manganese)”, MSIT Binary Evaluation Program, in MSIT







Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

( Al) hP2

P63/mmc

Mg

a = 269.3

c = 439.8

at 25°C, 20.5 GPa [Mas2]

( Al)

< 660.452

cF4

Fm3m

Cu

a = 404.96 at 25°C [Mas2]

( Mn)

1246 - 1138

cI2

Im3m

W

a = 308.0 [Mas2]

( Mn)

1138 - 1100

cF4

Fm3m

Cu

a = 386.0 [Mas2]

259


MSIT®

Al–Mn–Ti

( Mn)

1100 - 727

cP20

P4132

Mn

a = 631.52 [Mas2]

( Mn)

< 727

cI58

I43m

Mn

a = 891.26 at 25°C [Mas2]

( Ti) hP3

P6/mmm

Ti

a = 462.5

c = 281.3

at 25°C, HP 1 atm [Mas2]

( Ti)

1670 - 882

cI2

Im3m

W

a = 330.65 [Mas2]

( Ti)

< 882

hP2

P63/mmc

Mg

a = 295.06

c = 468.35

at 25°C [Mas2]

MnAl6< 705

oC28

Cmcm

MnAl6

a = 755.51

b = 649.94

c = 887.24

[V-C2]

Mn4Al11 (HT)

916 - 1002

oP160

Pnma

? [2003Pis]

< 1177

cI2

Im3m

W

a = 306.3 [V-C2] Also designated Mn55Al45

1

< 1048

2, Mn5Al8< 991

hR26

R3m

Cr5Al8

a = 1273.9

c = 1586.1

at 58 at.% Al [V-C2]

, Mn3Al2< 1312

hP2

P63/mmc

Mg

a = 270.5 - 270.5

c = 436.1 - 438

44.2 - 44.9 at.% Al [2003Pis]

TiMn

950

tP30

P42/mnm

CrFe

a = 888

c = 454.2

[Mas2], [V-C2]

TiMn

1200

- - [Mas2]

Ti(Mn,Al)2

TiMn2

< 1325

hP12

P63/mmc

MgZn2

a = 495.3

c = 805.8

a = 497.8

c = 815.1

a = 499.7

c = 890.8

a = 483.33 0.09

c = 793.84 0.11

at 17 at.% Al and 50 at.% Mn [1978Jac]

at 33.3 at.% Al and 33.3at.% Mn

[1974Dwi]

at 37.5 at.% Al and 25 at.% Mn

[1988Has]

[Mas2], [V-C2]. 60-70 at.% Mn

TiMn3

~1250 - 950

o**? - [Mas2]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

260


MSIT®

Al–Mn–Ti

TiMn4 hR53

R3m

Co5Cr2Mo3

a = 1100.3

c = 1944.6

[V-C2], [Mas2]. Phase referred to as

TiMn5 in [V-C2]. Prototype given as

~ (Mo,Ni) in [Mas2]

Ti3Al

< 1164

(up to 10 GPa at RT)

hP8

P63/mmc

Ni3Sn

a = 580.6

c = 465.5

a = 574.6

c = 462.4

~20 to 38.2 at.% Al [1992Kat]

D019 ordered phase (“ 2-Ti3Al”)

[1997Sah].

at 22 at.% Al [L-B]

at 38 at.% Al [L-B]

Ti3Al (I)

15 to > 41 GPa

hP16

P63/mmc

TiNi3

a = 531.2

c = 960.4

[1997Sah] at 16 GPa,

not confirmed by [2000Dub] (0-35 GPa,

25-2250°C)

TiAl

< 1463

tP4

P4/mmm

AuCu

a = 400.0

c = 407.5

a = 398.4

c = 406.0

46.7 to 66.5 at.% Al [1992Kat]

50 to 62 at.% Al at 1200°C [2001Bra]

L10 ordered phase (“ -TiAl”)

at 50.0 at.% Al, [2001Bra]

at 62.0 at.% Al, [2001Bra]

TiAl2< 1199

oC12

Cmmm

ZrGa2

tP4

P4/mmm

AuCu

tI24

I41/amd

HfGa2

tP32

P4/mbm

Ti3Al5

a = 1208.84

b = 394.61

c = 402.95

a = 403.0

c = 395.5

a = 397.0

c = 2497.0

a = 1129.3

c = 403.8



metastable modification of TiAl2, only

observed in as-cast alloys [2001Bra];

listed as TiAl2(h) by [1990Sch] (66 to 67

at.% Al, 1433-1214°C)

Ti1-xAl1+x; 63 to 65 at.% Al at 1250°C,

stable range 1445-1170°C [2001Bra];

listed as orthorhombic, Pmmm, with

pseudotetragonal cell by [1990Sch]

(range ~1445-1424°C).

at 1300°C [2001Bra]

stable structure of TiAl2 <1216

[2001Bra];

listed as TiAl2(r) by [1990Sch]

Ti3Al5, stable below 810°C [2001Bra]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

261


MSIT®

Al–Mn–Ti


“Ti2Al5”

1416 - 990

tetragonal

superstructure of

AuCu-type

[2001Bra]

tP28

P4/mmm

“Ti2Al5”

a* = 395.3

c* = 410.4

a* = 391.8

c* = 415.4

a = 390.53

c = 2919.63



Ti5Al11 stable range 1416 - 995°C

[2001Bra]

66 to 71 at.% Al at 1300°C [2001Bra]

(including the stoichiometry Ti2Al5!);

[1990Sch] claimed: 68.5 to 70.9 at.% Al

and range 1416 - 1206°C;

at 66 at.% Al [2001Bra]

* AuCu subcell only

at 71 at.% Al [2001Bra]

* AuCu subcell only

“Ti2Al5”

~1215 - 985°C [1990Sch];

included in homogeneity region of

Ti5Al11 [2001Bra]

TiAl3 (h)

< 1393

a = 384.9

c = 860.9

74.2 to 75.0 at.% Al [1992Kat]

74.5 to 75 at.% Al at 1200°C [2001Bra]

D022 ordered phase

stable above 735°C (Al-rich) [2001Bra]

TiAl3 (l)

< 950 (Ti-rich)

tI32

I4/mmm

TiAl3 (l)

a = 387.7

c = 3382.8

74.5 to 75 at.% Al [2001Bra]

* 1, Ti25Mn9Al66 cP4

Pm3m

AuCu3

a = 395.8

a = 395.9

Ti25Mn8Al67 [1991Nic]

Ti43Mn11Al66 [2001Mil]


Al Mn Ti

( Ti) + TiAl Ti3Al +

Ti(Mn,Al)2

~865 U1 - - - -

L + TiAl3 ( Al) + MnAl6 663 U2 L

TiAl3(Al)

MnAl6

98.99

75.0

99.37

85.7

0.94

0

0.09

14.3

0.07

25.0

0.54

0

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

262


MSIT®

Al–Mn–Ti

Al-Ti

L+TiAl3

(Al)+MnAl6

663 U2

l+TiAl3

(Al)

664 p

l (Al)+MnAl6

658.5 e

(Al)+TiAl3+MnAl

6

L+TiAl3+MnAl

6

Al-MnAl-Mn-Ti

Fig. 1: Al-Mn-Ti. Partial reaction scheme

Ti 2.30Mn 0.00Al 97.70

Ti 0.00Mn 2.30Al 97.70


Axes: at.%

700°C

750°C

800°C

850°C900°C

950°C

p (664°C)

(αAl)

U2 (663°C)

e (658.5°C)

MnAl6TiAl

3

Fig. 2: Al-Mn-Ti.

Partial liquidus

surface

263


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Al–Mn–Ti

50

60

10 20

40

50

Ti 70.00Mn 0.00Al 30.00

Ti 40.00Mn 30.00Al 30.00

Ti 40.00Mn 0.00Al 60.00 Data / Grid: at.%

Axes: at.%

(βTi)

(αTi)

TiAl

50

60

70

10 20 30

30

40

50

Ti 75.00Mn 0.00Al 25.00

Ti 40.00Mn 35.00Al 25.00


Axes: at.%

(βTi)

(αTi)

TiAl

Fig. 3: Al-Mn-Ti.

Partial isothermal

section at 1300°C

Fig. 4: Al-Mn-Ti.

Partial isothermal

section at 1200°C

264


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Al–Mn–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Mn


Axes: at.%L

TiAl3

TiAl2

TiAl

Ti3Al

(αTi)

βTiMn TiMn3 TiMn

4(αMn)

γ

ε

γ1

Mn4Al

11(h)

Ti3Al+(βTi)

(βTi) (βTi)+Ti(Mn,Al)2

Ti(Mn,Al)2

Ti(Mn,Al)2

+(βMn)

(βMn)

Ti2Al

5

τ1

20

40

60

20 40 60

40

60

80

Ti 70.00Mn 0.00Al 30.00

Ti 0.00Mn 70.00Al 30.00


Axes: at.%

TiAl

TiAl3

Ti2Al

5

TiAl2

τ1

Ti(Mn,Al)2

L

γ

ε

Fig. 6: Al-Mn-Ti.


1000°C

Fig. 5: Al-Mn-Ti.

Partial isothermal

section at 1150°C

(Al-rich part)

265


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Al–Mn–Ti

40

60

80

20 40 60

20

40

60

Ti Ti 30.00Mn 70.00Al 0.00


Axes: at.%TiAl2

TiAl

Ti3Al

(αTi)

(βTi)+αTiMn

Ti3Al+(βTi)+Ti(Mn,Al)

2

TiAl+Ti(Mn,Al)2+τ

1

(βTi)

αTiAl

Ti 1.00Mn 0.00Al 99.00

Ti 0.00Mn 1.00Al 99.00


Axes: at.%400°C

500°C

550°C

600°C

650°C

400°

C50

0°C

550°

C

600°

C

650°

C

(αAl)

(αAl)+TiAl3+MnAl

6

(αAl)+MnAl6

(αAl)+TiAl3

Fig. 7: Al-Mn-Ti.

Partial isothermal

section at 800°C

(Ti-rich part)

Fig. 8: Al-Mn-Ti.

Isothermals 650 to

400°C (Al-rich part)

266


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Al–Mo–Ni

Aluminium – Molybdenum – Nickel

Kostyantyn Korniyenko, Vasyl Kublii

Literature Data

Experimental investigations of the phase equilibria in the Al-Mo-Ni system were started by [1925Pfa] and

[1933Roe] and, as summarized in [1976Mon], concerned the Ni-rich corner as well as the influence of

additions of Mo and Ni on the Al solid solution, respectively. The investigations of the partial Ni-NiAl-Mo

system were further developed by [1959Gua2, 1960Bag, 1965Ram, 1976Jac, 1977Aig, 1977Pea, 1978Gul,

1983Nas, 1983Wak, 1984Kov1, 1984Kov2, 1984Mir, 1985Nas, 1986Mas1, 1986Mas2, 1988Mas,

1989Mas, 1989Hon1, 1989Hon2, 1991Mis]. Results of phase equilibria studies for the Al-rich corner are

presented by a series of isothermal sections [2002Gru]. The complete ternary system has been investigated

experimentally at 600°C [1971Pry], 800°C [1969Mar] and 950°C [1969Vir]. For preparation of the alloys

most of the authors used arc melting, while [1971Pry] and [2002Gru] applied levitation induction melting,

and [1984Mir] obtained specimens by both conventional arc-casting and powder metallurgy techniques.

The traditional methods of investigations were X-ray diffraction (XRD), metallography, differential

thermal analysis (DTA), electron microprobe analysis (EMPA). Some authors used scanning electron

microscopy (SEM) [1989Hon1, 1989Hon2, 1991Mis, 2002Gru], transmission electron microscopy (TEM)

[2002Gru], as well as energy-dispersive X-ray spectroscopy (EDXS) [1991Mis]. Calculations of phase

equilibria were carried out by [1974Kau, 1999Kau] and [1999Lu]. A critical review of literature data on

phase equilibria in the Al-Mo-Ni system was presented in the assessment of [1993Kub]. Further

experimental studies are necessary in order to construct the liquidus surface and the reaction scheme of the

complete ternary system as well as isothermal sections in the whole range of compositions [1969Mar,

1969Vir, 1971Pry].

Binary Systems

The Al-Mo, Al-Ni and Mo-Ni systems are accepted from [2003Sch], [2003Sal] and [Mas2], respectively.

Solid Phases

Crystallographic data on the known unary, binary and ternary phases are listed in Table 1. [1959Gua2]

reported the existence of a ternary phase of composition Mo7,5Ni58,0Al34,5 at 1175°C, but did not

determine its crystal structure. However, the data of [1960Bag, 1969Mar, 1971Pry, 1983Nas] and

[1984Mir] did not confirm its existence. A phase of similar composition was easily obtained by [1969Vir]

in the alloys (at.%) Mo50Ni25Al25, Mo43Ni31Al26 and Mo9Ni53Al38, the latter being fairly close to the

composition of the reported phase [1959Gua2]. Thus, [1969Vir] concluded that the phase in fact did

not belong to the Al-Mo-Ni system, but was easily stabilized by small amounts of impurities (low purity

99.8 mass% nickel was used for preparation of the specimens!). The phase was indexed as a MgZn2 type

Laves phase (a = 474, c = 770 pm). Two ternary compounds have been identified in the Al-rich range of

compositions, namely, 1, Mo(NixAl1-x)3 [1969Mar, 1969Vir, 1971Pry] and 2. The composition of 2 was

determined by [1969Mar] and [1971Pry] as Mo5Ni18Al77, but [2002Gru] corrected it and determined

crystal system and lattice parameters (Table 1). According to the findings of [1969Mar] at 800°C, [1971Pry]

at 600°C and [1965Ram, 1969Vir] at 950°C, the 1 phase is likely to exhibit a homogeneity range strongly

dependent on temperature. The compound Mo(Al2,75Ni0,25), observed in the aluminothermic preparation

of Al-Mo-Ni alloys from Mo- and Ni-oxides, is likely to be isotypic with the TiAl3 type, despite the fact

that the c-axis corresponds to a twofold superstructure [1969Rec]. [1984Och1, 1984Och2, 1985Mis,

1988Mas, 1989Hon1, 1989Hon2] investigated the influence of the addition of a third element on the lattice

parameter change of binary Ni solid solutions. The temperature dependence of the solid solubility is

reflected in the isothermal sections (see below) and additional information on the / ´ boundary may be

obtained from [1989Gai] and [1989Hon1]. A model based on X-ray measurements to show the effect of Mo

on the ´ structure has been suggested by [1977Aig].

267


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On the basis of X-ray, DTA as well as optical microscopy data [1986Mas2] plotted the phase diagram of

the partial pseudobinary NiAl-Mo system. The temperature of the L + equilibrium is equal to

1600 7°C; the maximum solubility of molybdenum in the phase is less than 4 at.%. The eutectic point is

placed at 10 at.% Mo, and its co-ordinates were later confirmed by [1991Sas]. Part of the quasibinary

section in the range of compositions 0 to 20 at.% Mo is presented in Fig. 1, with small changes according

to the melting temperature of the phase at 1651°C [2003Sal], whereas 1638°C was accepted by

[1986Mas2]. Similar compositions of the eutectic point were reported earlier by [1970Cli] (9 at.% Mo) and

[1971Pry] (10 at.% Mo), but considerably lower eutectic temperatures were presented (1427 and 1290°C,

respectively). According to the conclusion of [1993Kub], in view of the high melting temperature of the

phase and the reaction temperature of U1 (1340°C), the higher pseudobinary eutectic temperature (1600°C)

is recommended.

The vertical section Mo-Ni3Al, according to the data of [1971Pry], demonstrates a peritectoid reaction

+ ´+ , which is in contradictions to the observation of a eutectic solidification behavior in this area by

[1976Spr] and [1983Nas].


The reaction scheme of the partial Mo-NiAl-Ni system is presented in Fig. 2. One invariant three-phase

equilibrium as well as six invariant four-phase reactions have hitherto been observed in the system.

[1977Pea] reported the equilibrium L + + at 1300°C, but later it was established by [1977Aig,

1983Nas] and [1983Wak] that instead of the phase the ´ phase takes part in the eutectic reaction

L + + ´, and the various authors merely agree on the temperature of this reaction at 1300°C [1977Pea,

1983Wak, 1988Mas]. The reaction temperatures in Fig. 2 were measured and selected by [1986Mas1,

1986Mas2] and [1988Mas]. Table 2 presents the compositions of phases taking part in the invariant

equilibria, estimated on the basis of isothermal sections as well as on the data calculated by [1987Sve] and

experimentally determined by [1986Mas1] and [1986Mas2].

Liquidus Surface

Liquidus surface projection of the Ni-rich region (the Ni-NiAl-Mo partial system) is presented in Fig. 3. It

consists of five fields of primary crystallization corresponding to the , , , ´ and phases. It has been

constructed on the basis of constitution of the accepted binary phase diagrams and critically assessed

experimental data of different authors. So, the position of the U2E monovariant curve is established using

experimental data of [1977Pea, 1984Kov1] and [1984Kov2] on directionally solidified + eutectic

superalloys containing 8.58Al-27.22Mo (at.%) up to 18.66Al-15.50Mo (at.%) and 14.38Al-20.03Mo

(at.%), as well as data for two alloys, crystallized by [1987Sve] using the Bridgeman method. The position

of the U2 invariant point was accepted on the basis of data by [1977Pea] (Table 2), because data by

[1987Sve] do not agree with the estimated compositions of the , and phases participating in the

equilibrium LU2+ + at 1310°C. [1974Tho, 1976Nes, 1976Hen] and [1976Spr] discovered by

directional solidification studies the existence of eutectic reactions L + ´, L + and L + . The

Ni-NiAl-Mo liquidus surface projection was proposed by [1986Mas2] based on a rather schematic

projection given by [1983Nas], but the constitution of the Al-Ni binary as used by [1983Nas] contradicts

the assessment of [2003Sal]. A mathematical model was used to construct the isotherms at 1360 and

1340°C, as well as the monovariant curves p2U2 and U2E [1987Sve, 1989Gai]. Results of thermodynamic

calculations of the liquidus surface, carried out by [1999Lu] and based on the experimental data of

[1987Sve], were also used in our assessment, except the position of U2, which is placed by [1999Lu] at a

smaller Mo content. In Fig. 3 isotherms illustrating the shape of the surface, are added, in particular, the

isotherms at 1415, 1425 and 1445°C, using the experimental data of [1978Gul].

268


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Al–Mo–Ni

Isothermal Sections

Partial isothermal section at 1260°C is presented in Fig. 4 according to the data of [1984Mir]. Isothermal

sections at 1200°C constructed from the experimental results by [1983Nas] and [1988Mas] are in good

agreement with each other and with the calculation performed by [1974Kau], see Fig. 5. The character of

phase equilibria in the Ni-rich corner is similar to the character of the assessed equilibria, but the solubility

of Mo in the ´ phase is much smaller than experimental data, and also the position of the ´+ two-phase

region is different.

Phase equilibria at 1100°C [1988Mas] and 1000°C [2002Gru] are shown in Figs. 6 and 7, respectively.

Figure 8 presents a combination of data at close temperatures: for 1050°C in the Al-rich range by [2002Gru]

and for 1038°C in the Al-poor range [1984Mir]. Phase equilibria at 927°C [1984Mir] for Al-poor range and

at 950°C [2002Gru] for Al-rich corner are merged in Fig. 9. Partial isothermal sections at 880 (Fig. 10) and

700°C (Fig. 11) are accepted from [1988Mas].

In the assessed isothermal sections some minor modifications have been made taking into account the newly

determined position of the / ´ boundary according to SEM and DTA data by various groups (see “Solid

Phases”) and according to the constitution of the boundary systems. In particular, Fig. 11 reflects the

participation of the later determined Ni5Al3 phase in the equilibria at 700°C. [1989Hon1, 1989Hon2] and

[1991Mis] confirmed that the extent of the (Ni) solid solution area increases with rising temperature. The

position of the nickel-rich boundary of the ´ phase at 1200°C, calculated by [1991Eno] using the cluster

variation method, CVM (which utilizes the tetrahedron approximation and the phenomenological

Lennard-Jones pair interaction potential), practically coincides with the data of [1983Nas].


Figure 12 shows the partial isopleth at 14 at.% Al for a Ni content changing from 58 to 86 at.% according

to the data of [1989Mas]. This isopleth crosses two volumes of primary crystallization, corresponding to the

and phases, and four planes of invariant four-phase equilibria, in one of which (at 1300°C) the liquid

phase takes part, and the others are with participation of only the solid phases (at 1130, 890 and

730°C). The partial isopleth at 65 at.% Ni with Mo content changing from 15 to 35 at.%, as constructed

by [1983Wak], and the isopleths Mo60Al40 - Ni, Mo45Al55 - Ni constructed by [1986Mas2] do not comply

with the assessed liquidus surface.

Thermodynamics

No experimental thermodynamic data concerning the Al-Mo-Ni system are published in literature.

[1974Kau] calculated the isothermal sections at 1727, 1527 and 1200°C using symmetrical functions for the

excess free energies of mixing. There is a substantial disagreement between the calculated and the

experimental data; moreover constitution of the calculated binaries Al-Ni, Al-Mo and Mo-Ni contradict the

phase diagrams accepted in this assessment.

[1999Kau] and [1999Lu] assessed the experimental phase equilibria data in order to evaluate the

thermodynamic parameters of the ternary system by means of the CALPHAD method. A

substitutional-solution model is used to describe liquid, face-centered cubic (fcc) and body-centered cubic

(bcc) phases, while a sublattice model is used to describe the intermetallic phases. Two sets of

thermodynamic descriptions have been obtained, and comparison has been made between them. There is

satisfactory agreement between the calculations and experimental data. But phase diagrams of the boundary

systems Al-Ni, Mo-Ni and Al-Mo, accepted by [1999Kau] and [1999Lu], disagree to some extent with the

phase diagrams accepted in this assessment. [2000Bor] presented a general survey of the

diffusion-controlled transformations (DICTRA) software as an engineering tool for diffusion simulations

in multicomponent alloys. The model of coarsening of the ´ phase particles in ternary Al-Mo-Ni alloys was

used. In the calculation, the alloy composition was adjusted in order to have the same fraction of the ´

phase as experimentally observed. This gave a small difference in composition compared with the

experimental data.

269


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Al–Mo–Ni


The Al-Ni alloys with the addition of a refractory metal element (in particular, molybdenum) are interesting

as materials for production of in situ composites of eutectic superalloys that can serve, in particular, as

materials for specific hot section components of turbine engines, primarily blades or buckets and vanes as

well as nozzles [1976Jac]. In spite of very complicated alloys compositions, commercial superalloys

generally consist mainly of two phases, namely, and ´. The phase has potential applications such as

hot sections of gas turbine engines for aircraft propulsion systems, coats under thermal barrier coating,

electronic metallization compounds in advanced semiconductors [1998Mur] as well as surface catalysts

[1971Nal, 1998Mur].

The influence of molybdenum additions on the structure and hardness of the ´ phase based alloys has been

studied by [1959Gua1], and three general effects have been observed, namely, solid-solution hardening,

strain aging, and defect hardening arising from deviations from stoichiometry. A method for the

determination of site preference of substitutional elements in intermetallic compounds was proposed by

[2001Ter], and it was demonstrated that in the ´ based alloys molybdenum substitutes preferentially for

aluminium. The microstructure and chemical characteristics of the nanocrystalline phase are studied by

[2002Alb]. It was established that the addition of molybdenum tends to slightly refine the grain size of the

phase based alloy. The specimens with 2, 4 or 6 at.% Mo are polycrystalline containing, at the same time,

the phase, Ni, Al and Mo. The cast alloy NiAl-9Mo (at.%), prepared by [2002Guo], exhibited typical

deformation characteristics shown in conventionally superplastic materials, and possessed finely grained

structure. Properties of the directionally solidified eutectic superalloys were investigated by [1973Wal,

1974Tho, 1976Jac, 1976Nes, 1976Spr, 1981Sch, 1984Sch, 1985Nas, 1986Hor, 1986Kau] and [1987Sve].

It is shown that composites formed by directional eutectic solidification combined with a reinforcing

phase in the form of fibers have a considerable advantage over conventional superalloys [1973Wal,

1974Tho, 1976Jac, 1976Nes, 1976Spr]. Since the microstructure derives directly from the melt, the

composites are extremely stable at elevated temperatures. In addition improved oxidation and creep

resistances have been observed [1981Sch, 1985Nas]. The characteristic microstructure of the alloys consists

of / ´ matrix reinforced with faceted Mo fibres, which are primarily square in cross section and with the

following orientation relationship: [001] ´//[001] , (010) //(010) ´//(110) , (100) ´//(110) [1981Sch,

1984Sch, 1986Hor, 1986Kau, 1987Sve]. Precipitation in Ni-rich Al-Mo-Ni alloys has been investigated in

the temperature range 600 to 1100°C by transmission electron microscopy, selected-area electron

diffraction and hardness measurements [1987San]. Various stable and metastable phases ( , ´ and MoNi,

MoNi2 (MoPt2 type), MoNi3 (TiAl3 type), MoNi4, MoNi8 and SRO) have been observed and the ranges of

alloy composition and aging temperature for which each phase is formed have been determined and their

strengthening influence on the mechanical properties of ternary Al-Mo-Ni alloys has been discussed

[1987San]. Convergent-beam electron diffraction has been used by [1986Kau] to reveal local lattice

distortions in directionally solidified ´, Ni3Al type alloys with 12.8 at.% Al and 22.2 at.% Mo. TEM data

[1990Yam] from a Mo20Ni75Al5 alloy annealed at 800°C and quenched revealed the close-packed planes

of the ´´ and ´ phases to be parallel: [100] ´//[110] ´, (010) ´´//(111) ´ and [001] //[112] ´. [2001Kai]

studied the effect of molybdenum on the morphological stability of the interface between the ´ and

phases using Al-Mo-Ni ternary diffusion couples annealed at temperatures ranging from 900 to 1300°C.

Nonplanar interfaces with the Widmanstaetten-like structure were formed in the couples.

References

[1925Pfa] Pfautsch, H., “The System Aluminium-Molybdenum-Nickel” (in German), Z. Metallkd.,

19(4), 125-127 (1925) (Experimental, Equi. Diagram, 8) as quoted by [1993Kub]

[1933Roe] Roentgen, P., Koch, W., “Influence of Heavy Metals on Alloys of Aluminium” (in German),

Z. Metallkd., 25, 182-185 (1933) (Experimental, 8)

[1951Ham] Ham, J.L., “An Introduction to Arc-Cast Molybdenum and its Alloys”, Trans. Amer. Soc.

Mech. Eng. (ASME), 73, 723-732 (1951) (Experimental, 10) as quoted by [2003Sch]

270


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Al–Mo–Ni

[1952Tay] Taylor, A., Floyd, R.W., “The Constitution of Nickel-Rich Alloys of the

Nickel-Chromium-Aluminium System”, J. Inst. Met., 81, 451-464 (1952-1953)

(Experimental, Crys. Structure, Equi. Diagram, 15)

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[1958Woo] Wood, E.A., Compton, V.B., Matthias, B.T., Corenzwit, E., “ -Wolfram Structure of

Compounds Between Transition Elements and Aluminium, Gallium and Antimony”, Acta

Crystallogr., 11, 604-606 (1958) (Experimental, Crys. Structure, 13)

[1959Gua1] Guard, R.W., WestBrook, J.H., “Alloying Behavior of Ni3Al ( ´ phase)”, Trans. Met. Soc.

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Russian), Dokl. Akad. Nauk SSSR, 132, 339-342 (1960) (Experimental, Equi. Diagram, 14)

[1960Vig] Vigdorovich, V.N., Glazov, V.M., Glagoleva, N.N., “Investigation of the Solubility of Cr,

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Reaction”, Russ. Metall., (3), 61-63 (1969), translated from Izv. Akad. Nauk SSSR, Met., (3),


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Eutectic NiAl-Cr”, Metall. Trans., 1, 2907-2917 (1970) (Equi. Diagram, Experimental, 20)

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Networks with Lattice Mismatch in Eutectic Alloys”, Acta Metall., 19, 405-414 (1971)


[1971Nal] Nalibaev, T.N., Fasman, A.B., Inayatov, N.S., “Structure of Multicomponent Foraminate

Catalysts Based on Nickel”, Russian J. Phys. Chem., 45, 211-214 (1971), translated from

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[1971Pry] Pryakhina, L.I., Myasnikova, K.P., Markiv, V.Ya., Burnasheva, V.V., “Investigation of the

Molybdenum-Nickel-Aluminium Ternary System”, in “Phase Diagrams of Metal Systems”

(in Russian), Nauka, Moscow, 112-116 (1971) (Equi. Diagram, Experimental, 4)

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German), Z. Metallkd., 62, 844-848 (1971) (Experimental, Crys. Structure, Equi.

Diagram, 23)

271


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[1973Wal] Walter, J.L., Cline, H.E., “Stability of the Directionally Solidified Eutectics NiAl-Cr and

NiAl-Mo”, Metall. Trans., (4), 33-38 (1973) (Experimental, 10)

[1974Kau] Kaufman, L., Nesor, H., “Calculation of Superalloy Phase Diagrams, Part II”, Metall.

Trans., 5, 1623-1629 (1974) (Equi. Diagram, Thermodyn., 20)

[1974Tho] Thompson, E.R., Lemkey, F.D., “Metallic Matrix Composites”, Kreider, K.G. (Ed.),

Academic Press, New York, 101 (1974) as quoted by [1983Nas]

[1976Hen] Henry, M.F., “Precipitation of ' in - (Ni-Al-Mo) Eutectics”, Scr. Metall., 10, 955-957


[1976Jac] Jackson, M.R., Walter, J.L., “Superalloy Eutectic Composites with the VIA Refractory

Elements - Cr, Mo, W”, Superalloys-Metallurgy and Manufacture, AIME, N.Y., 341-350

(1976) (Review, Equi. Diagram, 42)

[1976Mon] Mondolfo, L.F., “Aluminium Alloys: Structure and Properties”, Butterworths, London,

598-599 (1976) (Review, 5) as quoted by [1993Kub]

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Properties of Some Directionally Crystallized Eutectics Based on Nickel”, Phys. Met.

Metallogr., 42, 117-123 (1976), translated from Fiz. Metall. Metalloved., 42, 1034-1041


[1976Spr] Sprenger, H., Richter, H., Nickel, J.J., “Directional Solidification of Ni-Mo-Al Eutectic

Alloys”, J. Mater. Sci., 11, 2075-2081 (1976) (Equi. Diagram, Experimental, 17)

[1977Aig] Aigeltinger, E.R., Bates, S.R., Gould, R.W., Hren, J.J., Rhines, F.N., “Phase Equilibria in

Rapidly Solidified Nickel-Rich Ni-Mo-Al Alloys”, Proc. Internat. Conf. Rapid

Solidification Processing. Principles and Technologies, Reston, Virginia, Claitor's

Publishing Div., Baton Rouge, 291-305 (1977), (Publ. 1978) (Equi. Diagram, Crys.

Structure, Experimental, Review, 20) as quotet by [1993Kub]

[1977Pea] Pearson, D.D., Lemkey, F.D., “Solidification and Properties of / ´- Mo Ductile/Ductile

Eutectic Superalloy”, Proc. Conf. Solidification and Casting of Metals, Metals Soc.

London, Sheffield, U.K., 526-532 (1977) (Publ. 1979) (Equi. Diagram, Experimental, 18)

as quoted by [1993Kub]

[1978Gul] Gulyaev, B.B., Grigorash, E. F., Efimova, M. N., “Solidification Range of Nickel Alloys”,

Heat-Resistant Steels and Alloys, (11), 914-917 (1978) translated from Metallov. Term.

Obrab. Met., 11, 34-37 (1978) (Experimental, Equi. Diagram, 8)

[1980Fer] Ferro, R., Marazza, R., “Crystal Structure and Density Data, Molybdenum Alloys and

Compounds other than Halides and Chalcogenides”, Atomic Energy Rev.: Spec. Iss. No.7,

IAEA, Vienna (1980) (Crys. Structure, Review, 961)

[1981Sch] Schwam, D., Dirnfeld, S.F., “Influence of Solidification Parameters on Microstructure of

/ '- (Mo) Eutectic Alloy”, Conf. Mat. Eng., Freund Publ. House Tel Aviv, 1981, 10-13

(1981) (Equi. Diagram, Experimental, 7) as quoted by [1993Kub]

[1982Shi] Shilo, I., Franzen, H.F., “High Temperature Thermodynamic Study of the

Molybdenum-Rich Regions of the Mo-Al System”, J. Electrochem. Soc., 129, 2613-2617


[1983Nas] Nash, P., Fielding, S., West, D.R.F., “Phase Equilibria in Nickel-Rich Ni-Al-Mo and

Ni-Al-W Alloys”, Met. Sci., 17(4), 192-194 (1983) (Equi. Diagram, Experimental, #, 20)

[1983Och] Ochiai, S., Oya, Y., Suzuki, T., “Solubility Data in Ni3Al with Ternary Additions”, Bull.

P.M.E. (T.I.T.), 52, 1-16 (1983) (Equi. Diagram, Review, Experimental, 4)

[1983Wak] Wakashima, K., Hoguchi, K., Suzuki, T., Umekawa, S., “Reinvestigation of Phase

Equilibria in the System Ni-Al-Mo and its Implication to the Elevated Temperature Stability

of / ' -Mo Aligned Eutectics”, Acta Metall., (11), 1937-1944 (1983) (Equi. Diagram,

Experimental, 19) as quoted by [1993Kub]

[1984Kov1] Kovacova, K., Kristin, J., “Morphological Properties of / '- Mo Eutectic Composite

Material” (in Czech), Kovove Mater., 22(3), 347-356 (1984) (Equi. Diagram, Experimental,

18) as quoted by [1993Kub]

272


MSIT®

Al–Mo–Ni

[1984Kov2] Kovacova, K., “Undirectional Solidification of Ni-Al-Mo Alloy”, J. Cryst. Growth, 66,


[1984Mir] Miracle, D.B., Lark, K.A., Srinivas, V., Lipsitt, H.A., “Nickel-Aluminium-Molybdenum

Phase Equilibria”, Metall. Trans. A, 15A, 481-486 (1984) (Equi. Diagram,


[1984Och1] Ochiai, S., Oya, Y., Suzuki, T., “Alloying Behavior of Ni3Al, Ni3Ga, Ni3Si and Ni3Ge”,

Acta Metall., 32, 289-298 (1984) (Equi. Diagram, Experimental, 90)

[1984Och2] Ochiai, S., Mishima, Y., Suzuki, T.S., “Lattice Parameter Data of Ni( ), Ni3Al ( ') and

Ni3Ga ( ') Solid Solutions”, Bull. P.M.E. (T.I.T.), 53, 15-28 (1984) (Crys. Structure,

Experimental, 66)

[1984Sch] Schwam, D., Dirnfeld, S.F., Nadiv, S., “Microstructural Instability of Ni-Mo-Al

Unidirectionally Solidified Eutectics”, J. Mat. Sci. Lett., (3), 363-366 (1984) (Experimental,

Crys. Structure, 6)

[1985Mis] Mishima, S., Ochiai, S., Suzuki, T.Y., “Lattice Parameters of Ni ( ), Ni3Al ( ') and Ni3Ga

( ') Solid Solutions with Additions of Transition and B-Subgroup Elements”, Acta Metall.,


[1985Nas] Nash, P., “Ni-Base Intermetallics for High Temperature Alloy Design “High-Temperature

Ordered Intermetallic Alloys”, Mat. Res. Soc. Symposia Proceedings, Kreh, C.C., Liu, C.T.,

Stoloff, N.S., (Eds.), MRS-Pennsylvania, Boston, Massachusetts (1984), 39, 423-427

(1985) (Equi. Diagram, Review, 15)

[1986Hor] Horita, Z., Sano, T., Nemoto, M., “Identification of Fine Particles in Unidirectionally

Solidified Ni-Al-Mo Eutectic Alloys by Means of EDX and SAD Analyses”, Acta Metall.,

34(8), 1525-1531 (1986) (Experimental, Crys. Structure, *, 27)

[1986Hua] Huang, S.C., Briant, C.L., Chang, K.-M., Taub, A.I., Hall, E.L., “Carbon Effects in Rapidly

Solidified Ni3Al”, J. Mater. Res., 1(1), 60-67 (1986) (Experimental, Mechan. Prop., 27)

[1986Kau] Kaufman, M.L., Pearson, D.D., Fraser, H.L., “The Use of Convergent beam Electron

Diffraction to Determine Local Lattice Distortions in Nickel Base Superalloys”, Philos.

Mag., A54, 79-92 (1986) (Crys. Structure, Experimental, 11)

[1986Mas1] Maslenkov, S.B., Udovskii, A.L., Burova, N.N., Rodimkina, V.A., “Phase Diagram of the

Nickel-Aluminium-Molybdenum System at 1300-2000°C”, Russ. Metall., (1), 203-209

(1986), translated from Izv. Akad. Nauk SSSR, Met., (1), 198-205 (1986) (Equi. Diagram,

Experimental, 9)

[1986Mas2] Maslenkov, S.B., Rodimkina, V.A., “Phase Equilibrium of the System Ni-Al-Mo in the

Composition Range Ni-NiAl-Mo”, Russ. Metall., (3), 215-220 (1986), translated from Izv.

Akad. Nauk SSSR, Met., (3), 218-223 (1986) (Equi. Diagram, Experimental, 7)

[1987Kha] Khadkikar, P.S., Vedula, K., “An Investigation of the Ni5Al3 Phase”, J. Mater. Res., 2(2),


[1987San] Sano, T., Nemoto, M., “Precipitates in Nickel-Rich Ni-Al-Mo Ternary Alloys”, Trans. Jpn.

Inst. Met., 28, 8-19 (1987), translated from J. Jpn. Inst. Met., 49(8), 690-698 (1985) (Crys.


[1987Sve] Svetlov, I.L., Udovski, A.L., Monastyrskaya, E.V., Oldakovskii, I.V., Nazarova, M.P.,

“Calculation of the Monovariant Liq/(Liq + + ) Line in the Ni-Mo-Al System and Plane

Front Solidification in / ´ - Alloys”, Russ. Metall., (6), 186-192 (1987), translated from

Izv. Akad. Nauk SSSR, Met., (6), 183-189 (1987) (Equi. Diagram, 14)

[1988Li] Li, X.Z., Kuo, K.H., “Decagonal Quasicrystals with Different Peridicities Along the

Tenfold Axis in Rapidly Solidified Al-Ni Alloys”, Phil. Mag. Lett., 58(3), 167-171 (1988)


[1988Mas] Maslenkov, S.B., Burova, N.N., Rodimkina, V.A., “The Ni-NiAl-Mo State Diagram in the

1200-700°C Temperature Range” (in Russian), Izv. Akad. Nauk SSSR, Met., (6), 183-190


[1989Ell] Ellner, M., Braun, J., Predel, B., “X-Ray Diffraction Investigation of Al-Cr Phases of the W

Family” (in German), Z. Metallkd., 80, 374-383 (1989) (Experimental, Crys. Structure, 38)

273


MSIT®

Al–Mo–Ni

[1989Gai] Gaidukov, A.M., Udovskii, A.L., Oldakovski, I.V., “Construction of Mathematical Models

in the Liquid System Surface” (in Russian), Dokl. Akad. Nauk SSSR, 305, 643-648 (1989)

(Review, Theory, 15)

[1989Hon1] Hong, Y.M., Nakajima, H., Mishima, Y., Suzuki, T., “The Solvus Surface in Ni-Al-X

(X = Cr, Mo and W) Ternary Systems”, ISIJ International, 29(1), 78-84 (1989) (Equi.


[1989Hon2] Hong, Y.M., Mishima, Y., Suzuki, T., “Accurate Determination of ' Solvus in Ni-Al-X

Ternary Systems”, Mat. Res. Soc. Symp. Proc., 733, 431-438 (1989) (Equi. Diagram,

Experimental, 35)

[1989Mas] Maslenkov, S.B., Rodimkina, V.A., “Phase Changes in Alloys of the System Ni-NiAl-Mo”

(in Russian), Izv. Akad. Nauk SSSR, Met., (1), 194-198 (1989) (Equi. Diagram, Crys.

Structure, #, 12)

[1990Yam] Yamamoto, M., Iada, J., Nenno, S., “The Microstructure of a Two-Phase Mixture in a

Ni75Mo20Al5 Alloy”, J. Mat. Sci. Lett., 9, 34-35 (1990) (Crys. Structure, Experimental, 5)

[1991Eno] Enomoto, M., Harada, H., Yamazaki, M., “Calculation of ´/ Equilibrium Phase

Compositions in Nickel-Base Superalloys by Cluster Variation Method”, Calphad, 15(2),

143-158 (1991) (Assessment, Calculation, Equi. Diagram, 34)

[1991Kim] Kim, Y.D., Wayman, C.M., “Transformation and Deformation Behavior of Thermoelastic

Martensite Ni-Al Alloys Produced by Powder Metallurgy Method” (in Korean), J. Korean

Inst. Met. Mater., 29(9), 960-966 (1991) (Mechan. Prop., Experimental, 15) as quoted by

[2003Sal]


Ternary Systems”, Mater. Sci. Eng. A, 146, 123-130 (1991) (Assessment, Equi. Diagram,

Experimental, 5)

[1991Sas] Sasaki, K., Morinaga, M., Yukawa, N., “Alloying Effect on the Solidified Structure of

NiAl”, Proc. Conf. Intermetal. Comp. - Struct. Mechan. Prop., 877-881 (1991) (Abstract,


[1991Sch] Schuster, J.C., Ipser, H., “The Al-Al8Mo3 Section of the Binary System Aluminium-

Molydenum”, Met. Trans., A22, 1729-1736 (1991) (Experimental, Crys. Structure, Equi.

Diagram, 20)

[1992Mur] Murakami, Y., Otsuka, K., Hanada, S., Watanabe, S., “Crystallography of Stress-Induced

B2 7R Martensitic Transformation in a Ni-37.0 at.% Al Alloy”, Mater. Trans. JIM, 33(3),


[1993Kha] Khadkikar, P.S., Locci, I.E., Vedula, K., Michal, G.M., “Transformation to Ni5Al3 in a 63.0

At. Pct Ni-Al Alloy”, Metall. Trans. A, 24A, 83-94 (1993) (Equi. Diagram, Crys. Structure,

Experimental, 28)

[1993Kub] Kubaschewski, O., “Al-Mo-Ni (Aluminium - Molybdenum - Nickel),” MSIT Ternary


International Services GmbH, Stuttgart; Document ID: 10.12789.1.20, (1993) (Crys.


[1994Mur] Murthy, A.S., Goo, E., “Triclinic Ni2Al Phase in 63.1 at.% NiAl”, Met. Mater., A, 25A(1),


[1995Gri] Grin, Y.N., Ellner, M., Peters, K., Schuster, J.C., “The Crystal Structures of Mo4Al17 and

Mo5Al22”, Z. Kristallogr., 210, 96-99 (1995) (Experimental, Crys. Structure, 11)

[1996Pau] Paufler, P., Faber, J., Zahn, G., “X-Ray Single Crystal Diffraction Investigation on

Ni1+xAl1-x”, Acta Crystallogr., Sect. A: Found. Crystallogr., A52, C319 (1996) (Crys.

Structure, Experimental, Abstract, 3)

[1996Vik] Viklund P., Haeussermann, U., Lidin, S., “NiAl3: a Structure Type of its Own?”, Acta

Crystallogr., Sect. A: Found Crystallogr., A52, C321 (1996) (Crys. Structure,

Experimental, Abstract)

[1997Bou] Bouche, K., Barbier, F., Coulet, A., “Phase Formation During Dissolution of Nickel in

Liquid Aluminium”, Z. Metallkd., 88(6), 446-451 (1997) (Thermodyn., Experimental, 15)

274


MSIT®

Al–Mo–Ni

[1997Jin] Jin, Y., Chaturvedi, M.C., Han, Y.F., Zhang, Y.G., “Crystal Structure of -NiMo Phase in

a Ternary Ni-Mo-Al Alloy”, Mater. Sci. Eng. A, A225, 78-84 (1997) (Crys. Structure,

Experimental, 13)

[1997Poh] Pohla, C., Ryder, P.L., “Crystalline and Quasicrystalline Phases in Rapidly Solidified Al-Ni

Alloys”, Acta Mater., 45, 2155-2166 (1997) (Experimental, Crys. Structure, 48)

[1997Pot] Potapov, P.L., Song, S.Y., Udovenko, V.A., Prokoshkin, S.D., “X-Ray Study of Phase

Transformations in Martensitic Ni-Al Alloys”, Metall. Mater. Trans. A, 28A, 1133-1142


[1998Mur] Murthy, B.S., Ranganathan, S., Int. Mater, Rev., 43(3), 101-141 (1998), as quoted by

[2002Alb]

[1998Rav] Ravelo, R., Aguilar, J., Baskes, M., Angelo, J.E., Fultz, B., Holian, B.L., “Free Energy and

Vibrational Entropy Difference between Ordered and Disordered Ni3Al”, Phys. Rev. B,

57(2), 862-869 (1998) (Thermodyn., Theory, Calculation, 43)

[1998Sim] Simonyan, A.V., Ponomarev, V.I., Khomenko, N.Yu., Vishnyakova, G.A., Gorshkov, V.A.,

Yukhvid, V.I., “Combustion Synthesis of Nickel Aluminides”, Inorg. Mater., 34(6),

558-561 (1998), translated from Neorgan. Mater., 34(6), 684-687 (1998) (Crys. Structure,

Experimental, 12)

[1999Kau] Kaufman, L., Dinsdale, A.T., “Summary of the Proceedings of the CALPHAD XXVII

Meeting, 17-22 May 1998, Beijing, China”, Calphad, 23(3-4), 265-303 (1999)

(Assessment, Calculation, Equi. Diagram, Thermodyn., #)

[1999Lu] Lu, X., Cui, Y., Jin, Z., “Experimental and Thermodynamic Investigation of the Ni-Al-Mo

System”, Metall. Mater. Trans. A, 30A, 1785-1795 (1999) (Equi. Diagram, Experimental,

Thermodyn., #, 28)

[2000Bor] Borgenstam, A., Engstroem, A., Hoeglund, L., Agren, J., “DICTRA, a Tool for Simulation

of Diffusional Transformations in Alloys”, J. Phase Equilib., 21(3), 269-280 (2000)

(Calculation, Kinetics, Thermodyn.)




[2001Kai] Kainuma, R., Ichinose, M., Ohnuma, I., Ishida, K., “Formation of '/ Interface

Morphologies in Ni-Al-X Ternary Diffusion Couples”, Mater. Sci. Eng. A, 312, 168-175

(2001) (Equi. Diagram, Experimental, Thermodyn., 21)

[2002Alb] Albiter, A., Bedolla, E., Perez, R., “Microstructure Characterization of the NiAl

Intermetallic Compound with Fe, Ga and Mo Additions Obtained by Mechanical Alloying”,

Mater. Sci. Eng. A, 328, 80-86 (2002) (Crys. Structure, Experimental, 14)

[2002Gru] Grushko, B., Mi, S., Highfield, J.G., “A Study of the Al-Rich Region of the Al-Ni-Mo Alloy

System”, J. Alloys Compd., 334, 187-191 (2002) (Crys. Structure, Equi. Diagram,

Experimental, 8)

[2002Guo] Guo, J.T., Du, X.H., Zhou, L.Z., Zhou, B.D., Qi, Y.H., Li, G.S., “Superplasticity in NiAl

and NiAl-Based Alloys”, J. Mater. Res., 17(9), 2346-2356 (2002) (Experimental, Mechan.

Prop., 17)





[2003Sch] Schuster, J.C., “Al-Mo (Aluminium - Molybdenum)”, MSIT Binary Evaluation Program, in


GmbH, Stuttgart; submitted for publication, (2003) (Crys. Structure, Equi. Diagram,

Assessment, 61)

275


MSIT®

Al–Mo–Ni


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Al)

< 660.452

MoxNiyAl1-x-y

cF4

Fm3m

Cu

a = 404.88 pure Al at 24°C [V-C]

x = 0, y = 0 to 0.004 [2003Sal]

y = 0, x = 0 to 0.00028 at 400°C [1960Vig]

y = 0, x = 0 to 0.00062 at 640°C [1960Vig]

y = 0, x = 0 to 0.0007 at 640°C, by

extrapolation [1960Vig]

, (Mo)

< 2623

,(Mo1-x-yNiyAlx)

cI2

Im3m

W

a = 314.7

a = 314.6

a = 314.5

pure Mo, at 25°C [V-C]

x = 0, y = 0.004 [1980Fer]

x = 0, y = 0.009 [1980Fer]

y = 0, x = 0 to 0.035, at 1000°C [1967Bel]

y = 0, x = 0 to 0.055, at 1205°C [1951Ham]

y = 0, x = 0 to 0.068, at 1316°C [1951Ham]

y = 0, x = 0 to 0.077, at 1317°C [1951Ham]

y = 0, x = 0 to 0.096, at 1482°C [1951Ham]

y = 0, x = 0 to 0.11, at 1572°C [1982Shi]

y = 0, x = 0 to 0.108, at 1600°C [1971Rex]

y = 0, x = 0 to 0.114, at 1604°C [1982Shi]

y = 0, x = 0 to 0.14, at 1700°C [1971Rex]

y = 0, x = 0 to 0.138, at 1748°C [1982Shi]

y = 0, x = 0 to 0.195, at ~2150°C [1951Ham]

, (Ni)

< 1455

, (MoxNi1-x)

, (Ni1-xAlx)

cF4

Fm3m

Cu

a = 352.40

a = 352.32

a = 355.8

a = 353.9

a = 355.2

a = 356.5

a = 356.3

a = 361.0

a = 352.8

a = 353.2

pure Ni at 25°C [1984Och2, Mas2]

pure Ni at 20°C [V-C]

quenched from 800°C [V-C]

x = 0.03, quenched from 1000°C [1984Och1,

1984Och2, 1985Mis]


1984Och2, 1985Mis]


1984Och2, 1985Mis]

x = 0.097 [1980Fer]

x = 0.218 [1980Fer]

x = 0 to 0.2 [2003Sal]

x = 0.2 at 1372°C [2003Sal]

x = 0.025 Slowly cooled alloy [1952Tay]

x = 0.05 Slowly cooled alloy [1952Tay]

, MoNi4< 870

tI10

I4/m

MoNi4

a = 572.0

c = 356.4

[V-C]

´´, MoNi3< 910

oP8

Pmmn

TiCu3

a = 506.4

b = 422.2

c = 444.8

[V-C]

276


MSIT®

Al–Mo–Ni

, MoNi

< 1362

(Mo,Ni,Al)1

(Ni,Mo,Al)1

oP112

Cmcm

MoNi

a = 910.8

b = 910.8

c = 885.2

a = 455

b = 1663

c = 873

46 to 48 at.% Ni [Mas2]

at 50.8 at.% Ni [1980Fer]

[1997Jin]

0 to 2 at.% Al [1969Vir]

0 to 1.6 at.% Al, T = 1260°C [1984Mir]

0 to 1.2 at.% Al, T = 1200°C [1988Mas]

0 to 1.1 at.% Al, T = 1171°C [1984Mir]

0 to 1.1 at.% Al, T = 1100°C [1988Mas]

0 to 0.6 at.% Al, T = 1093°C [1984Mir]

0 to 0.5 at.% Al, T = 1038°C [1984Mir]

0 to 0.3 at.% Al, T = 927°C [1984Mir]

MoAl12

< 712

cI26

Im3

WAl12

a = 757.3

a = 758.15

92.4 at.% Al [1991Sch]

[1954Ada]

[1980Fer]

MoAl5 (h2)

846 to 800 - 750

hP12

P63

WAl5

a = 491.2

c = 886.0

a = 489

c = 880

83.8 at.% Al [1991Sch]

[1980Fer]

MoAl5 (h1)

800 - 750 to ~648

hP60

P3

MoAl5 (h1)

a = 493.3

c = 4398

at 83.3 at.% Al [1991Sch]

MoAl5 (r)

648

hP36

R3c

MoAl5 (r)

a = 495.1

c = 2623

at 83.3 at.% Al [1991Sch]

Mo5Al22

964 to 831

oF216

Fdd2

Mo5Al22

a = 7382 3

b = 916.1 0.3

c = 493.2 0.2

81.7 at.% Al [1991Sch]

[1995Gri]

Mo4Al17

< 1034

mC84

C2

Mo4Al17

a = 915.8 0.1

b = 493.23 0.08

c = 2893.5 0.5

= 96.71 0.01

80.9 at.% Al [1991Sch]

[1995Gri]

MoAl41177 to 942

mC30

Cm

WAl4

a = 525.5 0.5

b = 1776.8 0.5

c = 522.5 0.5

= 100.88 0.06°

a = 525.5

b = 1176.8

c = 522.5

= 100.7°

79 to 80 at.% Al [1991Sch]

[1964Lea]

[1991Sch]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

277


MSIT®

Al–Mo–Ni

Mo1-xAl3+x

1260 to 1154

cP8

Pm3n

Cr3Si

a = 494.5 0.1

76 to 79 at.% Al [1991Sch]

[1991Sch]

MoAl31222 to ~818

mC32

C2/m

MoAl3

a = 1639.6 0.1

b = 359.4 0.1

c = 838.6 0.4

= 101.88 0.07

at 75 at.% Al [1991Sch]

Mo3Al8< 1555 10

mC22

Cm

Mo3Al8

a = 920.8 0.3

b = 363.78 0.03

c = 1006.5 0.3

= 100.78 0.05°

72 to 75 at.% Al [Mas2]

[1962For]

Mo2Al31570 to 1490

- - Called “ 1” (h) [1971Rex]

MoAl

1750 to 1470

cP2

Pm3m

CsCl a = 309.8

a = 309.8 to 309.9

46 to 51.7 at.% Al [Mas2]

Called “ 2” (h) [1971Rex]

[1971Rex]

[1980Fer]

Mo3Al

2150

(Mo,Ni,Al)3

(Al,Mo,Ni)1

cP8

Pm3n

Cr3Si

a = 495

a = 487.6

22 to 27 at.% Al [Mas2]

[1958Woo]

at 6 at.% Ni, 75 at.% Mo [1969Vir]

, NiAl3< 856

oP16

Pnma

NiAl3

oP16

Pnma

Fe3C

a = 661.15

b = 736.64

c = 481.18

a = 661.3 0.1

b = 736.7 0.1

c = 481.1 0.1

a = 659.8

b = 735.1

c = 480.2

[L-B]

[1996Vik]

[1997Bou, V-C]

Ni2Al3< 1138

hP5

P3m1

Ni2Al3

a = 403.63

c = 490.65

a = 402.8

c = 489.1

36.8 to 40.5 at.% Al [Mas2]

[L-B]

[1997Bou, V-C]

´, Ni3Al4< 702

cI112

Ia3d

Ni3Ga4

a = 1140.8 0.1 [1989Ell, V-C]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

278


MSIT®

Al–Mo–Ni

, NiAl

< 1651

(Ni,Mo,Al)1

(Al,Mo,Ni)1

cP2

Pm3m

CsCl

a = 287.04

a = 287.26

a = 286.0

a = 287.0

a = 288.72 0.02

a = 287.98 0.02

a = 289.0

a = 289.7

a = 290.4

a = 291.2

a = 291.9

a = 293.2

42 to 69.2 at.% Ni [Mas2]

57.7 at.% Ni [L-B]

46.6 at.% Ni [L-B]

[1987Kha]

63 at.% Ni [1993Kha]

50 at.% Ni [1996Pau]

54 at.% Ni [1996Pau]

[1971Cli]:

T = 0°C

T = 200°C

T = 400°C

T = 600°C

T = 800°C

T = 1000°C

0 to 1.5 at.% Mo, T = 1200°C [1983Nas]

0 to 0.3 at.% Mo, T = 1200°C [1988Mas]

0 to 0.2 at.% Mo, T = 1100°C [1988Mas]

0 to 4.0 at.% Mo, T = 1093°C [1984Mir]

Ni5Al3< 723

oC16

Cmmm

Pt5Ga3

a = 753

b = 661

c = 376

63 to 68 at.% Ni [1993Kha, Mas2]

at 63 at.% Ni [1993Kha]

´, Ni3Al

< 1372

Ni3(Al1-xMox)

cP4

Pm3m

Cu3Au

a = 356.6

a = 357.0

a = 356.77

a = 356.32

a = 357.92

a = 356.7

a = 357.0

a = 357.8

a = 356.8

a = 357.2

73 to 76 at.% Ni [Mas2]

[1952Tay]

[1984Och2, 1959Gua1]

[1986Hua]

disordered [1998Rav]

ordered [1998Rav]

at x = 0 [1963Arb] As scaled from diagram,

linear da/dx, alloys quenched from 1000C

[1984Och1, 1984Och2, 1985Mis]:

at x = 0

at 4 at.% Mo, 75 at.% Ni

at 1.5 at.% Mo, 75 at.% Ni [1963Arb]

at 1.5 at.% Mo, 73.5 at.% Ni [1963Arb]

0 to 4 at.% Mo, T = 1260°C [1984Mir]

0 to 4.6 at.% Mo, at 1200°C [1988Mas]

0 to 4.8 at.% Mo, at 1171°C [1984Mir]

0 to 4.9 at.% Mo, at 1100°C [1988Mas]

0 to 5.7 at.% Mo, at 1038 - 1093°C [1984Mir]

0 to 5 - 6 at.% Mo, at 1000°C [1977Aig,

1983Och, 1983Nas, 1984Och1, 1984Och2,

1984Mir, 1985Nas, 1985Mis, 1988Mas,

1989Hon1, 1989Mas, 1993Kub]

0 to 5.9 at.% Mo, at 927°C [1984Mir]

Ni2Al9 mP22

P21/c

Ni2Al9

a = 868.5 0.6

b = 623.2 0.4

c = 618.5 0.4

= 96.50 0.05°

Metastable;

[1988Li, 1997Poh]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

279


MSIT®

Al–Mo–Ni

NixAl1-x

0.60 < x < 0.68

tP4

P4/mmm

AuCu

m**

a = 383.0

c = 320.5

a = 379.5

c = 325.6

a = 379.5

c = 325.6

a = 379.5

c = 325.6

a = 379.9 to 380.4

c = 322.6 to 323.3

a = 371.7 to 376.8

c = 335.3 to 339.9

a = 378.00

c = 328.00

a = 418

b = 271

c = 1448

= 94.3°

Martensite, metastable

[1993Kha]

62.5 at.% Ni [1991Kim]

63.5 at.% Ni [1991Kim]

66.0 at.% Ni [1991Kim]

64 at.% Ni [1997Pot]

65 at.% Ni [1997Pot]

[1998Sim]

[1992Mur]

Ni2Al hP3

P3m1

CdI2

aP126

P1

a = 407

c = 499

a 1252

b 802

c 1526

90°

109.7°

90°

Metastable

[1993Kha]

[1994Mur]

D1 (Al-Ni) decagonal - Metastable [1988Li]

D4 (Al-Ni) decagonal - Metastable [1988Li]

* 1, Mo(NixAl1-x)3 tI8

I4/mmm

TiAl3

superstructure?

c = 2c0

a = 370.2

c = 836.1

a = 373.2

c = 843.0

a = 373

c = 1680

a = 376.1 0.6

c = 841.2 0.8

a = 373

c = 1680

[1971Pry], at 3 to 8 at.% Ni 25 at.% Mo,

600°C

at 1.6 to 6.0 at.% Ni [2002Gru] Called

“Mo2NiAl5” [1965Ram], from a three phase

alloy Mo25Ni25Al50

at 4 to 12 at.% Ni, 25 at.% Mo, 900°C

[1969Vir] From a three-phase alloy

Mo25Ni17Al58 [1969Vir] Called “N”

[2002Gru]

for Mo(Al2,75Ni0,25) [1969Rec]

from aluminothermic synthesis

* 2, Mo11Ni14Al75 orthorhombic a = 1005.4 0.4

b = 1528.8 0.4

c = 851.9 0.2

Called “X” [2002Gru]

Called “Mo5Ni18Al77” [1969Mar, 1971Pry]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

280


MSIT®

Al–Mo–Ni



Al Mo Ni

L + 1600 e1 (max) L 45

1

48.5

10

97.5

2.5

45

1.5

49

L + + ´ 1340 U1 L

´

20.5

<4

<33

<25

13

95.5

1.5

2

66.5

>0.5

>65.5

>73

L + + 1310 U2 L 8.58

<0.5

<2.5

<8

27.22

97.5

50.5

20

64.2

>2.0

>47

>72

L + + ´ 1300 E L

´

18

<2

<17

<20.5

16

96.0

10

5.5

66

>2

>73

>74

+ ´ + 1130 U3

´

1

2

10

20

98.2

49

14

5

0.8

49

76

75

+ ´+ ´´ 890 U4

´

´´

1

5.5

20.5

2.5

51

14

5.5

22.5

48

80.5

74

75

+ ´´ ´ + 730 U5

´

´´

5.5

20

3

1

11.5

5

22

19

83

75

75

80

281


MSIT®

Al–Mo–Ni

10

1100

1200

1300

1400

1500

1600

1700

1800

1900

Mo 20.00Ni 40.00Al 40.00

Mo 0.00Ni 50.00Al 50.00Mo, at.%

Te

mp

era

ture

, °C

L

1850°C

1651°C

1600±7°C

β

α+β

e1max

L+α

L+β

Fig. 1: Al-Mo-Ni.

Partial pseudobinary

system Mo-NiAl

Fig. 2: Al-Mo-Ni. Reaction scheme of the partial Mo-NiAl-Ni system

Mo-Ni

l + α δ1362 p

2

L α + β1600 e

1

Al-Mo-Ni

L + β α + γ´1340 U1

Al-Ni

l + γ γ´

1372 p1

l γ + δ1317 e

3

δ + γ γ´´

910 p3

γ + γ´´ θ870 p

4

l β + γ´

1369 e2

L + δ α + γ1310 U2

L α + γ + γ´ca.1300 E

δ + γ γ´ + γ´´ca.890 U4

γ + γ´´ γ´ + θca.730 U5

γ + α γ´ + δca.1130 U3

L+α+γ´

L+α+γ

δ+α+γ

α+γ+γ´

γ+γ´+γ´´

δ+γ+γ´

θ+γ´+γ´´ θ+γ+γ´

δ+γ´+γ´´

α+β+γ´

α+δ+γ´

282


MSIT®

Al–Mo–Ni

20

40

60

80

20 40 60 80

20

40

60

80

Mo Ni


Axes: at.%

α

e1max

U1

β

E

U2δ

γ

p2 e

3

p1

e2

1445

γ'

14251415

1369

1600

Fig. 3: Al-Mo-Ni.

Liquidus surface in

the Ni-rich region

20

40

60

80

20 40 60 80

20

40

60

80

Mo Ni


Axes: at.%

γ'

γ

δ

α

α+γ+γ'

α+δ+γ

Fig. 4: Al-Mo-Ni.

Partial isothermal

section at 1260°C

283


MSIT®

Al–Mo–Ni

20

40

60

80

20 40 60 80

20

40

60

80

Mo Ni


Axes: at.%

L

Mo3Al

8

Mo3Al

α

δ

γ

γ'

β

L+β+Mo3Al

8

Mo3Al+Mo

3Al

8+β

α+Mo3Al+β

α+γ'+β

γ'+γ+α

α+γ+δ

Fig. 5: Al-Mo-Ni.


1200°C, calculated by

[1974Kau]

20

40

60

80

20 40 60 80

20

40

60

80

Mo Ni


Axes: at.%

α

δ

γ

γ'

β

α+β+γ'

α+δ+γ'

δ+γ+γ'

Fig. 6: Al-Mo-Ni.

Partial isothermal

section at 1100°C

284


MSIT®

Al–Mo–Ni

10

20

30

10 20 30

70

80

90

Mo 40.00Ni 0.00Al 60.00

Mo 0.00Ni 40.00Al 60.00


Axes: at.%

L

τ1

τ2

Ni2Al

3

MoAl3

Mo3Al

8

τ1+τ

2+Ni

2Al

3

τ2+L+Ni

2Al

3

Fig. 7: Al-Mo-Ni.

Partial isothermal

section at 1000°C

20

40

60

80

20 40 60 80

20

40

60

80

Mo Ni


Axes: at.%

β

γ´

γα

δ

α+β+γ´

α+γ´+δ δ+γ+γ´

τ1

L

L+τ1+Ni

2Al

3

MoAl4

MoAl3

Mo3Al

8

Ni2Al

3

L+τ1

L+MoAl3

Fig. 8: Al-Mo-Ni.

Partial isothermal

section at 1038°C in

the Al-poor region

and at 1050°C in the

Al-rich region

285


MSIT®

Al–Mo–Ni

20

40

60

80

20 40 60 80

20

40

60

80

Mo Ni


Axes: at.%

τ2τ

1

MoAl4

MoAl3

Mo3Al

8

L

Ni2Al

3

β

αγ

δ

γ´α+β+γ´

δ+γ+γ´α+δ+γ´

Fig. 9: Al-Mo-Ni.

Partial isothermal

section at 927°C in

the Al-poor region

and at 950°C in the

Al-rich region

20

40

60

80

20 40 60 80

20

40

60

80

Mo Ni


Axes: at.%

β

γ'

γ

γ´´δα

α+γ'+δ δ+γ'+γ´´

α+β+γ'

γ´´+γ+γ'

Fig. 10: Al-Mo-Ni.

Partial isothermal

section at 880°C

286


MSIT®

Al–Mo–Ni

20

40

60

80

20 40 60 80

20

40

60

80

Mo Ni


Axes: at.%

β

γ'

γ

θγ´´δα

γ+γ'+θγ+γ'+γ´´

α+β+γ'

α+γ'+δδ+γ'+γ´´

Ni5Al

3

60 70 80

600

700

800

900

1000

1100

1200

1300

1400

1500

1600

Mo 28.00Ni 58.00Al 14.00

Mo 0.00Ni 86.00Al 14.00Ni, at.%

Te

mp

era

ture

, °C

L+α

L

L+γL+γ'+α

L+γ+α L+γ+γ'

γ+γ'+α

γ

γ'+δ+α

δ+γ'

γ+γ'+δ γ+γ'

890

γ+γ'+θ

γ+γ'+γ´´

γ'+γ´´

γ'+γ´´+θ

γ´´+γ'+δ

730

1300

1130

Fig. 11: Al-Mo-Ni.

Partial isothermal

section at 700°C

Fig. 12: Al-Mo-Ni.

Partial isopleth at 14

at.% Al

287


MSIT®

Al–Mo–Ti

Aluminium – Molybdenum – Titanium

Ludmila Tretyachenko

Literature Data

The phase equilibria in this system were assessed by [1993Bud] based on results published up to 1990.

Experimental data in these investigations have been interpreted from earlier versions of the Al-Ti phase

diagram, which essentially differs from those accepted at present. So, the earliest studies of the Ti rich alloys

did not take into account the Ti3Al based phase 2 [1958Boe, 1962Ere, 1962Ge, 1963Ge1, 1963Ge2,

1963Ge3]. Phase equilibria in the Ti rich alloys involving the 2 phase have been studied by [1969Cro,

1969Kor, 1969Nar, 1969Fed, 1978Ban, 1980Ban1, 1980Ban2]. However, the phase equilibria with phase

at temperatures higher than ~1200°C were not determined by [1958Boe, 1963Ge1, 1963Ge2, 1963Ge3,

1969Cro, 1969Kor, 1969Fed, 1978Ban, 1980Ban2, 1987Ere]. Phase equilibria involving phases, which

later were found to exist between TiAl and TiAl3 [1990Sch, 2001Bra], have not been considered by

[1962Ge, 1970Han, 1987Ere] and could not be shown in the isothermal sections at 1300 and 900°C given

by [1993Bud] in those days.

The following information on the Al-Mo-Ti system was available before 1990: an existence of a wide

region of bcc solid solutions and ordering of bcc solid solution that resulted in formation of the 2 CsCl type

phase with a composition of Ti2MoAl; significant solubilities of third component in some binary phases,

e.g. up to 20 at.% Mo in TiAl3 ( 1), 8 at.% Mo in TiAl ( ), more than 20 at.% in Mo3Al ( ) and small

solubilities of Mo in Ti3Al ( 2) and ( Ti). The ternary phase Ti1.5Mo1.5Al2 was found in alloys annealed

at 925°C for a week [1970Han]; the four-phase invariant equilibrium + 2 + 2 was suggested to exist at

550°C [1972Ham, 1973Ham, 1975Ham]. The partial ternary phase diagram in the Ti rich corner presented

by [1993Bud] has taken into account the coexistence of these four phases. A similar version of the phase

equilibria was used by [1981Tre] to describe phase transformations in Ti rich alloys.

The appearance of the 2+ + phase region [1980Ban2] has been discussed taking into account a version

of the Al-Ti phase diagram in which the 2 phase exists at high temperatures up to the melt.

The isothermal sections at 1600°C [1988Ere1] and 1300°C [1987Ere] were determined. An additional

investigation of the alloys resulted in an refining of some elements of the isothermal section at 1300°C, the

construction of the isothermal section at 1000°C, a preliminary version of the solidus surface and a reaction

scheme in the Al-Mo-Ti system up to 75 at.% Al. The reaction scheme takes into account new information

on the binary systems Al-Ti [1996Tre1] and Al-Mo [1991Sch]. The study was made using optical

microscopy (OM), later also electron microprobe (EMPA), X-ray diffraction (XRD) and differential

thermal (DTA) analyses [1988Ere2, 1990Ere, 1996Tre2].

The phase equilibria in the Al rich region of the Al-Mo-Ti system (> 65 at.% Al) have been studied by

[1994Sok] using OM and XRD and published as a partial isothermal section at 500°C. Crystallization of the

(Ti1-xMox)Al3 aluminides from dilute melts containing less than 0.5 at.% (Ti+Mo) was studied by

[1990Abd], who cooled very slowly from 1000 down to 700°C and then let the samples cool down to room

temperature inside a furnace.

Most of the investigations performed after the review by [1993Bud] concerned phase transformations and

microstructures of alloys adjacent to the Ti-Al side of the ternary phase diagram. The alloys based on Ti3Al

were studied by [1991Dja, 1992Dja1, 1992Dja2]. The alloys have been prepared by arc melting and, after

various heat treatments, were studied by means of OM, transmission electron microscopy (TEM), scanning

electron microscopy (SEM), selected area diffraction (SAD), anomalous small-angle X-ray scattering

(ASAXS). Mechanical properties were determined as well. The continuous cooling transformation

diagrams, from 1100°C down to room temperature were determined for different cooling rates and the phase

and structure transformations have been analyzed. The ( 2), at, 2´, 2 phases were observed.

Numerous investigations of AlTi based alloys have been carried out to obtain an information useful in

development of titanium aluminide alloys with improved mechanical properties and structural stability.

Such alloys have been studied using OM, XRD, and EMPA of arc melted, annealed at 1300°C for 5 h,

288


MSIT®

Al–Mo–Ti

1200°C for 48 h, 1100°C for 120 h and quenched alloys. Moreover in situ XRD at temperatures up to

1400°C [1992Kim] and studies of diffusion couples [1993Has, 1998Kim] have been made. The partial

isothermal sections were calculated [1998Kim] using the ThermoCalc program. The isothermal sections

were published for the region of the + + phase field at 1200 and 1300°C [1993Has, 1998Has, 1998Kim].

Earlier the thermodynamic calculations together with experimental studies of phase boundaries of the +

region were performed by [1975Zan, 1977Zan, 1986Gro, 1988Gro].

High temperature phase equilibria were studied by [1993Das1, 1993Das2] using OM, SEM, XRD, EMPA,

DTA and TEM of the Ti-50Al-5Mo and Ti-45Al-3Mo alloys; here and further compositions of alloys and

phases are given in at.%, if not stated differently. The location of the 2+ and + 2+ phase fields at

1175°C were determined using EMPA of the above alloys annealed at 1300°C for 3 d and then at 1175°C

for 6 d.

The microstructures of the Ti-48Al alloys containing 0.5 or 2 at.% Mo were studied as cast (plasma melted)

and quenched from temperatures between 1000 and 1350°C, by OM, SEM and TEM [1993Li].

Crystallographic analysis of the solidification microstructure of the Ti-48Al-2Mo alloy was used to

investigate high-temperature phase equilibria by OM, SEM, TEM, EMPA [1995Nak]. The mechanism of

phase transformations of the phase was studied on continuous cooling experiments.

The 2+ 2+ alloy, Ti-44Al-2Mo, prepared by plasma melting was studied by OM, TEM, SEM and

mechanical testing [1994Li, 1994Mor] on samples as cast, as HIPped (Hot Isostatic Pressed) at 1250°C, 150

MPa and as heat-treated at 1200 and 900°C for 120 and 500 h.

12 alloys containing 44 to 50 at.% Al and 2 to 6 at.% Mo were studied as cast and annealed in the

temperature range 1100 to 1400°C by means of TEM, XRD and EMPA [1997Sin1, 1997Sin2].

Solidification paths and postsolidification transformations were analyzed. Phases present after heat

treatments were determined and their compositions established. Partial Ti-rich isothermal sections at 1400,

1300 and 1200 - 1100°C were developed and projections of the liquidus and solidus surfaces involving ,

, and L phases were proposed.

A projection of the partial liquidus surface near Al-Ti side was constructed from microstructural analysis of

arc melted ingots of Ti alloys containing 45 to 60 at.% Al and 2 to 7.5 at.% Mo using OM, SEM [1998Joh].

The experimental data for the liquidus surface have been employed to calculate thermodynamically a

solidification path. There is a calculated partial isothermal section at 1500°C and a discussion on directional

solidification in the literature.

The partial liquidus surface in the regions of primary solidification of the and phases and directional

solidification of alloys have been analyzed by [2002Jun] too.

Two- and three-phase equilibria involving , 2, ( 2) and phases have been studied by [2000Kai] who

arc melted alloys, annealed them at 1000°C for 168 or 504 h, at 1200°C for 168 h and at 1300°C for 24 h

and characterized them by OM and EMPA. So partial phase diagrams at 1000, 1200 and 1300°C were

established. Similar phase relations were addressed by [1998Tak].

A detailed study of the Ti-50Al-15Mo alloy was made by [1997Che] using OM, XRD, SEM, TEM and

EMPA. The alloy was plasma arc melt and annealed at 1400°C for 1.5 h and at 1350°C for 2 h. The latter

samples were annealed additionally at 1200, 1000 or 800°C for 96, 144 and 504 h respectively and water

quenched after each of the heat treatments. The resulting phases, their compositions and crystal structure

were determined. In addition to the well known phases / 2 (hcp), / 2 (bcc/B2), (L10) and the phase

with D022 structure on the base of TiAl3, three new phases were reported and designated as L60, ´ and ´´.

The results by [1997Che] were used in the review by [1999Flo].

In Ti-(5.5-15)Mo-(2-7)Al (mass%) alloys, which were quenched from 1000°C, [1972Luz] studied

transformations during aging at 200 to 500°C and examined the influence of these transformations on the

mechanical properties.

[1980Sas] studied the crystal structures of martensites in Ti-(0-17)Mo-3Al (mass%) alloys quenched from

1000°C. The Ti-(0-30)Mo-3Al (mass%) were researched by [1971Kho] with respect to the

transformation temperature and the mechanical properties. Mechanical properties were also studied by

[1975Hid] together with the structure of the Ti-7Mo-(16,19)Al alloys, quenched from 960°C and aged at

600 and 400°C.

Physical properties and phase transformations were studied for the Ti3Al-1 % Mo alloy by [1976Zel].

289


MSIT®

Al–Mo–Ti

Site substitution behavior of Ti3Al and TiAl was calculated theoretically [1990Nan, 1993Rub, 1998Woo,

2000Yan, 2001Kan] and determined experimentally by [1999Hao] using the atom location channelling

enhanced microanalysis (ALCHEMI). The site occupancy of the alloying element in the and 2 phases

was used to estimate + 2 phase equilibrium and / 2 and 2/ phase boundaries [1999Yan, 2000Yan,

2001Kan]. The sublattice occupancy in B2 phases in the ALCHEMI experiments was analyzed by

[1995Che]. The local atomic order in the Ti2MoAl phase was determined from the EXAFS (Extended

X-Ray Absorption Fine Structure) study which revealed that this alloy has a pseudo-B2 structure, in which

Mo and Al atoms occupy one sublattice and Ti atoms the other one [1996Sik]. The relative stability of

different structures in the Ti50Mo25Al25 alloy was calculated theoretically by [2000Alo].

The stability of the aluminides Ti3Al, TiAl and the B2 phase in Ti2MoAl base alloys, has been considered

by [1992Nak, 1997Nak] as an information, which can be useful in developing Al-Mo-Ti based materials for

structural applications.

Binary Systems

The accepted Al-Mo system assessed by [2003Sch2] is based on the data of [1971Rex] for the Mo-Mo3Al8region and on the results of [1991Sch] for the Al rich part.

The Al-Ti phase diagram is accepted from the assessment of [2003Sch1], who has proposed a version based

on the results by [1992Kat, 1997Zha]. The TiAl-TiAl3 region shown by [1992Kat] summarizes complicated

phase relations in this concentration range as shown by [1990Sch] and recently reinvestigated by [2001Bra].

The data by [1996Tre1] are in good agreement with the results of recent studies, particularly as for the

Ti5Al11 phase. The Mo-Ti system is accepted as described by [Mas2].

Solid Phases

Data on solid phases observed in the ternary and relative binary systems are given in Table 1.

The bcc solid solutions existing in a wide range of compositions are the high temperature phase at the

Al-Ti side of the ternary phase diagram. They undergo a number of phase transformations as the

temperature decreases, 2, giving rise to a variety of microstructures depending on the temperature

of the heat treatment and on the cooling rate. Molybdenum is a strong stabilizer and its addition stabilizes

the bcc structure down to the room temperature [1991Dja]. Ordering of bcc solid solutions to ordered CsCl

type phase ( 2) was discovered by [1958Boe] and confirmed in the works of [1972Ham, 1991Dja,

1992Dja1, 1992Dja2, 1993Das1, 1993Das2, 1993Li, 1994Li, 1994Mor, 1995Nak, 1997Che, 1997Sin2].

Ordering takes place in a wide range of compositions. The temperature of ordering depends on the

composition of the phase and is supposed to be the highest at ~1400°C, for the composition Ti2MoAl.

An XRD study often is unable to recognize the ordered 2 phase owing to very weak superstructure

reflections. Therefore an electron diffraction analysis was used to identify the 2 phase [1993Das1].

The ternary phase detected by [1970Han] was confirmed by [1988Ere2, 1990Ere, 1996Tre2]. This

phase forms through a peritectoid reaction at ~1250°C.

The wide homogeneity range of the phase based on the binary TiAl3 compound earlier found by

[1970Han] was confirmed by [1987Ere, 1990Abd, 1990Ere, 1996Tre2]. The homogeneity range of TiAl3,

which is not more than ~1 at.% in the binary Al-Ti system, was found to extend up to ~22 at.% Mo at 75

at.% Al and up to ~16 at.% Mo along the 25 at.% Ti isopleths. The substitution of both Ti and Al atoms by

Mo atoms results in decreasing lattice parameters of the phase. The c/a ratio decreases insignificantly,

from 2.234 for TiAl3 to 2.214 for Ti3Mo22Al75, but the substitution of Al by Mo makes the c/a ratio

decrease to ~2.12.

The Mo solubility in the TiAl based phase increases with increasing Al content and reaches ~9 at.% at

~60 at.% Al. The lattice parameters of the phase were observed to decrease with c/a ratio increasing from

1.015 to ~1.035 with increasing Mo content [1990Ere, 1996Tre3].

[1997Che] observed a 1 phase with D022 type structure in the Ti-50Al-15Mo alloy which was annealed at

1200 - 800°C and water quenched. This work suggests that a transformation of the high temperature (L10)

phase to the 1 phase takes place, which can not be suppressed. Towards high Mo-contents in the phase

290


MSIT®

Al–Mo–Ti

region [1990Ere, 1996Tre3] observed a similar phase in alloys as cast and such annealed at 1300 and

1000°C.

The crystal structure of the ´ phase was supposed to be characterized by Mo and excess-Al ordering in Ti

layers in TiAl [1997Che]. The Mo solubilities in and phases were found to be small, about ~2 at.%

[1990Ere].

The ordered ´´ phase was observed by [1997Che] in the Ti-50Al-15Mo alloy after prolonged aging at

800°C. The crystal structure of the ´´ phase was found to be similar to that of TiAl3 (D022) and different

only by its sublattice. The proposed model of the ´´ phase is consistent with the chemical formula of

(Ti, Mo)3Al5. The ´´ phase was suggested to form from 2 (B2) phase or between 2 (B2) and ´ (D022).

The martensite phases ´, ´´ were observed in an alloy close to the Mo-Ti side of the ternary phase diagram

[1980Sas]. The metastable phase was reported by [1971Wil, 1972Ham, 1972Luz, 1980Sas] and also

observed in the research of [1991Dja, 1992Dja1, 1992Dja2] in Al-Ti base alloys of ~20 - 25 at.% Al and

3 - 4 at.% Mo, where also the 2´ martensite phase was observed which is based on Ti3Al. Additions of Al

to Mo-Ti alloys were found to suppress the formation of the phase [1972Luz].


The reaction scheme shown in Fig. 1a is based on results obtained by [1990Ere, 1996Tre2] mainly for the

Ti-TiAl3-MoAl3-Mo region. Temperatures of phase transformations were determined by DTA. Because of

the large losses of Al during heating at temperatures above ~1600°C, even for the time of an DTA

experiment, the temperature of the invariant equilibrium +L + 2 was developed from the Al-Mo binary

data and from temperatures determined on alloys of the nearest regions. As phase transformations in alloys

along the Al-Mo side could not be suppressed during cooling, only the phases existing at lower temperatures

have been observed. So, the equilibria involving the 1 and 2 phases were concluded to exist tentatively

from the analysis of results obtained from DTA, XRD and OM in as cast and annealed alloys. The reactions

in the region between the and phase fields is shown simplified because phase relations between the

phase (L10) and ´ (D022) are not determined.

The + ´+ (or + + / ´), + + , + + and + + phase fields were found to exist at 1000°C

[1990Ere], but the + + and + + phase fields were observed at 925°C by [1970Han]. So, the invariant

equilibria + + ´ and + ´ + (or summarized as + / ´+ ) were supposed to exist at

temperatures in the range of 925-1000°C.

According to [1972Ham, 1975Ham] the invariant equilibrium + 2 + 2 exists at 550°C in the Ti rich

region of the ternary system. However, the new version of the binary Mo-Ti phase diagram with a

monotectoid reaction + ´ existing at 675°C will lead to a three-phase region + + ´ in the ternary

system. It can be supposed that at lower temperature this three phase region and the + 2+ one will give

rise to the invariant four-phase equilibrium of + 2+ ´ rather than that proposed by [1972Ham,

1975Ham]. The phase taking part in this equilibrium may have an ordered B2 crystal structure.

Nevertheless, the invariant equilibrium + 2 + 2 suggested by [1972Ham, 1975Ham] takes place but at

a temperature between 675 and 850°C, which are the temperatures of the monotectoid reaction and the

maximum point of the binodal curve + ´ in the binary Mo-Ti system. One of the preceding three-phase

equilibria, 2+ + 2 may emerge from a contact of two-phase regions, 2+ ( 2) and + ´ based on the

+ ´ phase field in the Mo-Ti system (one of the phases may have the ordered B2 structure). One of the

equilibria succeeding the invariant equilibrium, + + ´, must move towards the binary Mo-Ti system down

to monotectoid line ´ at 675°C. The eutectoid reaction + 2+ ´ may be considered as one more

version of the invariant phase equilibrium in the Ti rich region of the ternary system.

Based on the data for the binary systems Al-Mo [1991Sch] and Al-Ti [2003Sch1] a tentative reaction

scheme for the Al rich region of the Al-Mo-Ti system is shown in Fig. 1b.

Liquidus and Solidus Surfaces

The solidus surface projected on the Ti-TiAl3-MoAl3-Mo region of the ternary system is shown in Fig. 2.

It mainly results from [1990Ere] and integrates additional data for the binary Al-Mo system by [1991Sch],

which were reported also by [1997Smi] and accepted by [2003Sch2].

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The liquidus surface has been determined only near the Al-Ti side. The projections of the boundaries

separating the fields of the primary crystallization of , and phases were constructed by [1997Sin1,

1998Joh]. An increasing stability of the phase or both the and phases with Mo addition was found. It

might be supposed that maxima on the boundary liquid curves corresponding to invariant three-phase

reactions L+ and L+ do exist. The boundary / liquid curve was established also by [2002Jun]

from analyzing the dendrite morphology of directionally solidified alloys. There is a good agreement

between the results obtained in the above studies.

The partial liquidus surface projection is shown in Fig. 3 with a maximum on the curve of the liquid

involving in the reaction L+ .

Earlier the liquidus surface was calculated by [1982Dan] using subregular solution approximation. The

calculation was performed without taking into account the existence of several Al-Ti and Al-Mo binary

phases.

Isothermal Sections

The isothermal section at 1600°C is shown in Fig. 4 [1988Ere1, 1988Ere2, 1990Ere, 1996Tre2]. Figure 5

shows the estimated partial section at 1500°C [1998Joh]. The tentative partial isothermal section at 1400°C

is given by Fig. 6 [1997Sin2]. The section was constructed from a study of 12 alloys annealed at 1400°C

for 1 h and quenched. Earlier the phase equilibria at 1400°C in the Ti rich region (Ti content > 50 mass%)

were reported by [1980Ban1], who has obtained similar results. Some discrepancies in phase boundaries

can be attributed to a different purity of alloys.

The phase equilibria at 1300°C are shown in Fig. 7 [1990Ere, 1996Tre2] and those between the , and

phases have been reported also by [1980Ban1, 1993Has, 1998Kim, 2000Kai]. A good agreement is

observed between obtained results.

Phase equilibria in the region between the and phase were not ascertained definitely. The Mo solubility

in (Ti5Al11) was found to be not more than ~1 at.%. The + + phase field was found to exist in a narrow

range at ~2 at.% Mo. The + equilibrium existing at higher Mo contents was observed to be replaced by

being in equilibrium with another phase. The crystal structure of this phase seems to be the same as that

of the phase, the D022 type, but with the c/a ratio close to 1.05, for a sublattice. A similar phase was

observed by [1997Che]. The phase relations involving this phase designated as ´ were not firmly

established and they are shown in Fig. 7 tentatively.

The phase equilibria between the , and phases at 1200°C were presented by [1980Ban1, 1993Has,

1997Sin2, 1998Has, 1998Kim, 2000Kai]. [1998Has, 1998Kim] have attempted to assess experimental

results by means of thermodynamic calculation. The partial phase diagram obtained for this region is shown

in Fig. 8 [2000Kai]. Similar phase diagrams were presented by [1993Has, 1998Has, 1998Kim] but another

location of apices of the + + phase triangle was proposed by [1997Sin2], especially for the and

phases. An ordered phase has not been detected by [1980Ban1, 1987Ere, 1990Ere], while the more recent

works have shown the ordered modification 2 of bcc solid solution [1992Kim, 1993Has, 1994Mor,

1995Nak, 1997Che, 1997Sin2, 1998Kim, 1998Tak].

The location of the three-phase + 2+ triangle at 1175°C established by [1993Das1, 1993Das2] is

consistent with that presented by [1993Has, 1998Has, 1998Kim, 2000Kai] for 1200°C.

The phase equilibria in the Ti rich alloys, i.e. with Ti content > 50 mass% have been presented by [1963Ge1,

1963Ge2, 1980Ban2, 1997Sin2]. The 2 phase was not identified in the earlier works, so the phase instead

of 2 was shown to coexist with the phase [1963Ge1, 1963Ge2]. The and 2 phases have not been

separated by [1980Ban2]. [1997Sin2] has reported that the Ti based phase which coexists with the and

2+ phases is an ordered 2 phase. A good agreement is observed as to the phase composition of the

studied alloys and compositions of the ( 2) and phases but there is a great difference between the

composition of the phase reported by [1997Sin2] and that shown by [1963Ge1, 1963Ge2, 1980Ban2].

Figure 9 shows the partial isothermal section at 1100°C developed mainly from that shown by [1963Ge1,

1963Ge2] and the accepted in this evaluation Al-Ti binary system. The data by [1963Ge1, 1963Ge2] were

preferred because a large number of alloys annealed at 1100°C for 100 h and water quenched were

investigated, while [1997Sin2] studied 12 alloys in the narrow composition range (44 to 50 at.% Al, 2 to

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Al–Mo–Ti

6 at.% Mo) and annealed only 6 h at 1100°C and [1980Ban2] examined only 4 alloys of the same

composition range.

The phase equilibria at 1000°C are shown in Fig. 10 [1988Ere2, 1990Ere, 1996Tre2]. A main peculiarity of

these phase equilibria is the ternary phase with a composition close to Ti3Mo3Al4. In the region between

the and phase fields, a coexistence of the phase and another phase with the D022 type crystal structure

( ´) was observed at a higher contents of molybdenum in the alloys. Phase relations between this ´ phase

and the and phase were not established. The coordinates of the 2+ / 2+ phase field agree well with

those determined by [2000Kai] as well as with the data by [1997Che] for the / 2 and phases in the

Ti-50Al-15Mo alloy. The region of the ordered 2 phase is shown mainly by [1958Boe], whose data are in

good agreement with those of [1993Das1, 1994Li, 1994Mor, 1997Sin2, 1997Che, 1997Nak].

The phase equilibria at 925°C are shown in Fig. 11 mainly from [1970Han] with corrections due to recent

data on the binary Al-Ti and Al-Mo and the ternary systems. The phase relations involving the phase,

which was discovered by [1970Han], are distinguished from those found at 1000°C. So, the invariant

reaction / 2+ + is supposed to take place at a temperature between 1000 and 925°C. The existence of

an + 2 phase field seems to be hardly probable as the ordering transformation 2 is believed to be of

second order. The region of the 2 phase is shown tentatively, the two-phase + 2 phase field is omitted in

Fig. 11. At lower temperatures the two-phase + 2 field would be possible, if attributed to a miscibility gap.

The phase equilibria at 800°C have been presented in the Ti rich part of the phase diagram by [1963Ge1,

1963Ge2]. The structure of alloys in the region of Ti-Al-(Ti ~30Mo) annealed at 800°C for 200 to 220 h

have been investigated by [1990Ere, 1996Tre2]. The partial section at 800°C shown in Fig. 12 was

constructed from the above works and information reported by [1958Boe, 1963Luz, 1972Ham, 1978Ban].

[1997Che] has reported the equilibria of the Ti-50Al-15Mo alloy annealed at 1350°C for 2 h, then at 800°C

for 504 h and water cooled. The alloy was found to consist of the 2+ ´+ ´´ phases, an information which

is not consistent with the isothermal section at 925°C shown above, because the equilibrium

2+ ´(TiAl)+ ´´(TiAl3) excludes the existence of the + + phase field, which was found earlier by

[1970Han] at 925°C.

The phase equilibria in the Ti rich corner at 700 and 600°C are similar to those at 800°C as it is shown in

Figs. 13 and 14 which incorporate compatibly data from [1972Ham, 1962Ge, 1963Ge1, 1963Ge2]

respectively, the binary Al-Ti and Mo-Ti phase diagrams, data by [1958Boe] for the / 2 boundary and data

by [1990Ere].

[1994Sok] studied the part of the system and published a partial isothermal section at 500°C for the Al rich

region; the Mo solubility in TiAl3 were found to be only 2 at.%; TiAl3 was found to coexist with MoAl3,

MoAl5 and MoAl12. The Ti solubilities in above aluminides were reported to be 2, 4 and 2 at.%,

respectively, but it is unknown what modification of MoAl3 was implied, no information on crystal

structures of the phases was reported. The presented data are not consistent with the data by [1990Abd],

who obtained the (Ti1-xMox)Al3 aluminides during very slow cooling, which allowed the equilibrium phase

to crystallize from the Al melt containing ~0.5 at.% (Ti+Mo). The (Ti1-xMox)Al3 aluminides with the TiAl3type structure were obtained up to x = 0.47 (~12 at.% Mo). The phase composition of +(Al) was found for

alloys in the Al-MoAl3-TiAl3 region almost up to x = 0.7 (the alloys were annealed at 600°C for 44 h,

solidus temperatures of these alloys were determined to be 650°C) [1990Ere].

Thermodynamics

Evaluated thermodynamic parameters used to asses isothermal sections in Al-Mo-Ti system by [1998Kim].

For modelling of individual phases the sublattice concept was applied. The calculated energy of formation

and the chemical potentials of elements, including that of Mo, in (TiAl) are given by [1998Woo] for low

temperatures and stoichiometric compositions.

The energy of formation for the A2 and B2 phases in the Ti50Al25Mo25 composition was evaluated by

[2000Alo].

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Ti alloys of the Al-Mo-Ti system are characterized by a variety of phase transformations, which can take

place during different heat treatments. Depending on a composition and a heat treatment, both equilibrium

and metastable phases can occur and result in various microstructures having an influence upon their

properties.

Mechanical properties of Ti based Al-Mo-Ti alloys have been studied depending on a composition and heat

treatments in earlier works. Composition - hardness relations of Ti rich alloys (Ti > 50 mass%) quenched

from various temperatures have been determined by [1963Ge1]. The maximum hardness (HV = 400 to 500

kg mm-2)has been observed for + and + + 2 alloys, solid solutions exhibited the minimum hardness,

(HV = 250 kg mm-2), at ~15 mass% Al near Ti3Al. High temperature hardness of Ti-1Mo-(5 to 20)Al

(mass%) has been determined by [1962Ge]. [1971Kho] studied the influence of a thermomechanical

treatment on tensile properties of Ti-3Al-(0 to 30)Mo (mass%) alloys, the maximum strengthening was

obtained for the Ti-3Al-15Mo alloy. [1972Luz] determined mechanical properties by tensile tests and

observed that additions of Mo gave rise to increasing strength in quenched alloys and suppressed the

formation of the phase. [1973Ham, 1975Hid] again observed a correlation between microstructures and

mechanical properties of Ti-(7-19)Al-7Mo. A study of Ti3Al based alloys containing up to 32 mass% Mo

was made by [1969Kor].

The TiAl and Ti3Al aluminides were a subject of recent investigations because they were found to have

potential use for high temperature applications in aerospace engines. These aluminides combine low

density, high specific strength, good resistance to oxidation, but they have low ductility at room

temperature. Molybdenum was found to be an alloying addition, which can have a favorable influence on

the properties of intermetallic alloys based on the Ti aluminides. [1991Mae] has found that Ti rich TiAl

modified by Mo exhibited higher tensile ductility at room temperature and improved creep strength. Room

temperature tensile tests have been carried out also by [1994Li, 1994Mor]. The high strength obtained at

room temperature for TiAl based alloys has been attributed to the presence of the ordered 2 phase.

Hardness measurements were carried out on the individual phases. The hardness values were measured to

be H 2 = 394 15 kg mm-2, H 2 = 430 20 kg mm-2 and H = 273 10 kg mm-2. Also 0.2 the stress

values at 0.2 % strain, the ductility , the maximum flow stress max were measured from tensile tests

[1994Mor]. Mechanical properties of TiAl based alloy at temperatures ranging from 77 to 1473 K were

examined by [1993Has]. The mechanical properties of TiAl can be greatly improved by control of

microstructure and morphology of secondary phases, which can be changed with Mo additions affecting the

stability of the phases.

Tensile properties of Ti3Al based alloys with Mo at room temperature have been examined on samples

thermomechanically processed (TMP) and heat treated (HT) [1992Dja2]. It has been shown that the tensile

properties of Al-Mo-Ti aluminides may be optimized by specific TMP and HT.

Electrical conductivity and a coefficient of thermal expansion in the temperature range from 20 to 1000°C,

hardness at 20 to 800°C, a modulus of elasticity and internal friction were measured on the Ti3Al-1 mass%

Mo alloy by [1976Zel]. An abrupt change of physical properties with a heat absorption has been observed

at 1080°C.

Calorimetric studies of superconducting (Ti0.75Mo0.25)1-xAlx alloys with x = 0 to 0.06 have revealed that

the superconducting transition temperature Tc decreases linearly from 3.9 0.1 K at x = 0 with a rate of

approximately 0.3 K per at.% Al [1985Ho].

Miscellaneous

Mo atoms tend to Ti sites [1999Hao, 2000Yan] in Ti3Al alloys. The Mo atoms were shown to occupy both

sublattices in TiAl [1990Nan, 1998Woo, 2000Yan, 2001Kan] and show different site preference of Mo in

TiAl alloys than in Ti3Al [1999Hao]. The formation of Ti3Al phase was shown to obey the electron

concentration rule. The experimental boundary of the 2 phase was found to agree with that calculated using

an electron model with N = 2.12 [1984Li].

[1995Che] studied two alloys, Ti-42Al-7.5Mo and Ti-50Al-15Mo, which were annealed at 1350°C for 2 h

and WQ, then the latter alloy was annealed at 800°C for 504 h and WQ. Compositions of three 2 phases

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of these alloys were determined (the first alloy was single phase 2) and sublattice occupancies were

established using ALCHEMI. The 2 phases were found to be Ti-48Al-15Mo, Ti-41.6Al-7.3Mo and

Ti-37.1Al-22.2Mo. The first two ones were described as (Ti,Mo)52Al48, Ti51(Al,Mo)49. In the third 2

phase, Mo was suggested to be distributed randomly on both sublattices, (Ti,Mo)50(Al,Mo)50. In all cases

2 contained more Al than Ti2AlMo [1958Boe].

Applying CVM, the cluster variation method [1993Rub] calculated from experimental binary data an

isothermal section at 1000°C and found a miscibility gap in the inner part of the section besides of bcc ( )

and B2 ( 2) fields. A comparison between the energy of formation of the A2 ( ) and B2 ( 2) phases of the

same composition Ti2AlMo calculated from first principles has shown the B2 phase to be more stable than

the A2 one [2000Alo].

[1980Sas] reported martensite phases in Ti-3Al-(0 to 17)Mo (mass%) alloys quenched from 1000°C (the

field). The crystal structure of the martensite at low Mo-content (4 mass%) was found to be hcp ( ´), at 7 -

12 mass% (3.5 - 6.2 at.%) Mo it was orthorhombic ( ´´). No martensite was observed at Mo contents higher

than 13 mass% (6.8 at.%). However, slight deformation caused orthorhombic martensite to occur at 13 to

17 mass% Mo. A distorted bcc phase was observed at 12 mass% Mo. [1971Wil] studied a decomposition

of a metastable phase in the alloys Ti-(3, 6)Al-20Mo (mass%) quenched from 1000°C and has found that

Mo additions reduced the volume fraction and time of stability of the phase. The influence of Mo

additions on the occurrence of the phase in the alloys containing 4 to 8 at.% Mo and 0 to 3 at.% Al was

studied by [1993Cui]. It was shown that formation of the phase obeys the electron concentration rule. The

- boundary was calculated and determined experimentally (at the valence electron number 4.10, from

~4.5 at.% Mo to ~6 at.% Mo at 0 and 3 at.% Al). A formation of an athermal phase (“tweed

microstructure”) has been observed in Ti3Al based alloys containing 3.4 and 4.4 at.% Mo quenched from

the field [1991Dja, 1992Dja2].

[1991Dja, 1992Dja2] have presented continuous cooling transformation diagrams for Ti3Al based alloys

with different Al and Mo contents, which have been annealed in the region and cooled with rates varying

from 80 to 0.1°C s-1.

The Ti-50Al-5Mo alloy was found to be single phase at 1400°C and to exhibit during cooling a sequence

of phase transformations + + + + 2. The + + phases were found in the alloy annealed at

1240°C for 150 h. The 2+ phase composition was established in the alloy annealed at 1175°C for 6 h. In

the Ti-45Al-5Mo alloy, the + structure observed in the alloy annealed at 1300°C for 3 d was found to be

changed to 2+ after annealing at 1175°C for 6 h. The 2+ alloys were found to be stable to a high

temperature exposure at 1240°C for 150 h, but some modifications took place at longer time.

The partitioning tendency of Mo into different phases ( , / 2 and ) was found to be as follows: > >

[2000Kai].

Sintering of elemental powders at 1150°C to obtain a ternary intermetallic compound of the L12 type has

resulted in the phase composition D022+(TiAl2) in the Ti-67Al-8Mo alloy [1993Nak].

Sulfidation properties of the TiAl-2Mo alloy at 900°C and 1.3 Pa sulphur pressure have been studied by

[2000Izu].

References

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Alloys by Step Aging”, Metall. Trans., 6A, 1541-1546 (1975) (Experimental, 7)

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[1990Abd] Abdel-Hamid, A.A., “Crystallization of Complex Aluminide Compounds from Dilute Al-Ti

Metals Containing One or Two Transition Metals of IVB to VIB Groups”, Z. Metallkd., 81,


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Structure of Alloys of the Ti-Mo-Al System” (in Russian), “Physico-Chemical Investigation

of Binary and Ternary Systems of Transition Metals of IV-VIII Groups of the Periodic

System and Development of Principles for Control of Mechanical Properties of Alloys on

Their Base (Theme 2.26.30, Final Report, State Regist. No. 01 86 0 060682)”, Akad. Nauk

Ukr. SSR, IPM, Kiev, 83-135, 141-143 (1990) (Crys. Structure, Equi. Diagram,




Review, 33)

[1991Dja] Djanarthany, S., Servant, S., Penelle, R., “Phase Transformations in Ti3Al and Ti3Al+Mo

Aluminides”, J. Mater. Res., 6, 969-986 (1991) (Crys. Structure, Equi. Diagram,

Experimental, 24)

[1991Mae] Maeda, T., Okada, M., Shida, Y., “Ductility and Strength in Mo Modified TiAl”, Mat. Res.

Soc. Symp. Proc., 213, 556-560 (1991) (Crys. Structure, Experimental, 15)

[1991Sch] Schuster, J., Ipser, H., “The Al-Al8Mo3 Section of the Binary System Aluminium-

Molybdenum”, Metall. Trans., 22A, 1729-1736 (1991) (Crys. Structure, Equi. Diagram,

Experimental, 20)

[1992Dja1] Djanarthany, S., Servant, S., Lyon, O., “Phase Separation in a Ti-Al-Mo Alloy Studied by

Anomalous Small-Angle X-Ray Scattering. A Synchrotron Radiation Experiment”, Philos.

Mag., 66A, 575-590 (1992) (Crys. Structure, Equi. Diagram, Experimental, Theory, 14)

[1992Dja2] Djanarthany, S., Servant, S., Penelle, R., “Influence of an Increasing Content of

Molybdenum on Phase Transformations of Ti-Al-Mo Aluminides - Relation with

Mechanical Properties”, Mater. Sci. Eng., A152, 48-53 (1992) (Equi. Diagram,

Experimental, 14)

[1992Kat] Kattner, U.R., Lin, J.C., Chang, Y.A., “Thermodynamic Assessment and Calculation of the

Ti-Al System”, Metall. Trans., 23A, 2081-2090 (1992) (Equi. Diagram, Review, Theory,

Thermodyn., 51)

[1992Kim] Kimura, M., Hashimoto, K., Morikawa, H., “Study on Phase Stability in Ti-Al-X Systems

at High Temperatures”, Mater. Sci. Eng., A152, 54-59 (1992) (Crys. Structure, Equi.


[1992Nak] Naka, S., Thomas, M., Khan, T., “Potential and Prospects of Some Intermetallic Compound

for Structural Applications”, Mater. Sci. Technol., 8, 291-298 (1992) (Equi. Diagram,

Experimental, Review, 26)

[1993Bud] Budberg, P., Schmid-Fetzer, R., “Aluminium - Molybdenum - Titanium”, MSIT Ternary




[1993Cui] Cui, Y., Li, D., Wan, X., “ Phase Formation in Ti Alloys” (in Chinese), Acta Metall. Sin.

(China), 29, A61-A67 (1993) (Crys. Structure, Experimental, Theory, 9)

[1993Das1] Das, S., Mishurda, J.C., Allen, W.P., Perepezko, J.H., Chumbley, L.S., “Development of a

( + 0) Lamellar Microstructure in a Ti45Al50Mo5 Alloy”, Scr. Metall. Mater., 28, 489-494


[1993Das2] Das, S., Jewett, T.J., Perepezko, J.H., “High Temperature Phase Equilibria of Some Ternary

Titanium Aluminides”, in “Structural Intermetallics”, Darolia, R., Lewandowski, J.J., Liu,

C.T., Martin, P.L., Miracle, D.B., Nathal, M.V., (Eds.), Min., Met., Mater. Soc., 420

298


MSIT®

Al–Mo–Ti

Commonwealth Dr., Warrendale, Pens. 15086, 35-43 (1993) (Equi. Diagram, Experimental,

Review, 48)

[1993Gam] Gama, S., “Aluminium - Niobium - Titanium”, MSIT Ternary Evaluation Program, in MSIT



Assessment, 22)

[1993Has] Hashimoto, K., Kimura, M., “Effects of Third Element Additions on Mechanical Properties

of TiAl”, in “Structural Intermetallics”, Darolia, R., Lewandowski, J.J., Liu, C.T.,

Martin, P.L., Miracle, D.B., Nathal, M.V., (Eds.), Min., Met., Mater. Soc., 420

Commonwealth Dr., Warrendale, Pens. 15086, 309-318 (1993) (Equi. Diagram,

Experimental, 18)

[1993Li] Li, Y.G., Loretto, M.H., “Antiphase Boundaries in Ti-48Al2Mo”, Acta Metall. Mater., 41,


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Alloys”, Intermetallics, 1, 41-48 (1993) (Equi. Diagram, Experimental, 40)

[1993Oka] Okamoto, H., “Al-Ti (Aluminium - Titanium)”, J. Phase Equilib., 14, 120-121 (1993)


[1993Rub] Rubin, G., Finel, A., “Calculation of Phase Diagrams of Ternary Systems with

Cluster-Variation-Method Entropy”, J. Phys.: Condens. Matter., 5, 9105-9120 (1993)

(Equi. Diagram, Theory, Thermodyn., 34)

[1994Kai] Kainuma, R., Palm, M., Inden, G., “Solid-Phase Equilibria in the Ti-Rich Part of Ti-Al

System”, Intermetallics, 2, 321-332, (1994) (Equi. Diagram, Experimental, 35)

[1994Li] Li, Y.G., Loretto, M..H., “Microstructure and Fracture Behaviour of Ti-44Al-xM

Derivatives”, Acta Metall. Mater., 42, 2913-2919 (1994) (Crys. Structure, Equi. Diagram,

Experimental, 12)

[1994Mor] Morris, M.A., Li, Y.G., Leboeuf, M., “Variation of the Phase Distribution in a Ti-44Al-2Mo

Alloy by Annealing: Influence on its Strength and Ductility”, Scr. Metall. Mater., 31,

449-454 (1994) (Crys. Structure, Equi. Diagram, Experimental, 11)

[1994Sok] Sokolovskaya, E.M., Kazakova, E.F., Poddyakova, E.I., Portnoy, V.K., Temirbayeva, A.A.,

“Isothermal Section of the Al-Mo-Ti System at 770 K” (in Russian), Vestn. Mosk. Univ.,

Ser. 2: Khim., 35, 95-97 (1994) (Equi. Diagram, Experimental, 6)

[1995Che] Chen, Z., Jones, I.P., “Sublattice Occupancy in Three Ti-Al-Mo B2 Phase”, Scr. Metall.

Mater., 32, 553-557 (1994) (Crys. Structure, Experimental, 5)

[1995Gri] Grin, Yu.N., Ellner, M., “The Crystal Structures of Mo4Al17 and Mo5Al22”, Z. Kristallogr.,


[1995Nak] Nakai, K., Ono, T., Ohtsubo, H., Ohmori, Y., “Phase Stability and Decomposition Processes

in Ti-Al Based Intermetallics”, Mater. Sci. Eng., A192, 922-929, (1995) (Equi. Diagram,

Experimental, 21)

[1996Sik] Sikora, T., Hug., G., Jaouen, M., Flank, A.-M., “EXAFS Study of the Local Atomic Order

in Ti2AlX (X = Nb, Mo) B2 Intermetallic Compounds”, J. Phys. IV, 6, C2-15-C2-20 (1996)


[1996Tre1] Tretyachenko, L.A., “On the Ti-Al System”, Phase Diagrams in Material Science, Fifth

International School, 1996, Katsyveli, Crimea, Ukraine, 118 (1996) (Equi. Diagram,

Experimental, #, 0)

[1996Tre2] Tretyachenko, L.A., “Phase Equilibria in the Ti-Mo-Al System”, Phase Diagrams in

Materials Science, Fifth International School, 1996, Katsyveli, Crimea, Ukraine, 119


[1996Tre3] Tretyachenko. L.A., unpublished data

[1997Bul] Bulanova, M., Tretyachenko, L., Golovkova, M., “Phase Equilibria in Ti-Rich Corner of the

Ti-Si-Al System”, Z. Metallkd., 88, 256-267 (1997) (Crys. Structure, Equi. Diagram,


299


MSIT®

Al–Mo–Ti

[1997Che] Chen, Z., Jones, I.P., Small, C.J., “The Structure of the Alloy Ti-50Al-15Mo between 800

and 1400°C”, Acta Mater., 45, 3801-3815 (1997) (Crys. Structure, Equi. Diagram,

Experimental, 18)

[1997Nak] Naka, S., Khan, T., “Designing Novel Multicomponent Intermetallics: Contribution of

Modern Alloy Theory in Developing Engineering Materials”, J. Phase Equilib., 18,


[1997Sau] Saunders, N., “The Al-Mo System (Aluminium - Molybdenum)”, J. Phase Equilib., 18,

370-376 (1997) (Crys. Structure, Equi. Diagram, Review, Thermodyn., 40)

[1997Sin1] Singh, A.K., Banerjee, D., “Transformations in 2+ Titanium Aluminide Alloys

Containing Molybdenum: Part I. Solidification Behavior”, Metall. Mater. Trans., 28A,


[1997Sin2] Singh, A.K., Banerjee, D., “Transformations in 2+ Titanium Aluminide Alloys

Containing Molybdenum: Part II. Heat Treatment”, Metall. Mater. Trans., 28A, 1745-1753


[1997Smi] Smith, J.F., “Appendix” to [1997Sau], J. Phase Equilib., 18, 376-378 (1997) (Crys.

Structure, Equi. Diagram, Review, 1)

[1997Zha] Zhang, F., Chen, S.L., Chang, Y.A., Kattner, U.R., “A Thermodynamic Description of the

Ti-Al System”, Intermetallics, 5, 471-482 (1997) (Equi. Diagram, Theory, Thermodyn., 45)

[1998Has] Hashimoto, K., Kimura, M., Mizuhara, Y., “Alloy Design of Gamma Titanium Aluminides

Based on Phase Diagrams”, Intermetallics, 6, 667-672 (1998) (Equi. Diagram,


[1998Joh] Johnson, D.R., Chihara, K., Inui, H., Yamaguchi, M., “Microstructural Control of

TiAl-M-B Alloys by Directional Solidification”, Acta Mater., 18, 6529-6540 (1998) (Equi.

Diagram, Experimental, Theory, Thermodyn., 33)

[1998Kim] Kimura, M., Hashimoto, K., “High-Temperature Phase Equilibria in Ti-Al-Mo System”,

J. Phase Equilib., 20, 224-230 (1998) (Equi. Diagram, Experimental, Theory,

Thermodyn., #, 19)

[1998Tak] Takeyama, M., Ohmura, Y., Kikuchi, M., Matsuo, T., “Phase Equilibria and Microstructural

Control of TiAl Based Alloys”, Intermetallics, 6, 643-646, (1998) (Equi. Diagram,

Experimental, Theory, Thermodyn., #, 33)

[1998Woo] Woodward, C., Kajihara, S., “Site Preferences and Formation Energies of Substitutional Si,

Nb, Mo, Ta and W Solid Solutions in L10 Ti-Al”, Phys. Rev. B, 57, 13459-13470 (1998)

(Crys. Structure, Theory, 45)


Elements in TiAl and Ti3Al Alloys”, Acta Mater., 47, 1129-1139 (1999) (Crys. Structure,

Experimental, 41)

[1999Flo] Flower, H.M., Christodoulou, J., “Phase Equilibria and Transformations in Titanium

Aluminides”, Mater. Sci. Technol., 15, 45-52 (Crys. Structure, Equi. Diagram, Review, 46)

[1999Yan] Yang, R., Hao, Y.L., “Estimation of ( + 2) Equilibrium in Two Phase Ti-Al-X Alloys by

Means of Sublattice Site Occupancies of X in TiAl and Ti3Al”, Scr. Mater., 41, 341-346


[2000Alo] Alonso, P.R., Rubiolo, G.H., “Relative Stability of bcc Structures in Ternary Alloys with

Ti50Al25Mo25 Composition”, Phys. Rev. B, 62, 237-242 (2000) (Crys. Structure, Equi.

Diagram, Theory, 19)

[2000Izu] Izumi, T., Yoshika, T., Hayashi, S., Narita, T., “Sulfidation Properties of TiAl-2 at.% X

(X = V, Fe, Co, Cu, Mo, Nb, Ag and W) Alloys at 1173 K and 1.3 Pa Sulfur Pressure in an

H2S-H2 Gas Mixture”, Intermetallics, 8, 891-901 (2000) (Experimental, 42)

[2000Kai] Kainuma, R., Fujita, Y., Mitsui, H., Ohnuma, I., Ishida, K., “Phase Equilibria Around

(hcp), (bcc) and (L10) Phases in Ti-Al Base Ternary Alloys”, Intermetallics, 8, 855-867


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Diagram, Review, 2)

300


MSIT®

Al–Mo–Ti

[2000Yan] Yang, R., Hao, Y., Song, Y., Guo, Z.-X., “Site Occupancy of Alloying Additions in

Titanium Aluminides and its Application to Phase Equilibrium Evaluation”, Z. Metallkd.,

91, 296-301 (2000) (Crys. Structure, Equi. Diagram, Review, 38)


Binary System Ti-Al”, Metall. Mater. Trans., 32A, 1037-1047 (2001) (Crys. Structure,

Equi. Diagram, Experimental, Review, 34)

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Ti-Al-X (X = V, Mn, Nb, Cr, Mo, Ni and Co) Systems by the Cluster Variation Method”,

J. Phase Equilib., 22, 424-430 (2001) (Equi. Diagram, Theory, 15)

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Alloys Containing Stabilizers by Directional Solidification”, Mater. Sci. Eng., A329-331,


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D.B., Legkaya, T.N., Voskoboynik, I.V., “The Influence of Zr Alloying on the Structure and

Properties of Al3Ti”, Intermetallics, 11, 241-249 (2003) (Crys. Structure, Experimental, 16)

[2003Sch1] Schmid-Fetzer, R., “Al-Ti (Aluminium-Titanium)”, MSIT Binary Evaluation Program, in



[2003Sch2] Schuster, J.C., “Al-Mo (Aluminium - Molybdenum)”, MSIT Binary Evaluation Program, in




Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Al)

< 664

< 660.452

cF4

Fm3m

Cu a = 404.96

0 to 0.6 at.% Ti [1992Kat, 2003Sch1]

0 to < 0.01 or 0.03 at.% Mo [2003Sch2]

pure Al at 25°C [1981Kin, Mas2]

, (Ti1-x-yMoxAly)

( Ti)(h)

1670 - 882

(Mo)

< 2623

cI2

Im3m

W

a = 330.65

a = 314.70

a = 314.2

a = 317.8

0 x 1 [Mas2]

pure Ti at 900°C;


at x = 0 [1992Kat, 1993Oka, 2003Sch1]


pure Mo [1981Kin, Mas2]

for Mo - ~20 at.% Al [1972Kam]

in Ti-50Al-5Mo annealed at 1240°C for 150 h

( + 2) [1993Das1]

* 2 cP2

Pm3m

CsCl

a = 321

a = 320.1

a = 321.5

ordered form of bcc (Ti,Mo,Al) solid

solution [1958Boe, 1972Ham, 1975Ham,

1991Dja, 1992Dja1, 1992Dja2, 1992Nak,

1993Das1, 1993Das2, 1994Li, 1994Mor,

1995Che, 1996Sik, 1997Che, 1997Nak,

1997Sin2, 1998Joh]

in the Ti-44Al-2Mo alloy

( 2+ 2+ ) HIPped at 1250°C, 150 MPa for

4 h [1994Li]

in as HIPped Ti-44Al-2Mo alloy [1994Mor]

in the Ti-44Al-2Mo alloy annealed at 1200°C

and 900°C [1994Mor]

301


MSIT®

Al–Mo–Ti

, ( Ti1-x-yMoxAly)

( Ti)(r)

< 882

hP2

P63/mmc

Mg

a = 295.03

c = 468.36

a = 294.9

c = 467.6

a = 293.1

c = 464.3

47.3 to 51.4 at.% Al at x = 0 at solidus

temperatures 1490 to 1463°C [1992Kat,

1997Zha, 2003Sch1]

~48 to 51 at.% Al at solidus temperatures

1520 to 1485°C [1996Tre1, 1997Bul]

pure Ti at 25°C [1981Kin, Mas2]

dissolves up to ~0.4 at.% Mo [Mas2]

for the single phase Ti-2.5Al-2.5Mo alloy

annealed at 800°C/222 h [1990Ere]

for ( + ) alloy (Ti-5Al-5Mo) annealed at

800°C/222 h [1990Ere]

, Mo3Al

2150

cP8

Pm3n

Cr3Si

a = 495

a = 496

a = 497

a = 498.7

a = 498.7

~23-28.5 at.% Al [2003Sch2]

dissolves up to ~14 at.% Ti at 1600°C, ~22

at.% Ti at 1300 and 1000°C [1987Ere,

1988Ere1, 1990Ere]

[V-C2]

in the Ti-40Al-50Mo ( + ) alloy annealed at

1000°C for 200 h [1990Ere]

in the Ti-40Al-40Mo alloy

( + + ) annealed at 1000°C for 200 h

[1990Ere]

in the Ti-19Al-55Mo alloy ( + ) annealed at

1600°C/53 h+1300°C/101 h [1990Ere]


1300°C for 101 h [1990Ere]

2, MoAl(h)

~1750 - 1470

cP2

Pm3m

CsCl

cI2

Im3m

W

a = 308.9 to 309.8

~46 to 52 at.% Al [2003Sch2]

[1971Rex]

1, Mo37Al63(h)

1570 - 1490

[1971Rex, Mas2, 1997Sau, 2003Sch2]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

302


MSIT®

Al–Mo–Ti

, Mo3Al8< 1555 10

mC22

c2/m

Mo3Al8

a = 920.8

b = 363.8

c = 1006.5

= 100.78°

a = 916.4

b = 363.9

c = 1004.0

= 100.50°

a = 920.7 0.3b = 364.1 0.1c = 1006.0 0.5

= 100.78

0.09°

a = 919

b = 363

c = 1008

= 101°

a = 913

b = 354

c = 1009

= 100.33°

a = 913

b = 362

c = 1002

= 100.62°

a = 916.2

b = 363.8

c = 1000.3

= 100.47°

a = 910

b = 364

c = 1005

= 100.82°

[V-C2]

72.7 at.% Al [2003Sch2]

[1991Sch]

[1990Ere], in the Ti-60Al-30Mo

( + + ) alloy annealed at 1300°C for 63 h


( + + ) alloy annealed at 1300°C for 107 h


( + ) alloy annealed at 1000°C


( + ) alloy annealed at 1000°C


( + + ) alloy annealed at 1000°C

MoAl3(h)

1222 - 818

mC32

Cm

MoAl3

a = 1639.6

b = 359.4 0.1c = 838.6 0.4

= 101.88°

[1991Sch, 1997Smi, 2003Sch2]

Mo1-xAl3+x(h)

1154 - 1260

cP8

Pm3n

Cr3Si

a = 494.5 76 to 79 at.% Al [1991Sch, 1997Smi,

2003Sch2]

MoAl4(h)

1177 - 942

mC30

Cm

WAl4

a = 525.5

b = 1776.8

c = 522.5

= 100.88°

79 to 80 at.% [1991Sch]

[V-C2]

Mo4Al17

< 1034

mC84

C2

Mo4Al17

a = 915.8 0.1b = 493.23 0.08

c = 2893.5 0.5 = 96.71 0.01°

[1991Sch] [1995Gri]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

303


MSIT®

Al–Mo–Ti

Mo5Al22(h)

964 - 831

oF216

Fdd2

Mo5Al22

a = 7382 3

b = 916.1 0.3c = 493.2 0.2

[1991Sch] [1995Gri]

MoAl5(h2)

846 -

(750 < T < 800)

hP12

P63

WAl5

a = 491.2 0.2c = 886.0 0.4a = 493.7

c = 924.3

[1991Sch]

[V-C2]

MoAl5(h1)

(850 - 750) - 648

hP60

P3

a = 493.3 0.1c = 4398 9

[1991Sch]

MoAl5(r)

< 650

hP36

R3c

a = 495.1 0.1c = 2623 1

[1991Sch]

MoAl12

712

cI26

Im3

WAl12

a = 758.15

a = 758.77

[V-C2]

[1991Sch]

,

(Ti1-xMox)1+yAl3-y

TiAl3

tI18

I4/mmm

TiAl3(h)

a = 384.9

c = 860.9

a = 385.3

c = 858.7

a = 384

c = 859

a = 380.7

c = 839.2

a = 379.8

c = 836.7

a = 380

c = 841

a = 384.0

c = 830.8

a = 387.1

c = 831.8

a = 389

c = 829

a = 383

c = 849

a = 387.4

c = 830.3

a = 390

c = 825

a = 386.5 0.3c = 843.9 0.1

[V-C2], D022 ordered phase

0 x 0.88; 0 y ~0.21

[1970Han, 1987Ere, 1990Ere, 1996Tre2,

1990Abd]

72.4 to 75.0 at.% Al [2003Sch1]

< 1425°C [1999Tre1, 1997Bul]

1385 to 735°C, 74.5-75 at.% Al at 1200°C

[2001Bra]

melting temperature 1408°C [2003Kar]

[1970Han]

Ti-75Al-12.5Mo [1970Han]

Ti-76Al-16Mo [1970Han]

Ti-75Al-20Mo annealed at 1000°C for 121 h

[1990Ere]



Ti-62.5Al-12.5Mo annealed at 1000°C

[1990Ere]

Ti-70Al-13Mo annealed at 1000°C for 100 h

[1990Ere]

Ti-65Al-15Mo annealed at 1300°C/50 h +

1000°C/147 h [1990Ere]

Ti-60Al-15Mo annealed at 1300°C/50 h +

1000°C/ 147 h [1990Ere]

Ti-67Al-10Mo, 1300°C [1996Tre3]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

304


MSIT®

Al–Mo–Ti

TiAl3(l)

< 950

tI32

I4/mmm

TiAl3(l)

a = 387.7

c = 3382.8

74.5 to 75 at.% Al [2001Bra]

tetragonal

superstructure

of AuCu type

tI16

I4/mmm

ZrAl3

tP28

P4/mmm

Ti2Al5

a* = 395.3

c* = 410.4

a* = 391.8

c* = 415.4

a = 398.81 to

392.3

c = 1646.69 to

1653.49

a = 399.1 1.3

c = 1646.6 0.5

a = 392.8 0.6

c = 1656.3 1.5

a = 390.53

c = 2919.63

summarizes several phases [2003Sch1]

Ti5Al11 [2001Bra]

stable in the range 1416-995°C,

66 to 71 at.% Al at 1300°C [2001Bra]

(including the stoichiometry Ti2Al5)

at 66 at.% Al, * AuCu subcell only [2001Bra]

at 71 at.% Al, * AuCu subcell only [2001Bra]

Ti5Al11, D023 type [V-C]

65.8 to 70.9 at.% Al, 1416-1206°C [1990Sch]

69 to 71 at.% Al, 1450-~990°C [1996Tre1,

1997Bul]

in the as cast Ti-68Al-2Mo alloy ( + )

[1996Tre3]


1300°C for 24 h [1996Tre3]

Ti2Al5, 1416-990°C [1992Kat],

~1215-985°C [1990Sch]; included in the

homogeneity range of Ti5Al11 [2001Bra]

, TiAl2< 1199

tP4

P4/mmm

AuCu

orthorhombic,

Pmmm, with

pseudotetragonal

cell

tI24

I41/amd

HfGa2

oC12

Cmmm

ZrGa2

tP32

P4/mbm

Ti3Al5

a = 403.0

c = 395.5

a = 402.62

b = 396.17

c = 402.62

a = 397.0

c = 2430.9

a = 396.7

c = 2429.68

a = 1208.84

b = 394.61

c = 402.98

a = 1209.44

b = 395.91

c = 403.15

a = 1129.3

c = 403.8

chosen stoichiometry [1992Kat] summarizes

several phases [2003Sch1]:

Ti1-xAl1+x, 63 to 65 at.% Al at 1300°C, stable

in the range 1445-1170°C [2001Bra]

for Ti36Al64 at 1300°C [2001Bra]

1445-1424°C [1990Sch]

for as arc melted Ti36Al64 [1990Sch]

stable structure of TiAl2 < 1216°C,

66 to 67 at.% Al at 1000°C [2001Bra];

shown as TiAl2(r) < 1214°C [1900Sch]

metastable modification of TiAl2 observed

only in as cast alloys [2001Bra]

TiAl2(h), 66 to 67 at.% Al, 1433-1214°C

[1990Sch]


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

305


MSIT®

Al–Mo–Ti

, TiAl

< 1463

tP4

P4/mmm

AuCu

a = 400.5

c = 407.0

a = 400.0 0.1c = 407.5 0.1

a = 398.4 0.1c = 406.0 0.1

a = 398.1

c = 407.5

a = 397

c = 408

a = 399.2

c = 405.6

a = 396

c = 408

a = 396.0 1.1

c = 407.5 0.2

a = 400.6

c = 405.7

a = 401.3

[V-C], L10 ordered phase

46.7 to 66.5 at.% Al [1992Kat, 1993Oka];

50 to 62 at.% Al at 1200°C [2001Bra]

~52 to 65 at.% Al at solidus temperatures,

~50 to 60 at.% Al at 1000°C [1996Tre1,

1997Bul]

at 50 at.% Al [2001Bra]

at 62 at.% Al [2001Bra]

in Ti-50Al-5Mo ( + 2) alloy annealed at

1240°C for 150 h [1993Das1]

Ti-55Al-5Mo annealed at 1300°C/ 111 h +

1000°C/ 150 h [1990Ere]

in Ti-65Al-5Mo ( + ) alloy annealed at

1300°C /13 h + 1000°C/ 26 h + 800°C/ 205 h

[1996Tre3]

the same ( + ) alloy annealed at 1300°C/

150 h [1990Ere]

in Ti-50Al-10Mo ( + ) alloy annealed at

1300°C/ 13 h + 1000°C 26 h + 800°C/ 205 h

[1996Tre3]

in Ti-34.5Al-1.5Mo(mass%) alloy

(Ti50.8Mo0.6Al48.6) annealed at 1000°C for 1

h [1991Mae]

c/a = 1.008, 1.015 or 1.013 in the

Ti-44Al-2Mo alloy as HIPpped (High

Isostatic Pressed), annealed at 900°C or

1200°C, respectively [1994Mor]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

306


MSIT®

Al–Mo–Ti

2, (Ti3Al)

Ti3Al

< 1164

hP8

P63/mmc

Ni3Sn

a = 578.2

c = 468.9

a = 580.6

c = 465.5

a = 574.6

c = 462.4

a = 577.5 0.4c = 463.7 0.5a = 579.5

c = 464.1

a = 567

c = 451

a = 606

c = 495

a = 576.2

c = 461.9

D019 ordered phase

~20 to 38.2 at.% Al, maximum at 30.9 at.% Al

[1992Kat, 1993Oka, 2003Sch1]

< 1180°C [1993Gam]

maximum at 32.5 at.% Al, ~1200°C

[1996Tre1, 1997Bul]

< 1210°C ( + 2) [1994Kai, 2000Oka]

[V-C]

at 28 at.% Al [L-B]

at 28 at.% Al [L-B]

at 32 at.% Al [1997Bul]

at 25 at.% Al, annealed at 1300°C/40 h

+ 1000°C/90 h + 800°C/222 h [1990Ere]

metastable 2 phase

(Ti-54.2Al-13.0Mo) in the Ti-50Al-15Mo

alloy annealed at 1400°C/2 h and water

quenched (WQ) [1997Che]

in the Ti-21.6Al-3.4Mo alloy aged at 450°C

[1992Dja1]

in the Ti-44Al-2Mo alloy ( 2+ 2+ )

[1994Mor]

* , ~Ti3Mo3Al4 tP30

P42/mnm

CrFe a = 966.7

c = 501.8

a = 965.1

c = 501.8

a = 963.6

c = 499.7

a = 959.1

c = 496.6

a = 966

c = 502

a = 966

c = 501

[V-C2], single phase (Ti26Al41Mo33

at 925°C) [1970Han]


[1970Han]


[1970Han]


[1970Han]


[1970Han]

(Ti~28Al~40Mo~32)

in the Ti-40Al-30Mo ( + + ) alloy annealed

at 1300°C/63.5 h+1000°C/200 h [1990Ere]

in the Ti-40Al-40Mo alloy ( + + ) annealed

at 1000°C [1990Ere]

L60 tP4

P4/mmm

a = 395

b = 403

b/a = 1.020

intermediate phase observed in the

Ti-50Al-15Mo alloy annealed at 1400°C for

1.5 h and WQ (composition of the phase

Ti-57.4Al-9.4Mo) [1997Che]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

307


MSIT®

Al–Mo–Ti

´ tI18

I4/mmm

a = 397

c = 815

c/a = 2.05

the ordered D022 type form of the phase

formed by diffusionless way, observed in the

Ti-50Al-15Mo alloy quenched from

temperatures in the range of 1400-800°C

[1997Che]

´´ tP18

P4/mmm

the ordered phase observed in the

Ti-50Al-15Mo alloy after prolonged aging at

800°C, supposed to be formed as a result of

further ordering of the ´ (D022) [1997Che]

´ hP2

P63/mmc

Mg

martensite phase in Ti-xMo-3Al alloys

(0 x 4) [1980Sas]

´´ oC4

P2221

U

martensite phase in Ti-xMo-3Al alloys

(7 x 12) [1980Sas]

hP3

P6/mmm

TiCr

metastable phase, appeared during quenching

of / 2 phases ( athermal) or aging of

metastable (quenched) / 2 phases (

isothermal) [1971Wil, 1972Luz, 1980Sas,

1991Dja, 1992Dja1, 1992Dja2, 1993Cui,

1997Che]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

308


MSIT®

Al–Mo–Ti

Fig

. 1

a:

Al-

Mo

-Ti.

Rea

ctio

n s

chem

e up t

o 7

5 a

t.%

Al

Al-

Ti

Mo

-Ti

A-B

-CA

l-M

o-T

iA

l-M

o

β +

l

α1

490

p4

L +

ρβ

+ ζ 2

1750>

T>

1600

U1

l + β

ρ2

150

p1

α +

l

γ1

463

p6

l +

γζ

14

16

p7

lζ

+ ε

13

93

p8

γ +

ζη

11

99

p9

l +

ζ2

ζ 1

15

70

p2

lδ

+ ζ

1

15

35

e 2

ζ 1δ

+ ζ 2

14

90

e 3

ζ 2δ

+ ρ

14

70

e 4

αα 2

+ γ

11

18

e 6

ζε

+ η

99

0e 7

lρ

+ ζ

2

17

20

e 1

βα

+ β´

67

5e 8

β +

l

α1

60

0>

T>

15

00

p3

L +

ζ2

β +

ζ1

15

50

U2

L +

ζ1

β +

δ1

500

U3

L+

δε

>1475

p5

β +

ζ 1δ

+ ζ

21

495

U4

δ +

L

β +

ε1

470

U5

ζ 2β

+ δ

+ ρ

14

55

E1

α +

L

β +

γ1

440

U6

L +

ζγ

+ ε

14

00

U7

Lβ

+ γ/

γ´ +

ε1

370

E2

β +

δε

+ ρ

13

25

U8

β +

ε +

δσ

12

50

P

γ +

ζε

+ η

11

45

U9

αα 2

+ β

+ γ

<1125

E3

β +

ε γ/

γ´ +

σ1000>

T>

925

U1

0 α +

βα 2

+ β

´<

550

U1

1

αα 2

+ β

>1125

e 5

L+

β+ζ 2 L

+β+

ζ 1

β+δ+

ζ 1

β+ζ 1

+ζ 2

Lβ+

δ

β+δ+

ζ 2

Lβ+

ε

β+δ+

ε

β+δ+

ρ

L+

β+γ

γ/γ´

+ε+

L

β+ε+

ρε+

δ+ρ

γ+ε+

ζ

β+γ/

γ´+

ε

α+α 2

+β

α+α 2

+β

γ+ε+

η

α+β+

γ

α 2+

β+γ

β+γ/

γ´+

σγ/

γ´+

ε+σ

α+α 2

+β´

α+β+

β´

β+ε+

σβ+

ρ+σ

ε+ρ+

σ

β+ρ+

ζ 2

309


MSIT®

Al–Mo–Ti

Fig

. 1

b:

Al-

Mo

-Ti.

Rea

ctio

n s

chem

e fo

r th

e A

l-ri

ch p

art

Al-

Ti

Mo

-Ti

A-B

-C

l +

ε (

Al)

66

4p

8

Al-

Mo

-Ti

L+

δε+

Mo

1-x

Al 3

+x

U1

Al-

Mo

l +

δ M

o1-x

Al 3

+x

12

60

p1

δ+M

o1-x

Al 3

+x

MoA

l 3

12

22

p2

l+M

o1-x

Al 3

+x

MoA

l 4

11

77

p3

Mo

8A

l 22

Mo

4A

l 17+

MoA

l 5

83

1e 3

l +

Mo

8A

l 22

MoA

l 5

84

6p

6

Mo

1-x

Al 3

+x

Mo

Al 3

+M

oA

l 4

11

54

e 1

l+M

o4A

l 17

Mo

8A

l 22

96

4p

5

Mo

Al 4

Mo

Al 3

+M

o4A

l 17

94

2e 2

MoA

l 3δ+

Mo

4A

l 17

81

8e 4

δ+M

o1-x

Al 3

+x

ε+M

oA

l 3U

2

L+

Mo

1-x

Al 3

+x

ε+M

oA

l 4U

3

ε+M

o1

-xA

l 3+

xM

oA

l 4+

Mo

Al 3

U4

L+

MoA

l 4ε+

Mo

4A

l 17

U5

ε+M

oA

l 4M

oA

l 3+

Mo

4A

l 17

U6

L+

Mo

4A

l 17

ε+M

o8A

l 22

U7

L+

Mo

8A

l 22

ε+M

oA

l 5U

8ε+

Mo

8A

l 22

Mo

4A

l 17+

Mo

Al 5

U9

ε+M

oA

l 3δ+

Mo

4A

l 17

U1

0

L+

ε+M

o1

-xA

l 3+

x

l+M

oA

l 5M

oA

l 12

71

2p

7

lM

oA

l 12+

(Al)

66

0e 5

Mo

Al 5

Mo

4A

l 17+

Mo

Al 1

2

64

8e 6

L+

MoA

l 5ε+

MoA

l 12

U1

1

Lε+

MoA

l 12+

(Al)

65

0E

ε+M

oA

l 5M

o4A

l 17+

MoA

l 12

U1

2

L+

δ+ε

δ+ε+

Mo

1-x

Al 3

+x

Mo

Al 3

+M

o1

-xA

l 3+

x+

ε

L+

ε+M

oA

l 4

ε+M

oA

l 3+

Mo

Al 4

l+M

oA

l 4M

o4A

l 17

10

34

p4

L+

ε+M

o4A

l 17

ε+M

oA

l 3+

Mo

4A

l 17

ε+M

o4A

l 17+

Mo

8A

l 22

L+

ε+M

o8A

l 22

ε+M

o4A

l 17+

Mo

Al 5

ε+M

o8A

l 22+

Mo

Al 5

ε+δ+

Mo

Al 3

ε+δ+

Mo

4A

l 17

L+

ε+M

oA

l 5

L+

ε+M

oA

l 12

ε+M

oA

l 5+

Mo

Al 1

2

ε+M

oA

l 12+

(Al)

ε+M

oA

l 12+

Mo

4A

l 17

Mo

Al 5

+M

o4A

l 17+

Mo

Al 1

2

ε+M

o1

-xA

l 3+

x+

Mo

Al 4

ε+M

oA

l 4+

Mo

4A

l 17

310


MSIT®

Al–Mo–Ti

40

50

10 20

50

60

Ti 60.00Mo 0.00Al 40.00

Ti 30.00Mo 30.00Al 40.00

Ti 30.00Mo 0.00Al 70.00 Data / Grid: at.%

Axes: at.%

γ

α

β

p6

p4

U6

p3

20

40

60

80

20 40 60 80

20

40

60

80

Ti Mo


Axes: at.%

~1400 >1415

~1500

1370

~1440 ~1550

δ

ζ1

ζ2

ζ

α

β

ρ

Mo1-x

Al3+x

γ

ε

1470

Fig. 3: Al-Mo-Ti.

Partial liquidus

surface projection

Fig. 2: Al-Mo-Ti.

Projection of the

partial solidus surface

311


MSIT®

Al–Mo–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Mo


Axes: at.%

L

ζ2

ρβ

L+β

40

50

10 20

50

60

Ti 60.00Mo 0.00Al 40.00

Ti 30.00Mo 30.00Al 40.00


Axes: at.%

α

L

β

L+α

α+β

L+β

L+β

Fig. 4: Al-Mo-Ti.


1600°C

Fig. 5: Al-Mo-Ti.

Partial isothermal

section at 1500°C

312


MSIT®

Al–Mo–Ti

40

50

10 20

50

60

Ti 60.00Mo 0.00Al 40.00

Ti 35.00Mo 25.00Al 40.00


Axes: at.%

γ

α

β

α+γ

α+β

γ+β

20

40

60

80

20 40 60 80

20

40

60

80

Ti Mo


Axes: at.%

εγ'

γ

ρ

L

β2

β

α

δζ

Fig. 6: Al-Mo-Ti.

Tentative partial

section at 1400°C

Fig. 7: Al-Mo-Ti.


1300°C

313


MSIT®

Al–Mo–Ti

50

60

70

10 20 30

30

40

50

Ti 80.00Mo 0.00Al 20.00

Ti 40.00Mo 40.00Al 20.00


Axes: at.%

α

β

γ

β+γ

50

60

70

80

10 20 30 40

20

30

40

50

Ti 90.00Mo 0.00Al 10.00

Ti 40.00Mo 50.00Al 10.00


Axes: at.%

γ

α2

α

β

α2+β

Fig. 8: Al-Mo-Ti.

Partial isothermal

section at 1200°C

Fig.9: Al-Mo-Ti.

Partial isothermal

section at 1100°C

314


MSIT®

Al–Mo–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Mo


Axes: at.%

ε

γ

γ1

β2

β

ρ

σα

2

α

MoAl4(h)

MoAl3(h)

δζη

L

Mo4Al

17

20

40

60

80

20 40 60 80

20

40

60

80

Ti Mo


Axes: at.%L

Mo5Al

22(h)

Mo4Al

17

MoAl3(h)

δ

ρ

σ

γ

εη

α2

α

β

β2

Fig. 10: Al-Mo-Ti.


1000°C

Fig. 11: Al-Mo-Ti.


925°C

315


MSIT®

Al–Mo–Ti

60

80

20 40

20

40

Ti Ti 40.00Mo 60.00Al 0.00


Axes: at.%

γ

β2

β'βα

α2

20

40

60

80

20 40 60 80

20

40

60

80

Ti Mo


Axes: at.%

η

α

β

β2

β+β'

ε

α2

γ

Fig. 12: Al-Mo-Ti.

Partial isothermal

section at 800°C

Fig. 13: Al-Mo-Ti.

Partial isothermal

section at 700°C

316


MSIT®

Al–Mo–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Mo


Axes: at.%

ε

γ

α

α2

σ

β'

(Al)

η

ββ

2

ρ

MoAl12

MoAl5(r)

Mo4Al

17

δ

Fig. 14: Al-Mo-Ti.


600°C

317


MSIT®

Al–N–Si

Aluminium – Nitrogen – Silicon

Hans Leo Lukas

Literature Data

During the investigation of the quaternary Al-Si-N-O system [1975Gau] found a phase Al5+xSi3-xN9-xOx,

which exists in the range 0 x 3. [1978Sch, 1980Sch] however, assumed this phase to be unstable in the

ternary Al-N-Si system, as it needs some oxygen to be stabilized. Other investigations regarding quaternary

systems with AlN-Si3N4 as boundary system [1978Lan, 1983Hua, 1986Hua, 1988Fuk, 1990Wei] do not

mention this phase and assume AlN to be in equilibrium with Si3N4. [1992Hil] thermodynamically

calculated the Al-N-Si system, assuming the ionic liquid model with ideal solution behavior for a nitride

liquid. These authors did not consider Al5Si3N9 to be a stable phase. Thus at 1 bar pressure the only stable

phases taken into account are liquid, the solid metals (Al) and Si, solid AlN and solid Si3N4. All these phases

have only small ranges of homogeneity, which for AlN and Si3N4 were neglected by [1992Hil] in their

calculation.

[2001Kas] synthesized Al1-xSixN solid solutions up to x = 0.12 by metalorganic vapor-phase epitaxial

growth. From the thermodynamic point of view it is very likely, that this solid solution has to be considered

as metastable supersaturated although the crystal quality is very perfect, measured by the full width at half

maximum of 100 arcsec ( = 0.028°) of an X-ray rocking curve (single crystal rotation technique focussed

on a single X-ray peak). The temperature during preparation (900°C) may be far too low to enable

equilibration. The solid solution was characterized as substitutional, one Si atom replaces one Al atom. The

same authors [2001Tan] reported lattice parameter measurements of Al1-xSixN in dependence of x,

extrapolated from the epitaxial layer to zero residual strain. [2002Wu] prepared Al containing solid

solutions of Si3N4 by Al ion implantation in order to study the influence of Al on the oxidation behavior of

Si3N4. No structural details of the solid solution were reported.

Binary Systems

The Al-Si system is accepted from [2003Luk]; it is based on the thermodynamic assessment of [1997Feu].

The N-Si and Al-N systems are accepted from the thermodynamic assessments of [1991Hil1] and

[1991Hil2], respectively. The calculation of the ternary system by [1992Hil] used the latter two binary

assessments and an older assessment of the Al-Si system without any ternary excess term. The calculated

results, except near the eutectic of the binary Al-Si system, do not show a visible dependence on the

selection of the binary Al-Si assessment.

Solid Phases

Stable binary phases are AlN and Si3N4. Pure Si3N4 is metastable but formed as the main product during

reaction of Si with N2. It is stabilized, however, by large cations, e.g. rare earth oxides. The phase Al5Si3N9

possibly exists only in the oxygen stabilized form Al5+xSi3-xN9-xOx with x > 0. All solid phases are

summarized in Table 1.


At 1 bar pressure the only four-phase equilibria are: (i) Gas+L Si3N4+AlN at 1839.5°C, which is nearly

degenerated and very near to the quasibinary three-phase equilibrium Gas+L Si3N4 at 1840.5°C; (ii)

L (Al)+(Si), AlN at 577°C, which is totally degenerated and identical to the binary Al-Si eutectic.

Isothermal Sections

Figures 1 and 2 show the isothermal sections at 1 bar and 2400 or 1800°C, calculated by [1992Hil].

318


MSIT®

Al–N–Si


The AlN-Si3N4 section, calculated for 1 bar [1992Hil], is shown in Fig. 3.


Epitaxially grown Al1-xSixN layers are promising candidates as materials for flat panel displays

(FE-displays), as the Si content in AlN decreases the electric field necessary for field emission (FE). As part

of the Al-Si-N-O system, Al-Si-N is interesting for high temperature materials based on SIALON.

References

[1975Gau] Gauckler, L.J., Lukas, H.L., Petzow, G., “Contribution to the Phase Diagram

Si3N4-AlN-Al2O3-SiO2”, J. Am. Ceram. Soc., 58, 366-367 (1975) (Experimental, Equi.

Diagram, 10)

[1976Jac] Jack, K.H., “Review: Sialon and Related Nitrogen Ceramics”, J. Mater. Sci., 11, 1135-1158

(1976) (Review. Equi. Diagram, Crys. Structure, 41)

[1978Lan] Land, P.L., Wimmer, J.M., Barns, R.W., Choudhury, N.S., “Compounds and Properties of

the System Si-Al-O-N”, J. Am. Ceram. Soc., 61, 56-60 (1978) (Experimental, Equi.

Diagram, 25)

[1978Sch] Schneider, G., “Equilibrium Investigations in the Si, Al, Be/C, N System” (in German),

Thesis, University of Stuttgart, Germany (1978) (Experimental, Equi. Diagram, Crys.

Structure, 71)

[1980Sch] Schneider, G., Gauckler, L.J., Petzow, G., “Phase Equilibria in the System AlN - Si3N4 -

Be3N2”, J. Am. Ceram. Soc., 63, 32-35 (1980) (Experimental, Equi. Diagram, 7)

[1983Hua] Huang, Z.K., Greil, P., Petzow, G., “Formation of -Si3N4 Solid Solutions in the System

Si3N4-AlN-Y2O3”, J. Am. Ceram. Soc., 66, C-96-C-97 (1983) (Experimental, Equi.

Diagram, 5)

[1986Hua] Huang, Z.K., Tien, T.-Y., Yen, T.-S., “Subsolidus Phase Relationships in Si3N4-AlN-Rare

Earth Oxide Systems”, J. Am. Ceram. Soc., 69, C-241-C-242 (1986) (Experimental, Equi.

Diagram, 5)

[1988Fuk] Fukuhara, M., “Phase Relationships in the Si3N4 Rich Portion of the

Si3N4-AlN-Al2O3-Y2O3 System”, J. Am. Ceram. Soc., 71, C359-361 (1988) (Experimental,

Equi. Diagram, 10)

[1990Wei] Weitzer, F., RemsChnig, K., Schuster, J.C., Rogl, P., “Phase Equilibria and Structural

Chemistry in the Ternary Systems M-Si-N and M-B-N (M = Al, Cu, Zn, Ag, Cd, In, Sn, Sb,

Au, Tl, Pb, Bi)”, J. Mater. Res., 5, 2152-2159 (1990) (Experimental, Equi. Diagram, Crys.

Structure, 39)

[1991Hil1] Hillert, M., Jonsson, S., “Report, Trita-Mac-465”, Royal Inst. of Technology, Stockholm,

Sweden, (1991) (Thermodyn., Equi. Diagram, Assessment, 0)

[1991Hil2] Hillert, M., Jonsson, S., “Report, Trita-Mac-466”, Royal Inst. of Technology, Stockholm,

Sweden, (1991) (Thermodyn., Equi. Diagram, Assessment, 0)

[1992Hil] Hillert, M., Jonsson, S., “Prediction of the Al-Si-N System”, Calphad, 16, 199-205 (1992)

(Thermodyn., Equi. Diagram, Assessment, 11)

[1997Feu] Feufel, H., Gödecke, T., Lukas, H.L., Sommer, F., “Investigation of the Al-Mg-Si System

by Experiments and Thermodynamic Calculations”, J. Alloys Comp., 247, 31-42 (1997)

(Experimental, Assessment, Thermodyn., Equi. Diagram, 38)

[2001Kas] Kasu, M., Taniyasu, Y., Kobayashi, N., “Formation of Solid Solution of Al1-xSixN

(0 x 12%) Ternary Alloy”, Jpn. J. Appl. Phys. 2, 40(10A), L1048-L1050 (2001)

(Experimental, 12)

[2001Tan] Taniyasu, Y., Kasu, M., Kobayashi, N., “Lattice Parameters of Wurtzite Al1-xSixN Ternary

Alloys”, Appl. Phys. Lett., 79(26), 4351-4353 (2001) (Experimental, Crys. Structure, 14)

319


MSIT®

Al–N–Si

[2002Wu] Wu, J., Wang, Y., Ye, J., Du, H.H., “The Cyclic and Continuous Oxidation of with and

without Aluminum Implantation”, Key Eng. Mater., 224-226, 803-806 (2002)

(Experimental, Corrosion, 14)





Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Al)

< 660.45

cF4

Fm3m

Cu

a = 404.93 at 23°C [V-C2]

(Si)

< 1414

cF8

Fd3m

C (diamond)

a = 543.06 at 30°C [V-C2]

Al1-xSixN

< 2800 50

hP4

P63mc

ZnS (wurtzite)

a = 311.15

c = 497.98

a = 311.13 - 14.12x

c = 498.18 - 22.99x

at 17°C, x = 0 [V-C2]

0 x 0.12 [2001Tan]

metastable ?

Si3N4 hP14

Be2SiO4

a = 760.8

c = 291.1

[V-C2]

Si3N4 hP28

Si3N4

a = 775 to 782

c = 562 to 559

metastable,

stabilized by rare earth oxides

three sets of parameters [V-C2]

Al5+xSi3-xN9-xOx hexagonal a = 307.9

c = 530

0(?) x 3 [1975Gau]

possibly not stable at x = 0 [1978Sch,

1980Sch];

parameters from [1976Jac] for

Si3Al7N11 formula

320


MSIT®

Al–N–Si

20

40

60

80

20 40 60 80

20

40

60

80

Al Si

N Data / Grid: at.%

Axes: at.%

L

Gas

L+Gas

L+AlN

L+Gas

+AlN

AlN

Fig. 1: Al-N-Si.


2400°C

20

40

60

80

20 40 60 80

20

40

60

80

Al Si

N Data / Grid: at.%

Axes: at.%

L

Si3N

4

AlN

L+AlN

L+AlN

+Si3N

4

Gas+AlN

+Si3N

4

Fig. 2: Al-N-Si.


1800°C

321


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Al–N–Si

10 20 30 40

1750

2000

2250

2500

2750

3000

Al 50.00Si 0.00N 50.00

Al 0.00Si 42.86N 57.14Si, at.%

Te

mp

era

ture

, °C

Gas

L+Gas

L+Gas+AlN

Si3N4+AlN Si3N4+L+Gas

1839.61840.6°C

Fig. 3: Al-N-Si.

Section from AlN to

Si3N4

322


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Al–N–Ti

Aluminium – Nitrogen – Titanium

Vasyl Tomashik and Pierre Perrot

Literature Data

A critical assessment of the Al-N-Ti ternary system has been published by [1993Jeh], which included the

literature data up to the year 1991. Thermodynamic data appearing up to 1997 are included in the

thermodynamic assessment made by [1998Che]. Subsequently this system was investigated in different

experimental approaches and for different temperatures. The present evaluation takes care of all data, from

the first publication to the present.

The investigations in this ternary system are concerned with (a) phase diagram studies, (b) preparation and

characterization of the ternary compounds and (c) the formation of metastable solid solutions in the

AlN-TiN pseudobinary system.

The equilibria in the Ti-rich part of the ternary system have been determined by [1954Thy] for 0 to 10

mass% Al and 0 to 1 mass% N. This study applied micrograph analysis and X-ray diffraction of samples

annealed at 600 to 1250°C, for 576 to 6 h. These samples were prepared from high purity arc molten alloys.

The obtained results are given as vertical sections for constant N content.

Annealing of Al-TiN bilayers on SiO2 for 15 h at 645°C leads to the formation of AlN and Al3Ti, as the

data of [1982Wit] show. These phases are also formed by reaction sintering of powder mixtures Al+TiN,

containing 10, 15, 20 and 30 mol% TiN [1992Koy]. Titanium specimens with embedded AlN particles, as

well as AlN-Ti and AlN-TiN diffusion couples were annealed at 900 to 1000°C up to 40 h by [2000Par]. It

was shown that in AlN-TiAl diffusion couples a ternary Ti2AlN phase is formed at the interface. A more

complex AlN-TiN-Ti3AlN-Ti3Al- Ti- Ti reaction zone was observed at the AlN-Ti interface.

Thermodynamic calculations give the same sequence of expected layers between AlN and pure Ti

[1998Lee] (the composition of Ti at the Ti3AlN/ Ti interface is close to the composition of Ti3Al). AlN

never is in contact with Ti3AlN [2000Par].

Nitriding the intermetallic TiAl3 in nitrogen and ammonium flow was studied by [1983Psh] in a temperature

range of 600 to 1200°C. This work states that Al and Ti are nitrated in fact simultaneously, which results in

formation of a heterogeneous mixture of practically not interacting binary nitrides.

Experimental results imply that AlN-TiN, TiAl3-AlN and TiAl3-TiN are stable tie lines in the Al-N-Ti

ternary system at low temperatures [1984Bey]. Phase equilibria in this ternary system were investigated at

1000 and 1300°C using previously prepared Al-Ti alloys, AlN, TiN and Ti powders [1984Sch]. About 30

ternary alloys were cold-pressed and sintered at the following conditions: 1000°C for 240 to 800 h in BN

crucibles sealed in evacuated quartz tubes, 1200°C for 60 h in Mo crucibles under dynamic vacuum, 1300°C

for 60 h in Mo crucibles under dynamic vacuum or for 50 h in BN (Mo) crucibles under argon and 1400°C

in Mo crucibles under dynamic vacuum. As the alloys sintered at 1000°C were initially not in equilibrium

they were powderized again, cold-pressed and sintered again. These two isothermal sections were included

in the reviews [1985Sch, 1992Sch, 1993Jeh, 1998Che]. The isothermal section at 900°C was constructed

by [1997Dur] which was supported by the thermochemical calculations. Based on such calculations the

1000°C isotherm is expected to be virtually not altered with respect to the 900°C isotherm, which disagrees

with [1984Sch]. It was concluded by [1997Dur] that the samples of [1984Sch] were not heat treated

sufficiently long to reach equilibrium at 1000°C. For the 850°C isotherm the thermochemical calculations

predict a three-phase field AlN+TiN+TiAl2 rather than AlN+TiN+TiAl3 [1997Dur]. The 1325°C

equilibrium isothermal section of the Al-N-Ti ternary system with accounting of Ti4AlN3-x formation was

constructed by [2000Pro2]. The isothermal sections of the Al-N-Ti system at 1200, 1400, 1580, 1600, 1900

and 2500°C were calculated thermodynamically by [1998Che] but in these calculations the existence of the

Ti3Al2N2 was taken into account. As the new investigations indicate that the more probable composition of

Ti3Al2N2 in this system is Ti4AlN3-x these isothermal sections must be recalculated and the Al-N-Ti ternary

system needs a revised thermodynamic assessment.

323


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Al–N–Ti

The phase diagram of the AlN-TiN pseudobinary system was calculated by [1988Hol, 1989Hol].

Unfortunately, they used a melting point of 2930°C for TiN instead of 3290°C. It was shown that the

eutectic temperature and eutectic composition depend on the size of both AlN and TiN particles [1991And].

Mixtures of AlN and TiN containing from 30 to 90 mass% AlN did not show evidence of a reaction between

these two materials [1976Kuz]. Similarly no reaction was found in annealing TiN powder and AlN plates

up to 2000°C for 6 h. AlN-TiN composite materials were prepared by pressureless sintering in N2

atmosphere at 1870°C for 6 h [2002Tan].

Ti1-xAlxN metastable solid solutions (0 x < 0.7) can be obtained in the AlN-TiN pseudobinary system

using cathodic arc plasma depositing process [1981Bee, 1986Mue, 1988Pen, 1988Ran, 1991Ike, 1992Tan,

1993Tan], or reactive dc and radio-frequency magnetron sputtering [1986Jeh, 1986Kno, 1987Hak,

1987Ina, 1987Kno, 1988Jeh, 1990McI, 1991Adi, 1993Pet, 1993Wah]. Such films could be prepared onto

polished flat high speed steel surfaces [1986Jeh, 1986Kno, 1988Jeh], or stainless-steel substrates [1988Pen,

1990McI] or stellite surfaces [1986Kno], or MgO(001) substrates [1991Adi, 1991Hul, 1993Adi, 1993Pet,

1993Wah], or oxidized silicon surfaces [1991Hul, 1993Adi, 1993Wah], or Si and WC-Co substrates

[1992Tan, 1993Tan]. These solid solutions based on TiN1-x phase crystallize in a cubic structure [1986Jeh,

1988Jeh] and the lattice parameter of the Ti1-xAlxN films linearly decreases with increasing Al content

[1986Kno, 1987Ina, 1987Kno, 1993Adi, 1993Tan, 1993Wah]. According to the data of [1986Jeh, 1988Jeh]

another phase was found in coatings deposited at low nitrogen pressures and in pure Ar atmosphere.

Although Ti0.5Al0.5N is thermodynamically metastable it exhibits a good high-temperature stability during

post annealing [1991Hul]. Such alloys deposited at 400°C were stable up to 1.5 h at 900°C [1990McI]). The

films which contain more than 70 mol% AlN crystallize in the wurtzite structure [1991Hul, 1992Tan,

1993Tan, 1993Wah]. According to the data of [1981Bee] the amorphous Ti1-xAlxN films can be obtained

when the N2 content in Ar-N2 atmosphere is greater than 20%. The existing experimental results and

thermodynamic calculations lead to a so-called vapor deposition phase diagram representing the range of

metastable phases which were established by [1988Hol, 1989Hol] and then refined by [2001Spe]. The

composition at which the structural transition takes place was experimentally verified at about 63 and 69

mol% AlN [2001Spe].

Binary Systems

Al-N: The solubility of nitrogen in Al(s) and Al(l) is very small. Only one compound AlN exists in the Al-N

binary system. The decomposition temperature of AlN under 0.1 MPa nitrogen pressure is 2437.4°C

[2003Fer]. AlN undergoes a congruent melting point towards 2800 50°C under a nitrogen pressure of

10 MPa [1984Jon]. On increasing nitrogen pressure above 1GPa, AlN undergoes a transition from the

wurtzite type to the rock salt type structure.

Al-Ti: Three ordered phases Ti3Al, TiAl and TiAl3 are stable in this system [2003Sch]. The composition

range between the phases TiAl and TiAl3, however, is still controversial, especially at temperatures above

1200°C because of the large number of long period structures. In total of five phases were suggested for this

region, some occurring in narrow temperature ranges only and/or with a range of solubility. These five

phases were subsumed in a simplified version by two stoichiometric compounds, TiAl2 and Ti2Al5[2003Sch].

N-Ti: The solubility of nitrogen both in ( Ti) and ( Ti) is significant. The congruently melting TiN1-x

compound with wide homogeneity range and incongruently melting Ti2N compound exist in this binary

system [Mas2]. According to the data of [1992Rog] the new phases Ti3N2-x and Ti4N3-x are also formed in

the N-Ti system.

Solid Phases

Three compounds ( 1, Ti2AlN, 2, Ti3AlN and 3, Ti4AlN3) are formed in this system among which 1,

Ti2AlN is the most stable [1995Wu] and belongs to the group of H phases [1964Now]. An excellent

agreement exists between the various determinations of the lattice parameters [1963Jei, 1976Ivc2, 1977Ivc,

1984Sch, 1985Sch, 1986Kau, 1995Wu, 1999Far, 2000Bar2, 2000Gam, 2001Per]. It has been observed to

exist over the temperature range from 700 to 1600°C and being deficient in nitrogen above 1300°C

324


MSIT®

Al–N–Ti

[1992Sch]. Its melting point is above at least 1950°C [1998Che], probably above 2500°C [1992Mab]. 1,

Ti2AlN is easily obtained by numerous ways: (I) from Al and Ti powders which react exothermically in

gaseous nitrogen to form Ti2AlN particles in a matrix of TiAl [1992Mab], (II) by hot pressing powder

mixtures of Al, Ti and TiN and homogenizing the samples at 850°C for 200 h [1963Jei], (III) from AlN and

Ti metal powder by sintering above 1500°C [1976Ivc2, 1977Ivc], (IV) by mixing elemental or binary

powders, followed by cold pressing, then hot pressing in sealed evacuated containers at temperatures from

1275 to 1600°C and pressures of up to 1100 MPa for up to 24 h [1999Far], (V) by nitriding Al-Ti alloys at

1000°C [1999Mag], (VI) by heating 2Ti+AlN mixtures at 1400°C for 48 h under a pressure of 40 MPa

[2000Bar2] and (VII) by reactive sintering AlN and Ti for 16 h under a vacuum of 10-3 Pa [2000Gam].

2, Ti3AlN which exhibits a negligible range of homogeneity [1984Sch] has a cubic structure with a lattice

parameter that varies only within experimental errors [1984Sch, 1985Sch, 1986Kau, 1992Sch]. This

compound becomes nitrogen deficient above 1300°C and melts incongruently at 1590 10°C,

decomposing presumably into either L+ TiN1-x+Ti2AlN or L+ TiN1-x [1998Che].

3, Ti4AlN3 is stable between ~1250 and 1500°C under Ar, but decomposes in air at 1400°C to form TiN

[2000Pro1]. It tends to be deficient in nitrogen Ti4AlN3- (where 0 < < 0.1) [1999Bar, 1999Ho, 2000Bar1,

2000Fin, 2000Pro2, 2000Raw]. The formulae Ti3Al2N2 [1984Sch, 1985Sch, 1992Sch, 1998Bar] and

Ti3Al1-xN2 [1997Lee] were initially accepted for this compound; however chemical analysis using energy

dispersive spectroscopy (EDS) unequivocally proved a stoichiometry of Ti4AlN3 [1999Bar]. Fully dense

polycrystalline samples of Ti4AlN3- were processed by mixing TiH2, TiN and AlN to the desired

stoichiometry [1999Bar, 1999Ho, 2000Bar1, 2000Pro1, 2000Pro2]. The mixed powders were cold-pressed

at ~200 MPa, sealed in evacuated borosilicate tubes and hot isostatically pressed at 1275°C for 24 h under

a pressure of ~70 MPa. To complete the reaction such samples were annealed further at a temperature of

1325°C for 168 h under an Ar atmosphere.

The solubility limit for nitrogen in TiAl alloys are lower than 0.1 at.%, because the precipitation of nitrides

occurs even at the smallest content of N in these alloys [1991Kaw]. Nitrogen solubility in Ti3Al should be

higher than 2.32 at.% [2001Per] and can be as high as 3.5 at.% [1997Dur]. Solid solution based on

aluminium does not hold detectable amount Ti, and TiN1-x dissolves very small amounts of Al [1984Sch].

Details of crystal structure of all solid phases are given in Table 1.


The phase diagram of the AlN-TiN pseudobinary sub-system, has been calculated using the model of

regular solutions for the solid phases and that of an ideal solution for the liquid phase [1988Hol, 1989Hol].

Figure 1 shows the calculated diagram modified to take into account the accepted melting point of TiN

(3290°C instead of 2930°C). The eutectic temperature and eutectic composition depend experimentally on

the dimension of both AlN and TiN particles [1991And] because of the possible formation of Ti1-xAlxN

metastable solid solutions. Hard coatings prepared by the cathodic arc ion plating method allow to form a

cubic solid solution Ti1-xAlxN (0 < x < 0.7) and a wurtzite type solid solution Ti1-xAlxN (0.8 < x < 1)

[1991Ike, 1992Tan]. The existing experimental results and thermodynamic calculations lead to the

so-called vapor deposition phase diagram, Fig. 2, [2001Spe].

Isothermal Sections

According to the calculations of [1984Bey] AlN-TiN, TiAl3-AlN and TiAl3-TiN are stable tie lines in the

Al-N-Ti ternary system at low temperatures. The sintering of Al with TiN powders leads to a hardening of

the alloy due to the formation of AlN and TiAl3 during sintering [1992Koy]. Figures 3 and 4 show

isothermal sections of the Al-N-Ti diagram at 900 and 1325°C, respectively. These sections were

constructed using the experimental data and accounting for the formation of Ti4AlN3 [1997Dur, 2000Pro2].

Ti2N does not coexist with any of the ternary compounds [1984Sch].

325


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Al–N–Ti


Substantial solid solubility of nitrogen in ( Ti) and ( Ti) solid solutions has been reported by [1954Thy].

Unfortunately, these ternary phase boundary data do not match with the currently established phase

boundaries in the Al-Ti and N-Ti binaries. Nitrogen raises the /( + ) phase boundaries toward higher

temperatures and widens the + field of the Al-Ti system.

Thermodynamics

Heat capacity of the Ti4AlN3- compound was measured between 2 and 10 K using a standard adiabatic

calorimeter in a liquid helium cryostat [1999Ho]. It was determined that Cp = 0.00812T + 0.033 10-3T3

J mol-1 K-1 and the characteristic Debye temperature ( D) equals 506°C. According to the data of

[2000Bar1] Cp = 232 - 24350T-1 from 25 to 1030°C and D = 498°C ( D = 489°C [2000Fin]. The molar

heat capacity at room temperature is 150 J mol-1 K-1 and increases monotonically with increasing

temperature, reaching a plateau at 220 J mol-1 K-1 at 1030°C. The Gibbs energy of formation of Ti2AlN

at 850°C equals -135.5 kJ mol-1 of atoms [1997Dur].


Ti2AlN is more wear-resistant than carbides of transition metals [1976Ivc1, 1977Ivc] and its abrasive ability

gives up only on diamond, B4C, B and BN. Composites containing 30 vol.% Ti2AlN and 70 vol.% TiAl

have a high strength at both room and elevated temperatures and show some intrinsic compressive ductility

at room temperature [1992Mab].

The yield strength and fracture stress increase with increasing nitrogen content in the TiAl phase

[1991Kaw]. At room temperature Young’s (ERT) and shear ( RT) moduli and Poisson’s ratio of Ti4AlN3-

are 310 2, 127 2 GPa and 0.2 respectively [2000Fin, 2000Pro1]. This ternary nitride is relatively soft

(Vickers hardness 2.5 GPa), lightweight (4.58 g cm-3) and machinable [2000Pro1].

Increasing the Al content in the Ti1-xAlxN metastable films leads to an increase coating roughness and a

change in color from gold to black-purple when the Al content increases from 13 to 27 mass% [1991Col].

Because of differences in chemical composition, the sputtered Ti1-xAlxN coatings show colors changing

from metallic silver for low nitrogen coatings to a very dark blue for layers with high nitrogen contents

[1986Jeh, 1988Jeh]; [1987Ina] indicates that these solid solutions in the composition range of

0.13 x 0.58 were greenish brown in color. These films have good decorative properties and excellent

wear as well, [1986Kno, 1987Kno, 1988Ran, 1992Tan]. The incorporation of Al into the nitride films

improves the oxidation resistance as well as the cutting performance of Ti1-xAlxN coated drills [1986Mue,

2001Spe]. It has been noted by [1990McI, 1991Ike, 1992Tan] that metastable single-phase polycrystalline

Ti0.5 Al0.5N alloy films exhibit much better high-temperature (750 - 900°C) oxidation resistance than

polycrystalline TiN1-x films grown under similar conditions. It was found that Ti1-xAlxN films upon

oxidation in air at 1000°C formed two-phase mixtures of TiO2 and Al2O3 [1991Ike, 2001Hug]. The

thickness of the oxide layer grown on these films decreases with increasing Al content in the films

[2001Hug]. The electric resistivity of Ti1-xAlxN metastable solid solutions raised with increasing Al content

[1987Ina]. Based upon resistivity and elevated-temperature interfacial reaction measurements, Ti1-xAlxN

appears to be a promising candidate for improved diffusion-barrier layers between Al and Si [1993Pet].

Metastable Ti1-xAlxN coatings with the cubic NaCl structure are already being produced commercially for

cutting tool applications [2001Spe].

When the amount of TiN particles was increased [2002Tan] the AlN-TiN composite materials showed an

increasing Vickers hardness (adding 21 vol.% TiN to AlN-ceramics increased the hardness more than 15%),

a decreasing fracture strength (20%) and a slightly increasing Young’s modulus (6%). Such composites

with high content of AlN (> 20 vol.%) have a great thermic stability against cyclic heating and cooling in

gas environments and in water [1976Kuz].

326


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Al–N–Ti

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327


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Al–N–Ti

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Al Based Hard Compound Film”, J. Vac. Sci. Technol., A4(6), 2695-2700 (1986)

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Defect Structure and Phase Transitions in Polycrystalline and Epitaxial Ti1-xAlxN and TiN

Films Grown be Reactive Magnetron Sputter”, Thin Solid Films, 205(2), 153-164 (1991)


[1991Ike] Ikeda, T., Satoh, H., “Phase Formation and Characterization of Hard Coatings in the

Ti-Al-N System Prepared by the Cathodic Arc Ion Platting Method”, Thin Solid Films,

195(1-2), 99-110 (1991) (Experimental, Crys. Structure, 17)

[1991Kaw] Kawabata, T., Tadano, M., Izumi, O., “Effect of Carbon and Nitrogen on Mechanical

Properties of TiAl Alloys”, ISIJ International, 31(10), 1161-1167 (1991) (Experimental,

Equi. Diagram, Mechan. Prop., 45)

[1992Koy] Koyama, K., Morishita, M., Suzuki, K., Yagi, S., “A New Ternary Al-Ti-N Alloy Prepared

by the Reaction Sintering Process” (in Japanese), J. Japan. Soc. Powder Powder Metall.,

39(10), 823-829 (1992) (Experimental, Equi. Diagram, 12)

328


MSIT®

Al–N–Ti

[1992Mab] Mabuchi, H., Tsuda, H., Nakayama, Y., “Processing of TiAl-Ti2AlN Composites and their

Compressive Properties”, J. Mater. Res., 7(4), 894-900 (1992) (Experimental, Mechan.

Prop., 21)

[1992Rog] Rogl, P., Schuster, J.C., “Ti-B-N (Titanium - Boron - Nitrogen)” in “Phase Diagrams of

Ternary Boron Nitride and Silicon Nitride Systems” (Monogr. Ser. of Alloy Phase Diag.),

Materials Park, Ohio: Materials Informations Soc., 103-106 (1992) (Review, Equi.

Diagram., Crys. Structure, Thermodyn., #, *, 19)

[1992Sch] Schuster, J.C., “System Aluminium - Nitrogen - Titanium: Summary of Constitution Data”,

Int. Report, (1992) (Review, Equi. Diagram, 18)

[1992Tan] Tanaka, Y., Guer, T.M., Kelly, M., Hagstrom, S.B., Ikeda, T., Wakihira, K., Satoh, H.,

“Properties of (Ti1-xAlx)N Coating Tools Prepared by the Cathodic Arc Ion Plating

Method”, J. Vac. Sci. Technol., A10(4), 1749-1756 (1992) (Experimental, Crys.

Structure, 21)

[1993Adi] Adibi, F., Petrov, I., Green., J.E., Wahlstroem, U., Sundaren, J.-E., “Design and

Characterization of a Compact Two-Target Ultrahigh Vacuum Magnetron Sputter

Deposition System: Application to the Growth of Epitaxial Ti1-xAlxN Alloys and

TiN/Ti1-xAlxN Superlattices”, J. Vac. Sci. Technol., A11(1), 136-142 (1993) (Experimental,


[1993Jeh] Jehn, H.A., “Aluminium-Nitrogen-Titanium”, MSIT Ternary Evaluation Program, in MSIT



Assessment, 29)

[1993Pet] Petrov I., Mojab, E., Adibi, F., Greene, J.E., Hultman, L., Sundgren, J.-E., “Interfacial

Reactions in Epitaxial Al/Ti1-xAlxN (0 x 0.2) Model Diffusion-Barrier Structure”, J. Vac.

Sci. Technol., A11(1), 11-17 (1993) (Experimental, Crys. Structure, 25)

[1993Tan] Tanaka, Y., Guer, T.M., Kelly, M., Hagstrom, S.B., Ikeda, T., “Strusture and Properties of

(Ti1-xAlx)N Films Prepared by Reactive Sputtering”, Thin Solid Films, 228(1-2), 238-241

(1993) (Experimental, Crys. Structure, 13)

[1993Wah] Wahlstroem, U., Hultman, L., Sundgren, J.-E., Adibi, F., Petrov, I., Greene, J.E., “Crystal

Growth and Microstructure of Polycrystalline Ti(1-x)AlxN Alloy Films Deposited by

Ultra-High-Vacuum Dual-Target Magnetron Sputtering”, Thin Solid Films, 235(1-2), 62-70


[1995Wu] Wu, Z.L., Pope, D.P., Vitek, V., “Ti2NAl in L12 Al3Ti-Base Alloys”, Metall. Mater. Trans.,

A26(3), 521-524 (1995) (Experimental, Crys. Structure, 15)

[1997Dur] Durlu, N., Gruber, U., Pietzka, M.A., Schmidt, H., Schuster, J.C., “Phases and Phase

Equilibria in the Quaternary System Ti-Cu-Al-N at 850°C”, Z. Metallkd., 97(5), 390-400

(1997) (Experimental, Review, Crys. Structure, Equi. Diagram, 32)

[1997Lee] Lee, H.D., Petuskey, W.T., “New Ternary Nitride in Ti-Al-N System”, J. Am. Ceram. Soc.,

80(3), 604-608 (1997) (Experimental, Crys. Structure, 8)

[1998Bar] Barsoum, M.W., Schuster, J.C., “Comment on “New Ternary Nitride in Ti-Al-N System”,

J. Am. Ceram. Soc., 81(3), 785-786 (1998) (Experimental, Crys. Structure, 10)

[1998Che] Chen, G., Sundman, B., “Thermodynamic Assessment of the Ti-Al-N System”, J. Phase

Equilib., 19(2), 146-160 (1998) (Assessment, Equi. Diagram, Thermodyn., 42)

[1998Lee] Lee, B.-J., “Predictive Analysis of Ti/AlN Interfacial Reaction Using Diffusion

Simulation”, Scr. Mater., 38(3), 499-507 (1998) (Calculation, Equi. Diagram, 15)

[1999Bar] Barsoum, M.W., Farber, L., Levin, I., Procopio, A., El-Raghy, T., Berner, A.,

“High-Resolution Transmission Electron Microscopy of Ti4AlN3, or Ti3Al2N2 Revisited”,

J. Am. Ceram. Soc., 82(9), 2545-2547 (1999) (Experimental, Crys. Structure, 23)

[1999Far] Farber, L., Levin, I., Barsoum, M.W., El-Raghy, T., Tzenov, T., “High-Resolution

Transmission Electron Microscopy of Some Tin+1AXn Compounds (n = 1, 2; A = Al or Si;

X = C or N)”, J. Appl. Phys., 86(5), 2540-2543 (1999) (Experimental, Crys. Structure, 23)

329


MSIT®

Al–N–Ti

[1999Ho] Ho, J.C., Hamdeh, H.H., Barsoum, M.W., El-Raghy, T., “Low Temperature Heat Capacities

of Ti3Al1.1C1.8, Ti4AlN3, and Ti3SiC2”, J. Appl. Phys., 86(7), 3609-3611 (1999)

(Experimental, Thermodyn., 15)

[1999Mag] Magnan, J., Weatherly, G.C., Cheynet, M.-C., “The Nitriding Behavior of Ti-Al Alloys at

1000°C”, Metall. Mater. Trans. A, 30A(1), 19-29 (1999) (Experimental, Equi. Diagram, 27)

[2000Bar1] Barsoum, M.W. Rawn, C.J., El-Raghy, T., Procopio, A.T., Porter, W.D., Wang, H.,

Hubbard, C.R., “Thermal Properties of Ti4AlN3”, J. Appl. Phys., 87(12), 8407-8414 (2000)

(Experimental, Crys. Structure, Phys. Prop., 33)

[2000Bar2] Barsoum, M.W., Ali, M., El-Raghy, T., “Processing and Characterization of Ti2AlC,

Ti2AlN and Ti2AlC0.5N0.5”, Metall. Trans. A, 31A(7), 1857-1865 (2000) (Experimental,

Crys. Structure, Phys. Prop., 36)

[2000Fin] Finkel, P., Barsoum, M.W., El-Raghy, T., “Low Temperature Dependencies of the Elastic

Properties of Ti4AlN3, Ti3Al1.1C1.8, and Ti3SiC2”, J. Appl. Phys., 87(4), 1701-1703 (2000)

(Experimental, Mechan. Prop., 22)

[2000Gam] Gamarnik, M.Y., Barsoum, M.W., El-Raghy, T., “Improved X-Ray Powder Diffraction

Data for Ti2AlN”, Powder Diffr., 15(4), 241-242 (2000) (Experimental, Crys. Structure, 7)

[2000Par] Paransky, Y., Gotman, I., Gutmanas, E.Y., “Reactive Phase Formation at AlN-Ti and

AlN-TiAl Interfaces”, Mater. Sci. Eng. A, A277, 83-94 (2000) (Experimental, Equi.

Diagram, 28)

[2000Pro1] Procopio, A.T., Barsoum, M.W., El-Ragny, T., “Characterization of Ti4AlN3”, Metall.

Mater. Trans. A, 31A(2), 333-337 (2000) (Experimental, Crys. Structure, Phys. Prop., 24)

[2000Pro2] Procopio, A.T., El-Raghy, T., Barsoum, M.W., “Synthesis of Ti4AlN3 and Phase Equilibria

in the Ti-Al-N System”, Metall. Mater. Trans. A, 31A(2), 373-378 (2000) (Experimental,

Equi. Diagram, Crys. Structure, 24)

[2000Raw] Rawn, C.J., Barsoum, M.W., El-Raghy, T., Procopio, A., Hoffmann, C.M., Hubbard, C.R.,

“Structure of Ti4AlN3 - A Layered Mn+1AXn Nitride”, Mater. Res. Bull., 35, 1785-1796


[2001Hug] Hugon, M.C., Varniere, F., Letendu, F., Agius, B., Vickridge, I., Kingon, A.I., “18O Study

of the Oxidation of Reactively Sputtered Ti1-xAlxN Barrier”, J. Mater. Res., 16(9),

2591-2599 (2001) (Experimental, Crys. Structure, Phys. Prop., 24)

[2001Per] Perdix, F., Trichet, M.-F., Bonnentien, J.-L., Cornet, M., Bigot, J., “Influence of Nitrogen

on the Microstructure and Mechanical Properties of Ti-48Al Alloy”, Intermetallics, 9,


[2001Spe] Spencer, P.J., “Computational Thermochemistry: from its Early Calphad Days to a

Cost-Effective Role in Materials Development and Processing”, Calphad, 25(2), 163-174

(2001) (Calculation, Equi. Diagram, 31)

[2002Tan] Tangen, I.-L., Grande, T., Yu, Y.D., Hoier, R., Einarsrud, M.-A., “Preparation and

Mechanical Characterisation of Aluminium Nitride-Titanium Nitride and Aluminium

Nitride-Silicon Carbide Composites”, Key Eng. Mater., 206-213, 1153-1156 (2002)

(Experimental, Mechan. Prop., 2)

[2003Fer] Ferro, R., Bochvar, N., Sheftel, E., Ding, J.J., “Al-N (Aluminum-Nitrogen)”, MSIT Binary







330


MSIT®

Al–N–Ti


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

( Al)

< 660.452

cF4

Fm3m

Cu

a = 404.96 at 25°C [Mas2]

( N)

< -237.54

cP8

Pa3

N

a = 566.1 [Mas2]

( Ti)

1670 - 882

cI2

Im3m

W

a = 330.65 [Mas2]

dissolves up to 6.2 at.% N at 2020°C


( Ti)

< 882

hP2

P63/mmc

Mg

a = 295.06

c = 468.35

at 25°C [Mas2]

dissolves up to 23 at.% N at 1050°C


AlN

< 2434.7

hP4

P63mc

ZnS (wurtzite)

a = 311.14

c = 497.92

at 25°C [2003Fer]

Ti2N

< 1100

tP6

P42/mnm

TiO2

a = 494.52

c = 303.42

at 33 to 34 at.% N [V-C2]

, TiN1-x

< 3290

cF8

Fm3m

NaCl

a = 423.9 0.1 [V-C2] From 28 at.% N at 2350°C to

> 50 at.% N

Ti3N2-x

1103 - 1066

hR2

?

VTa2C2

a = 297.95

c = 2896.5

at 29 at.% N [1992Rog]

Ti4N3-x

1291 - 1078

hR2

?

V4C3

a = 298.09

c = 2166.42

at 31.5 at.% N [1992Rog]

Ti3Al

1164

hP8

P63/mmc

Ni3Sn

a = 580.6

c = 465.5

a = 574.6

c = 462.4

at 22 at.% Al [2003Sch]

at 38 at.% Al [2003Sch]

TiAl

< 1463

tP4

P4/mmm

AuCu

a = 400.0

c = 407.5

a = 398.4

c = 406.0

at 50.0 at.% Al, [2003Sch]

at 62.0 at.% Al, [2003Sch]

TiAl2< 1199

tI24

I41/amd

HfGa2

a = 397.0

c = 2497.0

[2003Sch]

“Ti2Al5”

1416 - 990

tP28

P4/mmm

“Ti2Al5”

a = 390.53

c = 2919.63

[2003Sch]

331


MSIT®

Al–N–Ti

TiAl3(h)

< 1393

tI8

I4/mmm

TiAl3(h)

a = 384.9

c = 860.9

[2003Sch]

TiAl3(l)

< 950 (Ti-rich)

tI32

I4/mmm

TiAl3 (l)

a = 387.7

c = 3382.8

[2003Sch]

* 1, Ti3AlN cP5

?

CaTiO3

a = 411.20

a = 411.70 0.07

[1984Sch, 1985Sch]

[1992Sch]

* 2, Ti2AlN hP8

P63/mmc

Cr2AlC

a = 298.9

c = 1361.4

a = 299.9

c = 1365.0

a = 300.9

c = 1365.0

at 25°C [2000Bar2]

at 400°C [2000Bar2]

at 800°C [2000Bar2]

* 3, Ti4AlN3 hP16

P63/mmc

Ti4AlN3

a = 299.05 0.01

c = 2338.0 0.1

a = 300.45 0.02

c = 2348.1 0.2

a = 302.22 0.02

c = 2360.8 0.2

a = 298.80 0.02

c = 2337.2 0.2

a = 299.10 0.02

c = 2339.6 0.1

Nitrogen deficient Ti4AlN3-x [1999Bar]

at 25°C [2000Bar1]

at 570°C [2000Bar1]

at 1094°C [2000Bar1]

Ti4AlN2.78, neutron powder diffraction

[2000Raw]

X-ray powder diffraction [2000Raw]

* Ti1-xAlxN

metastable

cF8

Fm3m

NaCl

a = 424

a = 422.6

a = 420.6

a = 419.9

a = 416.9

a = 416

at x = 0.1 [1993Pet]

at x = 0.2 [1993Pet]

at x = 0.3 [1993Tan]

at x = 0.42 [1993Tan]

at x = 0.5 [1993Tan]

at x = 0.7 [1991Ike]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

332


MSIT®

Al–N–Ti

10 20 30 40

250

500

750

1000

1250

1500

1750

2000

2250

2500

2750

3000

3250

Ti 50.00Al 0.00N 50.00

Ti 0.00Al 50.00N 50.00Al, at.%

Te

mp

era

ture

, °C

L

TiN+AlN

TiN+L

L+AlN

TiN

AlN

2800°C

3290°C

2500°C

Fig. 1: Al-N-Ti.

Calculated phase

diagram of the

AlN - TiN

pseudobinary system

10 20 30 40

0

100

200

300

400

500

600

700

800

900

Ti 50.00Al 0.00N 50.00

Ti 0.00Al 50.00N 50.00Al, at.%

Te

mp

era

ture

, °C

(Ti,Al)N (Al,Ti)N

cubic hexagonal

cubic+hexagonal

Fig. 2: Al-N-Ti.

Metastable TiN - AIN

phase diagram

333


MSIT®

Al–N–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Al

N Data / Grid: at.%

Axes: at.%

AlN

TiN1-x

Ti2N

(αTi)

(βTi) Ti3Al TiAl TiAl

2TiAl

3

τ1

τ2

L

20

40

60

80

20 40 60 80

20

40

60

80

Ti Al

N Data / Grid: at.%

Axes: at.%

AlN

TiN1-x

(βTi) TiAl TiAl3

τ1

τ2

Ti2Al

5L

(αTi)

τ3

Fig. 3: Al-N-Ti.


900°C

Fig. 4: Al-N-Ti.


1325°C

334


MSIT®

Al–Nb–Ti

Aluminium – Niobium – Titanium

Ludmila Tretyachenko

Literature Data

Titanium aluminide based alloys are candidate materials for high temperature structural applications;

among alloying elements particularly niobium is expected to exert a favorable influence on low temperature

ductility. Data on phase equilibria in the Al-Nb-Ti system are a prerequisite to promote the development of

appropriate alloys.

The status of investigations in the Al-Nb-Ti system was summarized by [1993Gam] in a critical assessment

comprising all constitution-relevant literature data up to 1990. As a result the experimental data by

[1990Hel] and [1990Per] were chosen for the liquidus projection as well as for the isothermal sections at

1200 and 1000°C. Furthermore, a liquidus projection and several isothermal sections calculated by

[1992Kat1] were given. Phase relations at that time were characterized by a series of equilibrium phases:

(i) a wide region of the bcc disordered solid solution (Ti,Nb,Al) [1962Pop, 1970Nar, 1972Nar, 1974Nar,

1983Tro, 1984Zak, 1989Jew, 1989Kal, 1990Hel, 1990Per], which transforms to an ordered ternary solid

solution phase (B2 or 0) in a wide range of compositions [1987Ban, 1989Ben, 1990Hel, 1990Per]; (ii)

extended solid solution phases on the base of binary compounds TiAl ( ), Nb3Al ( ) and Nb2Al ( ), (iii) a

continuous solid solution between the binary boundary phases TiAl3 and NbAl3 (from now on designated

as ), (iv) solid solutions based on the low-temperature modification of titanium Ti ( ) and the Ti3Al based

phase ( 2) as well as (v) the ternary compounds Ti2NbAl (so-called O phase, discovered by [1988Ban])

[1990Moz, 1990Mur, 1990Wey, 1991Ben] and Ti4NbAl3 with the Ni2In type (B82) [1990Ben1, 1990Ben2,

1991Ben].

Although the general features of the phase relations remained unchanged, new investigations refined

various details in the constitution of the ternary system and furthermore solved a series of controversies,

which essentially concerned (a) the stability of ternary phases and (b) the extension of solid solution phases.

A listing of recent and some earlier experiments and the techniques used is presented in Table 1.

One of the problems is linked to the two ternary phases, T1 (Ti-18Nb-34Al) and T2 (Ti-11Nb-44Al),

reported by [1989Jew] in an isothermal section at 1200°C, which turned out to be part of ternary solutions:

T1 was shown to have the structure of the ordered bcc phase (B2 or 0 in this assessment [1990Per]), whilst

the T2 phase was supposed to be an isolated region of the same phase. The authors of [1990Per, 1990Ben2,

1990Kno, 1990Mis, 1990Wey] meanwhile agree that due to numerous phase transformations alloys in the

area of T1 and T2 are very sensitive to composition, temperature and the cooling rate.

The second problem is related with the so-called 1 phase. Although the TiNbAl3 ( 1) phase was reported

in the Ti-NbAl3 section by [1962Pop, 1983Tro, 1984Zak], which in the review by [1984Arg] was assumed

to be pseudobinary, the 1 phase was, however, not observed by the authors of [1989Jew, 1989Kal,

1990Per]. A study of diffusion couples at 1000°C [1990Hao1, 1990Hao2] again was interpreted in terms of

two ternary compounds, TiNbAl3 ( 1) and Ti5NbAl2 each with a large solubility range. Whilst the second

phase is to be identified with the O phase, Ti2NbAl [1988Ban, 1989Kes], the existence of the 1 phase was

denied in an investigation of partial isothermal sections at 1100, 900 and 800°C [1991Smi, 1992Smi,

1991Zak, 1992Zak, 1992Pav1, 1992Pav2]. Nevertheless, claim for the existence of the 1 phase was again

raised by [1993Zha] and [1994Che1] and a model of its crystal structure was reported by [1994Che2,

1994Wan]. Furthermore 1 phase fields were shown in the isothermal sections at 1000, 1150 and 1400°C

by [1996Che].

Despite [1997Jew] studied in detail the alloy Ti-23Nb-51Al (which was prepared by arc melting, annealed

at 1200°C for 180 h and then at 1150°C for 50 h and water quenched) by backscattered electron imaging

(BSEI), energy dispersive X-ray analysis (EDX) and XRD, and could not confirm the existence of 1, the

previous authors did not agree with the comment of [1997Jew] and again presented (i) the 1 phase in their

isothermal section at 1400°C [1998Wan], (ii) in refined versions of the sections at 1000 and 1150°C

[1998Din] and (iii) the crystal structure of the 1 phase [1998Che]. However, in a detailed reinvestigation

335


MSIT®

Al–Nb–Ti

of the isothermal sections at 1200 and 1000°C by [1998Hel] employing optical microscopy, EMPA, TEM

and XRD on diffusion couples and bulk samples neither TiNbAl3 ( 1) nor the phases T1 or T2 could be

traced. No other ternary compounds were observed. Considerable solid solubility of the third element in the

most of the binary phases were confirmed and refined. A separate area of the ordered B2 phase was detected

at 1000°C.

The third problem area covers (i) the so-called (orthorhombic) O phases near the composition Ti2NbAl and

(ii) the Ti4NbAl3 phase. As these problems are related to the crystallography of the phases mentioned, a

detailed discussion is included in the section “Solid Phases”.

Vaporization of solid alloys has been studied by Knudsen-effusion mass spectroscopy in the temperature

range between 897 and 1362°C to derive Ti, Al partial pressures and thermodynamic activities of Ti and Al,

partial enthalpies and entropies of mixing at 1200°C. Data on the phase compositions of 25 alloys in the

range adjoining to the Al-Ti side and containing up to ~30 at.% Nb are given for 1200, 1100 and 1000°C

[1999Eck].

Based on the experimental phase equilibria and thermodynamic data, thermodynamic assessments of the

Al-Nb-Ti system were performed by [1998Ser] using the Redlich - Kister polynomial to describe the excess

Gibbs energies of liquid, bcc and hcp phases. The intermetallic compounds, which exhibit a homogeneity

range, were modeled using two or three sublattices. The sublattice model was also used to describe the order

- disorder transformations D019 - hcp and A2 - B2. Both O1 and O2 forms were modeled as separate phases

with two and three sublattices, respectively. As a result a liquidus projection has been calculated, as well as

partial isothermal sections of the Nb rich corner at 700, 900 and 1200°C (in weight fractions), isothermal

sections at 700, 800, 900, 1000, 1020, 1060, 1100, 1150, 1175, 1200, 1400 and 1650°C (in at. fractions) the

isopleth at 27.5 at.% Al up to 35 at.% Nb. The representation of the thermodynamic properties of two states

of the orthorhombic phase, ordered O1 and disordered O2, with a unique function was proposed by

[2001Ser]. Two other models were proposed for thermodynamic modeling of the orthorhombic phase. The

two sets of thermodynamic parameters obtained according to both models were used to calculate the

isothermal sections at 990 and 700°C. Fields of B2 and bcc phase stability in the isothermal section at

1000°C were calculated using the CPA-GPM (coherent potential approximation - generalized perturbation

method) within the cluster variation method (CVM) [1993Rub] and with application of linear muffin-tin

orbitals (LMTO) [1995Rub]. The CVM in the irregular tetrahedron approximation was furthermore used to

calculate the limits of the B2 phase field at 800, 1000, 1200 and 1400°C [1996Jac, 1999Cha1] and in the

vertical section at 50 at.% Ti and 50 at.% Nb [1996Jac]. The results obtained were proven by experimental

studies [1999Cha2]. The results of [1993Rub] and [1999Cha1] were included in a review by [2001Col] and

used for the mixed CVM-CALPHAD method to calculate the phase equilibria in ternary system (isothermal

section at 1000°C). [2001Kan] applied the CVM in the octahedron and tetrahedron approximation to

calculate the / 2 and / 2 phase equilibria at 1000°C. The grand potential approach was applied to obtain

thermodynamic parameters used to calculate the / and / phase equilibria at 1150 and 1400°C

[2001Li1].

Binary Systems

The Al-Nb and Nb-Ti systems are accepted from [Mas2] and [1987Mur], respectively. A critical assessment

of the Al-Ti phase diagram is due to [2003Sch]. The version accepted therein and in [1993Oka1] is primarily

based on the work of [1992Kat2], which is in essential agreement with recent data by [1996Tre]. However,

the Ti5Al11 stoichiometry was shown in the latter phase diagram. Recently the Al-rich part of the system

has been reinvestigated by [2001Bra], who also has shown the Ti5Al11 phase to exist.

Solid Phases

Data of the solid phases in the Al-Nb-Ti system are given in Table 2. The bcc solid solution ( ) exists in a

wide range of composition up to 40 at.% Al [1995Zdz, 1998Hel, 2000Leo2, 2002Leo1]. The transformation

of the disordered (A2) phase to ordered 0 (B2) has been observed by many research groups [1987Ban,

1989Ben, 1992Men, 1994Hou, 1996Men, 1996Vas, 1998Rho, 1999Cha2] and others. The transition

temperatures were shown to be sensitive to composition [1999Cha2] with the highest ordering temperatures

336


MSIT®

Al–Nb–Ti

(>1600°C, close to melting temperature) found for alloys in the vicinity of Ti2NbAl [1996Vas]. There is

still some discrepancy on the ordering temperature for the Ti2NbAl alloy for which [1989Ben] reported a

temperature higher than 1400°C, [1990Hel] estimated 1000 to 1200°C but [1999Cha2] from in situ neutron

diffraction recorded only 1182 5°C.

The A2 B2 transformation temperature decreases from Ti2NbAl toward the Nb corner. The ordered 0

phase can be obtained in metastable form on quenching from the high temperature field and decomposes

at aging.

There are two well established ternary phases with the stoichiometries Ti2NbAl [1988Ban, 1990Moz] and

Ti4NbAl3 [1990Ben1, 1990Ben2, 1992Ben, 1996Sad, 2000Sad]. The orthorhombic Ti2NbAl based phase

O arises from the phase as a result of a sequence of phase transformations [1992Mur1, 1994Ben2,

1995Mur2, 1999Boe, 2001Sad, 2003Sad]. The formation of the O phase was suggested to occur

immediately from the B2 ( 0) phase [1989Kes, 1991Ben, 1994Ben2] as well as through the peritectoid

reaction 0+ 2 O [1992Mur1, 1992Mur2, 1995Mur1, 1999Boe, 2003Sad] below ~1000°C. The

orthorhombically distorted phase was observed at the 2/ 0-interface with the same composition and site

occupancy as the 2 phase; as a similar structure has been obtained in hydrogenated Ti3Al-Nb alloys, the

authors [1990Mur] concluded that the O phase appears as a result of hydrogen absorption during thin foil

preparation in the acid-containing electrolyte.

The homogeneity range of the O phase extends preferably at constant Al content of 26 - 27 at.%. The

orthorhombic phase was shown to exist in two forms, O1 and O2, with crystal lattice of the same space

group and lattice parameters, but with different site occupancies [1990Mur, 1990Wey, 1992Mur2,

1995Mur1, 1995Mur2, 2002Wu]. In the O1 form, which was observed to exist at higher temperatures from

~1000 down to ~900°C, Ti and Nb atoms randomly occupy the same sublattice (as in hexagonal 2), while

Nb atoms occupy a distinctive sublattice in the O2 form detected at temperatures below 900°C. The

transformation between these two forms was shown to be reversible. A first order transition was suggested

for the O1 O2 reaction [1995Mur2]. A very “weak” first order transition was predicted using the Bragg -

Williams model [2002Wu].

As to the O1 phase, it was suggested that the disordered orthorhombic martensite observed in the binary

Nb-Ti system can be stabilized into an equilibrium phase at certain Al and Nb contents in the ternary

Al-Nb-Ti system around the Ti2NbAl composition [1995Mur1].

[1994Ben1, 1994Ben2] have outlined possible paths for the constant composition coherent transformation

of bcc Ti2NbAl high temperature phases to the hcp or orthorhombic low temperature phases employing

crystallographic group-subgroup relations.

The Ni2In (B82) type phase Ti4NbAl3 is formed from the CsCl (B2) type phase at ~900°C. This phase was

found in the Ti-20Nb-30Al alloy annealed at 900°C by [1992Ben], however, in an in situ neutron diffraction

of the Ti-12.9Nb-36.5Al by [2000Sad] it was only revealed at temperatures at or below 800°C. A

thermodynamic calculation of the phase transformation in the Ti-10.8Nb-36.9Al alloy yielded Ti4NbAl3below 1060°C [1996Sad]. The transformation of the B2 phase to Ti4NbAl3 involves the formation of

metastable ´´ with trigonal structure [1990Ben1, 1990Ben2, 1990Sho, 1996Sad]. From TEM-analysis

[1990Ben2] reported also a new phase with a tripled hexagonal lattice for which he assumed further

substitutional ordering of the B82 type phase in terms of either a possible Ti5Ga4 type phase with 18 atoms

per unit cell and (Ti3Al3)(AlNb2) stoichiometry or in terms of the Mn5Si3 type structure (16 atoms/u.c.)

with (Ti3Al3)Nb2 stoichiometry. A phase with hexagonal structure (a = 579 pm, c = 1409 pm) was found

in the as cast alloy Ti4NbAl4 by means of TEM [1995Zdz]. It was supposed to be a superstructure of 2.

Formation of metastable phases ´ and ´´ during rapid cooling was observed in Ti3Al-Nb alloys containing

up to 5 at.% Nb [1988Str, 1990Wey, 1995Xu]. At higher Nb contents various metastable related phases,

both athermal and isothermal, have been detected in alloys rapidly cooled from high temperatures or aged

at ~350-550°C [1978Zak, 1982Str, 1988Str, 1991Li, 1992Hsi1, 1992Sur, 2000Leo2, 2000Sad, 2001Sad]

and others, as well as in Nb-Ti alloys with low Al content [1992Voz, 1996Men]. The ´ and ´´ phases are

described by [1990Ben1] as two configurations of the same trigonal P3ml phase. They are related to the

ordered B2 type phase and are distinguished by site occupancies. The ´ modification is considered as the

idealized state with the B2 chemical order inherited in a diffusionless transition. The chemical order in the

´´ configuration is changed but the space group is the same. This configuration is more stable.

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Al–Nb–Ti

The 1 phase (TiNbAl3), which earlier was reported by [1962Pop, 1983Tro, 1984Arg, 1984Zak] and not

confirmed by [1989Kal, 1989Jew, 1990Hel, 1990Per], was again reported by [1990Hao1, 1990Hao2,

1994Che2, 1994Wan, 1996Che, 1998Wan] and even later by [1998Che, 1998Din] in spite of the fact that

[1997Jew] once more disproved the existence of this phase. The crystal structure of 1 was identified as

tetragonal with a = 558 to 584 pm, c = 815 to 845 pm and has been considered as a superstructure of the

L10 structure of (TiAl). The transformation (L10) 1 was suggested to be a continuous ordering process

taking place with increasing Nb content in the (TiAl) phase. The ordering process has been presumed to

proceed as a first order transition at 1000 and 1150°C but as a second order transition at 1400°C. The

ordered 1 phase has been considered to be formed at the Nb content of 18 at.%, when Nb atoms occupy a

specific sublattice. A possible relation of 1 with Ti2Al5 (tP32, P4/mbm) [1982Mii] was suggested.

[1993Jac] detected in the Ti-20.3Nb-42.9Al alloy a high temperature phase with a Ti3Cu (L60) type lattice,

a modification of the CuAu (L10) type lattice. The L60 phase (tP4, space group P4/mmm) was found to

differ from the L10 phase in the site occupancy and was suggested to be an intermediate phase between the

high temperature phase and the lower temperature and or 2 phases.

Besides major amounts of the (Ti,Nb)Al3 phase, [1991Spa] claimed the formation of a cubic Cu3Au (L12)

type phase (composition Ti27.8Nb12.3Al60.9, a = 397.8 pm) using XRD, SEM and EMPA on the

Ti-12Nb-63Al alloy arc melted and annealed at 1200°C for 16 h. However, [1993Nak2] from optical

microscopy, XRD and SEM analyses did not confirm the L12 phase in the Ti-8Nb-67Al alloy sintered at

1150°C for 24 h.

A metastable ordered tetragonal transition phase T with a composition of Ti5NbAl2 arising during the B2

to 2 transition in a plasma sprayed Ti-11Nb-24Al alloy after aging for 10 min at 650°C was reported by

[1992Hsi1, 1992Hsi2, 1992Hsi3]. On prolonged aging the T phase transforms to an ordered O phase and

further to 2. The phase was detected by means of XRD, SEM and TEM. The crystal structure of the T phase

was found to be similar to the D03 type structure but with a tetragonal distortion (P4/mmm, a =

650 10 pm, c/a 1.02) and structural relationships and habit plane between T and O phases were

established. [1994Ban] analyzed the diffraction patterns obtained by [1992Hsi2, 1992Hsi3] and found that

they were not consistent with the proposed structure, but can be attributed to the structure of the metastable

O phase proposed by [1990Moz]. Recently a new phase evolution path during aging at 650°C was proposed

to be B2 B19 O´´ O´ 2, with O´´ and O´ phases instead of T and O (Ti2AlNb) phases involved in the

previous phase evolution path: B2 T O 2 [1995Hsi]. The phases taking part in the newly proposed phase

transition sequence were the following: B2 (Pm3m), a = 325; B19 (Pmmm), a = 325, b = c = 460; O´´

(Cmcm, previously T), a = 660, b = 920, c = 460; O´ (Cmcm, previously O), a = 605 b = 980, c = 473; 2

(P63/mmc), a = 580, c = 465 (the lattice parameters in pm). The B19 and O´´´ phases can only be resolved

with difficulties owing to overlapping peaks and weak reflection intensity, however a tetragonal distortion

of the B2 phase was detected.

A novel tetragonal phase, designated as , was observed by [2000Leo1] in Ti-Nb-40Al alloys (Ti from 24

to 36 at.%) aged below 1000°C for times up to 3600 h followed by water quenching. Phase identification

was performed by XRD and EMPA. Convergent beam electron diffraction yielded a bct cell and space

group I41/amd, a = 510.6 pm, c = 2816.8 pm, on the basis of which indexing of the X-ray powder pattern

was satisfactory. The composition was evaluated as ~25Ti-45Nb-30Al and orientation relationships

between and (TiAl) were determined. From the low concentration of elements (O, N, C) interstitial

contamination was ruled out. The phase was reported to be thermodynamically stable.

A hydride phase with the same crystal structure and nearly the same lattice parameters as the phase was

observed to replace the 2 phase in a Ti-48Al-2Cr-2Nb duplex alloy at hydrogen charging for 60 h at 12.8

MPa and 800°C. However, there is no reason to suppose a high content of H in the studied samples

[2000Leo1], in particular, for crushed powder samples analyzed by means of XRD.

Precipitates, which occurred in the single phase alloy Ti-5Nb-54Al containing < 900 ppm O2, were shown

by SEM and EDS analysis to be a cubic ternary Al-O-Ti compound with a = 690 pm [2001Cao].

A stress induced orthorhombic 9R phase was observed at incoherent twin or incoherent pseudotwin

boundaries of the phase in the Ti-10Nb-45Al alloy, which was hot-forged at 1050°C [1997Wan]. The

lattice parameters of the 9R phase were obtained from HRTEM as follows: a = 490 pm, b = 282 pm,

c = 2080 pm.

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Al–Nb–Ti

Formation of a martensite type fcc phase with a = 437 pm was observed at electrical polishing of thin foils

[1978Zak].


Continuous solid solubility between TiAl3 and NbAl3 was confirmed by experimental investigations

[1996Che, 1995Zdz, 1998Din, 1998Wan] and was accepted in the thermodynamic assessment of the

Al-Nb-Ti system by [1998Ser]. No new experimental data were reported on the melting temperatures within

the (Ti,Nb)Al3 solid solution. The TiAl3-NbAl3 section is shown in Fig. 1 taking into account calculated

liquidus temperatures reported by [1998Ser] for the Ti rich solid solutions. They are somewhat higher than

those shown earlier by [1990Per] and [1992Kat1], and are more reasonable for the phase with

stoichiometric composition. The highest values of solidus temperatures shown by [1990Per] are used to

draw the solidus.


There are five invariant equilibria with a liquid phase, but the type of only one of them is well established:

L+ + (U type). Various types have been proposed for other invariant equilibria (Table 3). In addition

the temperatures of the invariant equilibria are not well established and compositions of phases participating

in equilibria are not known. Different stoichiometry for the third aluminide, Ti5Al11, Ti2Al5 or Ti9Al23, has

been accepted in equilibria including the phases on the base of Al-rich titanium aluminides TiAl ( ) and

TiAl3 ( ).

The existence of a three phase invariant equilibrium, L + , was shown by [1990Per] and [1995Zdz] but

the temperature of this equilibrium was not established and the position of the maximum point on the

liquidus curve is different in [1990Per] and [1995Zdz].

The existence of an invariant equilibrium + + at ~1100°C was reported by [1989Kal].

[2002Leo1] suggested that the four-phase equilibrium + + +O occurs at 900°C from a convergence of

the + + and +O+ phase fields in the alloy Ti-37.5Nb-25Al.

The existence of the invariant peritectoid reaction 2+ 0 O at about 1000°C was proposed by [1995Mur1],

whilst an eutectoid reaction 0 O+ was considered by [2001Mis].

Liquidus Surface

The liquidus surface was presented earlier by [1989Kal, 1990Per] from experimental studies. A

thermodynamic calculation was performed by [1992Kat1]. The liquidus surface, shown in Fig. 2, was

constructed by [1995Zdz] and is similar to that of [1990Per]. The liquidus surface presented by [1992Pav1]

has not been constructed for the part of the phase diagram adjoining to the Al-Ti side. A peritectic reaction

L+ + was proposed. Recently [2000Leo1] reinvestigated the liquidus surface and has found that the

field of primary crystallization of the phase is wider than earlier reported.

Figure 3 shows the liquidus surface projection calculated by [1998Ser]. There are four maximum points,

which indicate the existence of three-phase pseudobinary reactions.

Isothermal Sections

Figure 4 shows the calculated section at 1650°C [1998Ser]. Experimental data [1992Men, 1996Men] for

Nb rich alloys show good agreement with calculated boundaries for the + region.

The isothermal section at 1400°C was presented by [1996Che, 1998Wan] (Fig. 5) from results of an

experimental study and was calculated by [1992Kat1] (shown also in [1993Gam]) and [1998Ser] (Fig. 6).

The calculated versions are in good agreement with each other, the existence of the ordered 0 phase is

shown in the latter version. The field of the questionable 1 phase is shown in Fig. 5. Boundaries / and

/ calculated by [2001Li1] are in better agreement with the data of [1996Che, 1998Wan] than with the

boundaries calculated by [1998Ser].

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Al–Nb–Ti

It should be noted that the (Ti,Nb)Al3 solid solutions ( ) have to remain in solid state in the whole

homogeneity range at 1400°C taking into account recent data on the melting temperature of TiAl3 (1408°C

[2003Kar], 1425°C [1996Tre, 1997Bul]).

The calculated isothermal section at 1300°C [2001Sad] is shown in Fig. 7. The structure of two alloys,

Ti-21.8Nb-29.7Al and Ti-31.7Nb-23.4Al, studied by [2001Sad], is consistent with the calculated section.

The experimental region in the vicinity of the phase field [2000Kai] (with insignificant corrections to

adjust to the accepted binary Al-Ti system) is presented in Fig. 8. There is agreement between calculated

and experimental Nb solubility in the phase. The isothermal section at 1260°C [2001Sad] is similar to that

at 1300°C.

The isothermal section at 1200°C shown in Fig. 9 was taken from [1995Zdz] with minor changes to comply

with the accepted binary Al-Ti system. This shows a good agreement with the calculated section (Fig. 10)

[1998Ser]. [1992Pav1] presented the isothermal section at 1200°C for the Nb rich side, where the Ti

solubility in Nb3Al and Nb2Al were found to be less than shown formerly and confirmed later [1989Jew,

1989Kal, 1990Per, 1990Hel, 1992Sur, 1998Hel, 1993Ebr, 1998Che, 2002Leo1].

Satisfactory agreement also exists between phase compositions of alloys investigated at 1200°C by

[1992Jac, 1993Ebr, 1993Nak1, 1994Che1, 1999Eck] and the phase equilibria proposed by [1995Zdz].

[2000Kai] suggested a slightly different configuration of the phase field and adjacent phase fields, but the

same Nb solubility in the phase, ~10 at.%.

The isothermal section at 1150°C was constructed from results of a diffusion couple study [1996Che,

1998Din] and calculated by [1998Ser] (Fig. 11). Unlike the predicted phase equilibria shown in Fig. 11,

those obtained by [1996Che] and modified by [1998Din] are characterized with an existence of equilibrium

between and phases and + + and + + phase fields as well as a separate region of the questionable

1 phase coexisting with the , and phases. The data by [1998Yu] on the + + phase field are

consistent with the prediction of [1998Ser].

The isothermal section at 1100°C has been represented by [1991Smi, 1992Smi, 1996Che, 1998Din] and

calculated by [1992Kat1, 1998Ser]. Opposite results were obtained for the phase equilibria in the Ti rich

part of the system by [1992Kat1] and [1998Ser]. The coexistence of the 2 and phases was shown by

[1992Kat1], while according to [1998Ser] (Fig. 12) the 2 and phase fields are separated by the 0 phase

field and the 0 phase coexists with the phase. [1991Smi, 1992Smi], who studied the Nb rich part of the

phase diagram, reported the existence of a + + field though none of the studied alloys was in this region

and directions of tie-lines show better correlation with the version by [1992Kat1] rather than [1998Ser]. The

version proposed by [1992Che, 1998Din] satisfactorily agrees with [1992Kat1] in the part adjacent to the

Nb-Ti side. The above mentioned 1 phase also was shown at higher Al contents [1992Che, 1998Din].

However, numerous results obtained for certain alloys are in agreement with the version by [1998Ser,

1989Ben, 1989Mur, 1990Ben1, 1990Ben2, 1991Ben, 1992Qua, 1994Hou, 1994Ben2, 1999Boe, 1999Eck,

2001Mis, 2002Leo1].

The calculated isothermal section at 1020°C [1998Ser] shown in Fig. 13 is consistent with the experimental

section at 1000°C [1998Hel] (Fig. 14). These versions were found to be more reliable than those reported

by [1990Hao1, 1994Kum, 1996Che, 1998Din]. [1990Hao1, 1996Che, 1998Din] have shown the not well

established 1 phase. [1994Kum] reported only a small part of the section including the O phase, what can

be explained by the temperature of formation of O slightly below 1000°C. The ternary phase in the region

of the existence of the O phase also was shown by [1990Hao1]. It should be noted that the coexistence of

the 2+ phases shown in Figs. 13 and 14 has been observed by [2001Sad], but this has not been reported

by a majority of researchers who studied phase transformations and structures of alloys in the appropriate

region.

The occurrence of the disordered O phase is shown in the calculated isothermal section at 990°C [2001Ser]

(Fig. 15).

Phase equilibria in the Nb rich part at 900°C have been studied by [1991Smi, 1992Pav1, 1992Pav2,

1992Smi, 1992Zak] and have been presented as a partial isothermal section. The Ti solubility in Nb3Al and

Nb2Al found seems to be too low. The calculated isothermal section at 900°C [1998Ser] is shown in Fig. 16.

[2003Sad] reported the content of Al to be ~22 at.% in the O phase coexisting with the 0 phase. According

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MSIT®

Al–Nb–Ti

to [1995Mur1], the largest extension of the homogeneity range of the O phase is along the isopleth at

~27.5 at.% Al.

An occurrence of the Ti4NbAl3 ternary compound, designated here as , is shown in the calculated section

at 800°C [1998Ser] (Fig. 17). The experimental partial isothermal sections were presented in [1991Smi,

1992Smi].

The calculated isothermal section at 700°C is accepted from [2001Ser] (Fig. 18).

The calculated Nb rich corner at 600°C is shown in Fig. 19 after [1998Ser].


The experimental isopleth for 27.5 at.% Al is shown in Fig. 20 on the base of that earlier proposed by

[1995Mur1]. The calculated version of the same isopleth [1998Ser] is presented in Fig. 21. The latter

version proceeds from the existence of the 2+ equilibrium, while the first one does not suggest this

equilibrium. The presented versions of the isopleth also differ by the position of the homogeneity range of

the O phase, which according to [1998Ser] is supposed to be located at an Al content less than 27.5 at.%.

The existence of the O phase in two forms is shown in both versions.

Thermodynamics

The thermodynamic activities of Ti and Al, as well as partial enthalpies and entropies of mixing were

evaluated from measurements of Ti and Al partial pressures using Knudsen effusion mass spectrometry

[1999Eck]. Among the twenty four Al-Nb-Ti alloys studied, more attention was paid to those within the ,

2+ and 2 phase fields. The measurements were carried out in the temperature range between 897 and

1362°C and thermodynamic properties were evaluated for the mean temperature of 1200°C. Figures 22 to

24 summarize the thermodynamic activities, partial enthalpies of mixing and partial entropies of mixing for

the TiAl based alloys. Additional data for the alloy series (Ti0.48-xNbxAl0.52, Ti0.35NbxAl0.65-x,

(Ti0.8Al0.2)1-xNbx, (Ti0.7Al0.3)1-xNbx and Ti0.67NbxAl0.33-x) are given in [1999Eck]. Thermodynamic

activities of Al and Ti were calculated using a two-sublattice quasi-subregular solution model for based

alloys (Ti0.32Al0.08)1-xNbx (0 < x < 0.2), Ti0.48-xNbxAl0.52 and Ti0.44-xNbxAl0.56 (0 < x < 0.15) [2001Wan].

The Gibbs free energies of the , and phases were described by a subregular solution model; interaction

parameters were calculated and used to calculate / and / phase equilibria at 1150 and 1400°C by a

grand potential approach [2001Li1].

The Gibbs energy of formation of phases in the Al-Nb-Ti system were derived by [1998Ser] from an

optimization procedure using all the available experimental data on thermodynamics and phase equilibria.

A multi-sublattice model was used to describe the ordered compounds, whilst solution phases were

described by means of Redlich - Kister polynomials.

The thermodynamic modeling of the orthorhombic phase was reanalyzed by [2001Ser]. A representation of

the thermodynamic properties of ordered and disordered states with a continuous function was applied. Two

different models of the orthorhombic phase were performed. The thermodynamic parameters used to model

the order/disorder transformation in the orthorhombic O phase were reported.


The increased interest in titanium aluminides is due to their promising properties, which make them

attractive for potential application as aerospace materials, in particular, for jet engine components. These

intermetallics are characterized with low density, good strength at elevated temperatures, high resistance to

oxidation, good creep properties. However, they exhibit poor ductility at room temperature and low fracture

toughness, which can both be significantly improved by additions of niobium.

An increased high temperature strength was reported for Nb additions to Ti3Al [1970And, 1972And], but

the variation of high temperature strength versus composition was found to exhibit a maximum at 3 mass%

Nb and a minimum at 15 mass% Nb. A Ti3Al based alloy with ~5 at.% Nb at 760°C after various heat

treatments exhibited a fine acircular Widmanstaetten structure yielding a very high mechanical strength

[1977Sas]. However, this structure is unstable at high temperatures and the strength decreases with time.

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Al–Nb–Ti

Nevertheless, the alloy revealed higher strength and ductility than Ti3Al. A similar structure was proposed

for Ti based alloys containing 13.5 - 15.3 % Al and 23.4 - 30 % Nb [1981Bla]. These alloys were suggested

for application up to 750°C.

Tensile tests at room and elevated temperatures and a study of the creep behavior at 650°C were carried out

by [1991Row1, 1991Row2] on Ti2NbAl based alloys: the best heat resistance was found after heat

treatment in the field. Significant strengthening and resistance to fracture have been achieved in alloys

with a Widmanstaetten O+ 0 structure. Similar results were obtained by [1990Kno] for Ti-11Nb-24Al

alloys. [1991Row3] proposed the Ti3Al based alloys Ti-(18-34)Nb-(18-30)Al, which were reported to

exhibit an elevated heat resistance and a hot stamping ability, for gas turbine components.

[1997Nak] reported a large tensile elongation (~16 - 28 %) at room temperature for Ti3Al based alloys.

Elongations up to 810 % have been achieved for a Ti-10Nb-25Al alloy after a deformation rate of 5 10-5 s-1

at 980°C [1992Yan].

The dislocation structure and deformation behavior of the O and 2 phases at RT and at 650°C were

examined as a function of the Nb concentration in the alloys Ti-21Nb-26Al and Ti-16Nb-25Al. The O phase

was found to deform on all slip systems observed in 2 in spite of the lower (orthorhombic) symmetry

[1991Ban, 1995Ban].

The alloys of Ti-11Nb-(24-26)Al are the most studied. Creep testing at 650°C was carried out to evaluate

the influence of cooling rate from the field on the steady state strain rate and time to rupture [1990Mis].

Deformation and fracture processes were examined by [1991Akk]. [1992Aco] studied microstructure and

microhardness of spot welds. Phase transformations resulting from laser and gas-tungsten-arc welding and

solid state processing have been characterized to optimize mechanical properties [1990Cie].

Plasticity of the Ti-25Nb-25Al alloy was improved after rapid quenching and disappeared after annealing

[1991Cha]. Slow cooling from the region followed by aging in the 2+ phase field resulted in the

formation of relatively stable Widmanstaetten structure and a good balance of compressive and tensile

properties of the forged Ti-11Nb-24Al alloy produced by powder metallurgy [1993Sob].

Dynamic material modeling (DMM) was used to analyze the mechanical behavior of the Ti-11Nb-25Al

alloy [1993Lon]. Unstable and stable flow zones were predicted by DMM and attributed to the O 2

transformation. Data of hot compression tests have been used to construct instability maps for

Ti-11Nb-25Al [1994Sag] and Ti-15Nb-25Al alloys [1998Sag]. [2000Mur] determined regimes of unstable

material flow during hot deformation of the Ti-15Nb-25Al alloy.

[1995Sem] reported on microstructure evolution during rolling of sheets of Ti-23Nb-22Al.

A significant increase of hardness (from ~270 VHN to ~440-470 VHN) was observed in the solution treated

Ti-22.8Nb-11.1Al alloy as a result of precipitation hardening [1992Qua]. The age hardening occurred in the

disordered matrix in the temperature range of 575 - 675°C due to the formation of lath-like 2 precipitates.

A similar increase of hardness as observed for the quenched Ti-60Nb-8Al alloy annealed at 600°C has been

attributed to the precipitating O phase [1992Voz].

Alloys on the base of (TiAl) have been discussed by [1989Kim] (phase relations, microstructure,

processing, mechanical properties, deformation and fracture, factors affecting ductility).

A possibility to improve the oxidation resistance of based alloys has been reported earlier by [1962Pop].

[1993Zha] reported two heat resistant alloys Ti-10Nb-45Al and Ti-8Nb-48Al, which were developed for

high temperature application. The specific strength of these alloys at 800 - 1100°C was found to be higher

than that of TiAl and superalloys (the compressive yield strength was about ~700 MPa at 800°C, 350 MPa

at 1100°C, the density was ~4.3 g cm-3). The alloys showed some ductility at room temperature and

oxidation resistance better than that of TiAl and Ti3Al. The 2 phase transformation, which occurred at

grain boundaries during high temperature stress rupture deformation, has been studied by [2000Che].

Internal friction at high temperature and creep measurements were carried out for a Ti-4Nb-46.5Al alloy

[2000Wel]. Planar fault energies and sessile dislocation configurations were studied in (Ti1-xAlx)1-yNby

alloys, 0.48 < x < 0.51, 0 < y < 0.02 [1996Woo].

High temperature strength (compression testing up to 1100°C) and oxidation behavior (at 900 - 1200°C in

air) of alloys in a wide composition range (Ti3Al - TiAl3 - NbAl3 - TiNbAl3) have been investigated by

[1992Che]. The alloys with 55 - 64 at.% Al and a Ti:Nb ratio of 2 to 5 yielded the highest oxidation

resistance besides high tensile strength.

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Al–Nb–Ti

Fracture toughness measurements and fractographic analysis were carried out to evaluate the toughening

mechanism of the phase with particles of and phases [1993Ebr].

Superconducting properties of solid solutions on the base of Nb3Al were studied by [1975Pan, 1975Sha,

1977Ale]. The critical temperature Tc of the superconducting transition was found to decrease down to ~9 K

with increasing Ti content up to ~13 at.% [1975Sha]. [1981Ish] investigated the influence of Al additions

on the critical current density in superconducting Nb-60Ti alloys.

Electrical resistivity and its temperature dependence in the range of 20 to 220°C, as well as emf in a couple

with Cu have been studied for Ti alloys containing up to 50 mass% Nb and 10 mass% Al. Aluminium

additions to Nb-Ti alloys resulted in a decrease of heat conductivity [1965Kal]. Electrical resistivity,

hardness and density of Ti3Al-Nb alloys (up to 50 mass% Nb) have been studied by [1970And].

Temperature dependence of the 0.2 % proof stress at a compression rate 10-4 s-1 for Ti0.25Nb0.75Al3 with

the D022 structure was presented by [1990Sau].

Miscellaneous

The effect of Nb on the phase equilibria and transformation behavior in Al-Ti alloys based on / 0, 2 and

phases has been discussed for development of advanced high temperature materials [1999Flo]. It was

pointed out that data reported on the phase transformation in the appropriate field of the phase diagram are

fragmentary and often they are mutually incompatible. This may be due to limitations in experimental

techniques or interstitial contamination. It can be added that elements of phase diagrams often contradict

requirements of the phase equilibria theory.

CCT - Curves-(Continuous Cooling Transformation)

Schematic curves of continuous cooling transformations were derived from a study of microstructure

occurring in the Ti-11Nb-24Al alloy during continuous cooling from 1230°C ( field) down to room

temperature by immersing a wedge-shaped specimen with a narrow end into ice water [1990Wey].

[1995Lon] used DTA (600 - 1300°C) and in situ high temperature XRD (600 - 1300°C) to investigate phase

stability during continuous heating/cooling of Ti-11Nb-25Al alloy. The sequence of the phase fields

2+ +O 2+ + was established at heating, the same fields were identified at cooling. The alloy was

in the 2+ +O region up to 850°C, the field was found to exist above 1200°C.

[2001Sad] constructed CCT diagrams for Ti-21.8Nb-27.9Al and Ti-31.7Nb-23.4Al alloys from samples,

which had been cooled from 1260°C with the rates from 100 to 0.25 K s-1, using dilatometry, DTA, XRD,

SEM, TEM and microhardness measurements. Out-of-equilibrium phase transformations were observed for

fast cooling, while quasi-equilibrium transformations were detected for lower cooling rates. The sequence

of transformations at a cooling rate of 0.25 K s-1 was established to be

0 + 0 2+ 0+ O+ 0+ + +O for Ti-21.8Nb-27.9Al and 0 + 0 2+ 0+ O+ 0+ for

Ti-31.7Nb-23.4Al. The CCT diagrams for the alloys Ti-21.8Nb-27.9Al and Ti-31.7Nb-23.4Al,

respectively, are shown in Figs. 25, 26. Three non-equilibrium phases, ´, ´´ and Om (a massive

orthorhombic phase, which formed by a diffusion-less mechanism and had the chemical composition of the

parent B2 phase) were observed. According to an in-situ neutron diffraction study [2000Sad] the transition

from 0+ 2+ to 2+ + in Ti-12.9Nb-36.5Al occurs between 800 and 960°C.

Atomic Structure and Electronic Structure

The electronic structure and the total energy of Ti2NbAl in B2 ( 0), D019 ( 2) and O structure were

calculated with the self-consistent tight binding linear muffin-tin orbital method [1999Rav]. The obtained

results were used to study the phase stability and cohesive properties of these phases. The B2 phase was

shown to be the most stable one. The presence of all these phases in equilibrium over a range of temperature

is possible because they are close in energy. The heats of formation H were calculated to be -0.239,

-0.208 and -0.036 (eV/atom) for the B2, D019 and O phases, respectively.

The linear muffin-tin orbital method was also employed to elucidate the atom site distribution in ordered

(TiAl) compounds (L10), TiXAl2 and Ti2AlX (X = transition metal) via calculation of the electronic

structure and total energies from first principles [1993Ers]. Niobium was found to preferentially substitute

on Ti sites thereby increasing c/a. Accordingly, preferential Nb substitution for Ti in TiAl was established

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Al–Nb–Ti

experimentally [1986Kon, 1991Moh, 1999Hao] and has been predicted using a thermodynamic approach

based on a Bragg - Williams model [1990Nan], a plane-wave pseudopotential method [1996Woo,

1998Woo] and CVM [2001Kan]. Ti/Nb substitution is also supported from the partial entropy of mixing for

the (Ti0.38Al0.62)1-xNbx (0 < x < 0.2 alloy series [1999Eck]. Ti substitution for Nb on Nb sublattice sites was

determined for the ordered 0 phase in five Nb rich alloys [2002Leo2]. The local atomic order in 0 phase

of Ti2NbAl composition has furthermore been studied using the Extended X-Ray Absorption Fine Structure

(EXAFS) [1996Sik]. The site composition was shown to be written as (Ti1.5Nb0.5)A(Ti0.5Nb0.5Al)B.

CVM in the irregular tetrahedron approximation, used to calculate the 800, 1000, 1200 and 1400°C

isothermal partial sections, revealed a Heusler type phase instead of the CsCl (B2) phase for the Ti rich

region with a miscibility gap between the ordered Heusler phase and the disordered W(A2) type phase

[1996Jac, 1999Cha1]. CVM was furthermore employed by [1999Cha1] to calculate the sublattice

occupation of the 0 phase by Nb, Ti and Al atoms. It was shown that Ti atoms occupy one sublattice, Al

atoms on the other but Nb atoms prefer one or both sublattices depending on the composition. In agreement

with the CVM calculation, a neutron diffraction study [1999Cha2] at room temperature and in situ up to

1600°C has shown that sublattice occupation of the 0 phase is sensitive to the concentration. [1987Ban]

determined the site occupancy in the ordered solid solution phase 0 in the Ti-10Nb-25Al alloy using

ALCHEMI. It was shown that Ti atoms occupy one of two possible sublattices whilst Al and Nb atoms are

found in the other one.

Theoretical and experimental investigations of sublattice substitution of Nb in (TiAl) and 2 (Ti3Al)

based alloys [1999Hao, 1999Yan] comparing binding energy data and the Bragg-Williams model with

ALCHEMI (Atom Location Channeling Enhanced Microanalysis) measurements were summarized by

[2000Yan]. ALCHEMI data prompted a strong preference of Nb atoms to substitute for Ti in both TiAl and

Ti3Al [1999Hao]. [1986Kon] confirmed the Nb/Ti substitution in the Ti3Al lattice. The ordering tie-line

(OTL) approach to represent sublattice occupations was adopted by [2000Ama]: the OTLs were determined

via the ALCHEMI method. It was suggested that the order-disorder transformation is a second-order phase

transformation.

Studies of corrosion

An addition up to 15 mass% Al to Nb alloys containing 20-40 mass% Ti significantly decreases the

oxidation rate at 1100°C [1991Pav]. Oxidation kinetics of a Ti-25Nb-50Al alloy was studied using

thermogravimetry in air, pure O2 and their mixture at 1300°C at the pressure of 100 kPa [1992Bra]. A study

of cyclic oxidation of a Ti-24Nb-14Al alloy by [1988Sub] demonstrated the benefits of a protecting TiAl3coating.

Stress corrosion cracking (SCC) was shown to occur for a Ti-11Nb-24Al ( 2+ ) alloy in methanol and

aqueous solutions and needs to be taken into account in developing and applying Ti3Al-Nb alloys

[1992Zha].

Electro-spark deposition (ESD) was used to produce crack-free TiAl3 aluminide coating on a Ti3Al-Nb

alloy (Ti-10.8Nb-24.1Al) to improve its high temperature oxidation resistance [2001Li2]. An Al plate was

used as an electrode material. Isothermal oxidation tests at 800 and 900°C in air proved the low oxidation

rate of the coating.

The use of Ti hydride instead of pure Ti for the synthesis of O phase based alloys by ball-milling resulted

in a reduced contamination with oxygen and nitrogen, in considerable particle refinement and it accelerated

the amorphization of the powders [2002Bou].

[1989Shi] investigated the hydrogenation behavior in Ti3Al observing a “hydride” phase in the

Ti-11Nb-24Al alloy. The crystal structure of this phase was not established but an orthorhombic distortion

of the hexagonal base structure was reported. [1992Roz] studied the influence of hydrogen on phase

transformations in Ti-11Nb-24Al. Cathode charging hydrogen resulted in the formation of a Ti3Al-H

hydride in a thin surface layer and induced cracking. Temperature and pressure dependencies of hydrogen

solubility in a Ti-11Nb-24Al alloy were reported by [1992Chu] and the hydrogenization behavior of three

alloys with compositions in the vicinity of Ti2NbAl was investigated by [2001Zha]; a beneficial effect of

the O phase on the hydrogenization properties was established, i.e. Hf becomes more negative with

increasing volume fraction of the O phase.

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Al–Nb–Ti

[2002Hod] studied the behavior of Ti-7Nb-6Al (mass%) alloy under simulated biological conditions

(specific ions, pH, temperature) i.e. the electrochemical characterization by impedance spectroscopy and

photoelectrochemistry of the passive film.

An investigation of the sulfidation process of TiAl-2Nb (at.%) alloy was undertaken in order to find out the

alloying element, which would improve oxidation resistance [2000Izu]. The sulfidation amount was found

to be close to that for binary TiAl.

Disordering of the phase with tetragonal lattice and a new phase formation with a smaller c/a ratio were

observed at a neutron irradiation treatment of a TiAl-Nb alloy [1986Ibr].

The diffusivity in the phase was estimated at 1200 and 1400°C using the diffusion couple method

[1996Ebr]: Ti seems to be the fastest species, Al having a mobility close to Ti and Nb being the slowest

species.

A bulk Ti-19.9Nb-14.6Al nanophase material with the structure of the O phase was synthesized and

consolidated from powders with structure produced by ball milling [1991Chr]. The grain size of the

consolidated material was ~10 nm, the density was 4.48 g cm-3 and Vicker´s hardness was 498 VHN.

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345


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Al–Nb–Ti

[1978Zak] Zakharova, M.I., Khatanova, N.A., Kozlovskaya, N.A., “An Investigation of Supersaturated

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at.% Ti Alloys” (in Japanese), J. Jpn. Inst. Met., 45, 517-524 (1981) (Experimental, 8)

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401-42, (1981) (Crys. Structure, Review, 5)

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the System Nb-Ti-Al at 1200-600ºC” (in Russian), Izv. Akad. Nauk SSSR, Met., 1, 200-202


[1986Ibr] Ibragimov, S.S., Kofanov, B.A., Melikhov, V.D., “Change of Intermetallic Phases Structure

in Three-Phase TiAl-Nb Alloy at Neutron Irradiation” (in Russian), Izv. Akad. Nauk

Kazakh. SSR, Ser. Fiz. Mat., (2), 29-33 (1986) (Experimental, 8)

[1986Kon] Konitzer, D.G., Jones, I.P., Fraser H.L., “Site Occupancy In Solid Solutions of Nb in the

Intermetallic Compounds TiAl And Ti3Al”, Scr. Metall., 20, 265-268 (1986) (Crys.


[1987Ban] Banerjee, D., Nandy, T.K., Gogia, A.K., “Site Occupation in the Ordered -Phase of

Ternary Ti-Al-Nb Alloys”, Scr. Metall., 21, 597-600 (1987) (Crys. Structure,

Experimental, 21)

[1987Mur] Murray, J.L., “Nb-Ti (Niobium - Titanium)”, in “Phase Diagrams of Binary Titanium

Alloys”, ASM Publication, 188-194 (1987) (Crys. Structure, Equi. Diagram, Review, 44)

[1988Ban] Banerjee, D., Gogia, A.K., Nandi, T.K., Joshi, V.A., “A New Ordered Orthorhombic Phase

in a Ti3Al-Nb Alloy”, Acta Metall., 36, 871-882 (1988) (Crys. Structure, Experimental, 22)

[1988Has] Hashimoto, K., Doi, H., Kasahara, K., Tsujimoto, T., Suzuki, T., “Effects of the Third

Elements on the Structures of TiAl-Based Alloys” (in Japanese), J. Jpn. Inst. Met., 52,


[1988Str] Strychor, R., Williams, J.C., Soffa, W.A., “Phase Transformations and Modulated

Microstructures in Ti-Al-Nb Alloys”, Metall. Trans., 19A, 225-234, (1988) (Equi. Diagram,

Experimental, 44)

[1988Sub] Subrahmanyam, J., “Cyclic Oxidation of Aluminized Ti-14Al-24Nb Alloy”, J. Mater. Sci.,

23, 1906-1010 (1988) (Experimental, 7)

[1989Ben] Bendersky, L.A., Boettinger, W.J., “Investigation of B2 and Related Phases in the Ti-Al-Nb

Ternary System”, Mater. Res. Soc. Symp. Proc., 133, 45-50 (1989) (Equi. Diagram,

Experimental, 6)

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Al–Nb–Ti

[1989Jew] Jewett, T.J., Lin, J.C., Bonda, N.R., Seitzman, L.E., Hsieh, K.C., Chang, A.Y., Perepezko,

J.H., “Experimental Determination of the Titanium-Niobium-Aluminum Phase Diagram at

1200°C”, Mater. Res. Soc. Symp. Proc., 133, 69-74 (1989) (Equi. Diagram,

Experimental, 8)

[1989Kal] Kaltenbach, K., Gama, S., Pinatti, D.G., Schulze, K., Henig, E.-T., “A Contribution to the

Ternary System Al-Nb-Ti”, Z. Metallkd., 80, 535-539 (1989) (Equi. Diagram,

Experimental, 13)

[1989Kes] Kestner-Weykamp, H.T., Ward, C.H., Broderick, T.F., Kaufman, M.J., “Microstructures

and Phase Relationships in the Ti3Al+Nb System”, Scr. Metall., 23, 1697-1702 (1989)


[1989Kim] Kim, Y.-W., “Intermetallic Alloys Based on Titanium Aluminide”, JOM, 41, 24-30

(1989) (Review, 61)

[1989Mur] Muraleedharan, K., Banerjee, D., “Alloy Partitioning in Ti-24Al-11Nb Analytical Electron

Microscopy”, Metall. Trans., 20A, 1139-1142 (1989) (Equi. Diagram, Experimental, 10)

[1989Shi] Shih, D.S., Scarr, G.K., Wasielewski, G.E., “On Hydrogen Behavior in Ti3Al”, Scr. Metall.,

23, 973-978 (1989) (Experimental, 13)

[1990Ben1] Bendersky, L.A., Boettinger, W.J., Burton, B.P., Biancaniello, F.S., “The Formation of

Ordered -Related Phases in Alloys of Composition Ti4Al3Nb”, Acta Metall. Mater., 38,


[1990Ben2] Bendersky, L.A., Burton, B.P., Boettinger, W.J., Biancaniello, F.S., “Ordered

-Derivatives in a Ti-37.5Al-12.5Nb (at.%) Alloy”, Scr. Metall. Mater., 24, 1541-1546


[1990Cie] Cieslak, M.J., Headly, T.J., Baeslack III, “Effect of Thermal Processing of the

Microstructure if Ti-26Al-11Nb: Application to Fusion Welding”, Metall. Trans., 21A,

1273-1286 (1990) (Experimental, 27)

[1990Hao1] Hao, S., Zhao, Q., “Investigation of the 1000°C Isothermal Section of Ti-Al-Nb Ternary

Phase Diagram” (in Chinese), Proc.: 6th National Symp. Phase Diagrams, Shenyang,

China, 1990, 141-143 (1990) (Equi. Diagram, Experimental, 4)

[1990Hao2] Hao, S., Zhao, Q., “A New Ternary Intermetallic Compound in Ti-Al-Nb System”, Proc.:

6th National Symp. Phase Diagrams, Shenyang, China, 1990, 144-145, 149 (1990) (Crys.


[1990Hel] Helwig, A., “Experimental Study About the Constitution of the Aluminium - Titanium -

Niobium System” (in German), Ph.D. Thesis, University of Dortmund (1990)

(Experimental, 68) as quoted by [1993Gam]

[1990Kno] Knorr, D.B., Stoloff, N.S., “Effect of Heat Treatment on Microstructure and Texture in

Ti-24 at.% Al-11at.% Nb”, Mater. Sci. Eng., A123, 81-87 (1990) (Experimental, 23)

[1990Mis] Misra, R.S., Banerjee, D., “On the Influence of Cooling Rate in Solution Treatment for a

Ti-25Al-11Nb Alloy”, Scr. Metall. Mater., 24, 1477-1482 (1990) (Experimental, 18)

[1990Moz] Mozer, B., Bendersky, L.A., Boettinger, W.J., “Neutron Powder Diffraction Study of the

Orthorhombic Ti2AlNb Phase”, Scr. Metall. Mater., 24, 2363-2368 (1990) (Crys. Structure,

Experimental, 10)

[1990Mur] Muraleedharan, K., Naidu, C.V.N., Banerjee, D., “Orthorhombic Distortion of the 2 Phase

in Ti3Al-Nb Alloys: Artifacts and Facts”, Scr. Metall. Mater., 24, 27-32 (1990) (Crys.


[1990Nan] Nandy, T.K., Banerjee, D., Gogia, A.K., “Site Substitution of TiAl Intermetallic”, Scr.

Metall. Mater., 24, 2019-2022 (1990) (Crys. Structure, Theory, Thermodyn., 13)

[1990Per] Perepezko, J.H., Chang, Y.A., Seitzman, L.E., Lin J.C., Bonda, N.R., Jewett, T.J.,

Mishurda. J.C., “High Temperature Phase Stability in the Ti-Al-Nb System”, in “High

Temperature Aluminides and Intermetallics”, Wang, S.H., Liu, C.T., Pope, D.P., Stiegler,

J.O., (Eds.), The Minerals, Metals and Materials Society, 19-47 (1990) (Equi. Diagram,

Experimental, 20)

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Al–Nb–Ti

[1990Sau] Sauthoff, G., “Intermetallic Alloys-Overview on New Materials Developments for

Applications in West Germany”, Z. Metallkd., 81, 855-861 (1990) (Review, 36)



Review, 33)

[1990Sho] Shoemaker, C.B., Shoemaker, D.P., Bendersky, L.A., “Structure of -Ti3Al2.25Nb0.75”,

Acta Crystallogr., Sect. C: Cryst. Struct. Commun., C46(3), 374-377 (1990) (Crys.


[1990Wey] Weykamp, H.T., Baker, D.R., Paxton, D.M., Kaufman, M.J., “Continuous Cooling

Transformations in Ti3Al+Nb Alloys”, Scr. Metall. Mater., 24, 445-450 (1990), (Crys.


[1991Akk] Akkurt, A.S., Liu, G., Bond, G.M., “Micromechanisms of Deformation and Fracture in a

Ti3Al-Nb Alloy”, Mater. Res. Soc. Symp. Proc.: High-Temp. Ordered Intermetallic

Alloys IV, 213, 455-460 (1991) (Crys. Structure, Experimental, 11)

[1991Ban] Banerjee, D., Rowe, R.G., Hall, E.L., “Deformation of the Orthorhombic Phase in Ti-Al-Nb

Alloys”, High-Temp. Ordered Intermetallic Alloys IV, Mater. Res. Soc. Symp. Proc.:

Johnson, L.A., Pope, D.P., Stiegler, J.O., (Eds.), 213, 285-290 (1991) (Crys. Structure,

Experimental, 16)

[1991Ben] Bendersky, L.A., Boettinger, W.J., Roytburd, A., “Coherent Precipitates in the

B.C.C./Orthorhombic Two-Phase Field of the Ti-Al-Nb System”, Acta Metall. Mater., 39,


[1991Cha] Chang, C.P., Loretto, M.H., “The Decomposition Process of Rapidly Solidified Ti-25 at.%

Al-25 at.% Nb”, Philos. Mag. A, 63, 389-406 (1991) (Crys. Structure, Experimental, 23)

[1991Chr] Christman, T., Jain, M., “Processing and Consolidation of Bulk Nanocrystalline Titanium

Aluminide”, Scr. Metall. Mater., 25, 767-772 (1991) (Crys. Structure, Experimental, 32)

[1991Li] Li, D., Zhou, J., Chang, X., Guan, S., “On the Ordering Transformations in Ti3Al-Nb

Alloy”, Acta Metall. Sin. (China), 4A(3), 204-208 (1991) (Equi. Diagram, Experimental, 6)

[1991Moh] Mohandas, E., Beaven, P.A., “Site Occupation of Nb, V, Mn and Cr in -TiAl”, Scr. Metall.


[1991Pav] Pavlov, A.V., Zakharov, A.M., Karsanov, G.V., Vergasova, L.L., “An Influence of Al and

Si upon Heat Resistivity of Nb-Ti Alloys at 1100°C” (in Russian), Izv. Vyss. Uchebn.

Zaved., Tsvetn. Metall., (5), 89-94 (1991) (Experimental, 17)

[1991Row1] Rowe, R.G., Hall, E.L., “Stress-Assisted Discontinuous Precipitation during Creep of

Ti3Al-Nb Alloys”, High-Temp. Ordered Intermetallic Alloys IV, Mater. Res. Soc. Symp.

Proc., 213, 449-454 (1991) (Experimental, 11)

[1991Row2] Rowe, R.G., Konitzer, D.G., Woodfield, A.P., Chesnutt, J.C., “Tensile and Creep Behavior

of Ordered Orthorhombic Ti2AlNb-Based Alloys”, High-Temp. Ordered Intermetallic

Alloys IV, Mater. Res. Soc. Symp. Proc., 213, 703-708 (1991) (Experimental, 10)

[1991Row3] Rowe, R.G., “Tri-Titanium Aluminide Alloys Containing at Least Eighteen Atom Percent

Niobium”, Pat. 5032357 USA, Cit. by Ref. J. Metallurgiya, (10), Abs. 10I449P (1992) (in

Russian)

[1991Smi] Smirnova, T.R., Zakharov, A.M., Oleinikova, S.V., Filipyeva, O.A., “Phase Composition of

Alloys in the Nb-Ti-Al System with 0-20 % Al and Ti:Nb 1 at 1100-800°C” (in Russian),

Izv. Vyss. Uchebn. Zaved., Tsvetn. Metall., 4, 93-100 (1991) (Crys. Structure, Equi.


[1991Spa] Sparks, C.J., Porter, W.D., Schneibel, J.H., Oliver, W.C., Golec, C.G., “Formation of Cubic

L12 Phases from Al3Ti and Al3Zr by Transition Metal Substitutions for Al”, Mater. Res.

Soc. Symp. Proc., 186, 175-180 (1991) (Crys. Structure, Experimental, 15)

[1991Zak] Zakharov, A.M., Pavlov, A.V., Kachanova, T.L., “The Molybdenum Influence on the Phase

Composition of the Nb-Ti-Al Alloys at 1400 - 1600°C” (in Russian), Izv. Akad. Nauk SSSR,

Met., (3), 102-106 (1991) (Equi. Diagram, Experimental, 10)

348


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Al–Nb–Ti

[1992Aco] Acoff, V.L., Thompson, R.G., Griffin, R.D., Radhakrishnan, B., “Effect of Heat Treatment

on Microstructure and Microhardness of Spot Welds in Ti-26Al-11Nb”, Mater. Sci. Eng.,

A152, 304-309 (1992) (Abstract) (Experimental, 5)

[1992Ben] Bendersky, L.A., Boettinger, W.J., Biancaniello, F.S., “Intermetallic Ti-Al-Nb Alloys

Based on Strengthening of the Orthorhombic Phase by -Type Phases”, Mater. Sci. Eng.,

152A, 41-47 (1992) (Experimental, 15)

[1992Bra] Brady, M.P., Nanrahan, R.J. (Jr.), Elder, R.S.P., Verink, E.D. (Jr.), “The Effect of Nitrogen

on the Oxidation Behavior of 25Nb-25Ti-50Al”, Scr. Metall. Mater., 26, 767-770 (1992)

(Experimental, 6)

[1992Che] Chen, G., Sun, Z., Xhou, X., “Oxidation and Mechanical Behavior of Intermetallic Alloys

in the Ti-Nb-Al Ternary System”, Mater. Sci. Eng., 153, 597-601 (1992) (Experimental, 6)

[1992Chu] Chu, W.-Y., Thompson, A.W., Williams, J.C., “Hydrogen Solubility in a Titanium

Aluminide Alloy”, Acta Metall. Mater., 40, 455-462 (1992) (Experimental, 38)

[1993Gam] Gama, S., “Aluminium - Niobium - Titanium”, MSIT Ternary Evaluation Program, in MSIT



Assessment, 22)

[1992Hsi1] Hsiung, L.M., Cai., W., Wadley, H.N.G., “Microstructure and Phase Evolution in

Rapidly-Solidified Ti-24Al-11Nb”, Mater. Sci. Eng., 152, 295-303 (1992)

(Experimental, 14)

[1992Hsi2] Hsiung, L.M., Wadley, H.N.G., “A New Ordered Tetragonal Phase in the Ti3Al+Nb

System”, Scr. Metall. Mater., 26, 35-40 (1992) (Crys. Structure, Experimental, 10)

[1992Hsi3] Hsiung, L.M., Wadley, H.N.G., “Structural Relationships between the T and O Phases in

Ti-24Al-11Nb”, Scr. Metall. Mater., 26, 1071-1076 (1992) (Crys. Structure, Experimental,

Theory, 7)

[1992Hsi4] Hsiung, L.M., Wadley, H.N.G., “Stability of the Ordered Orthorhombic Phase in

Ti-24Al-11Nb”, Scr. Metall. Mater., 27, 605-610 (1992) (Crys. Structure, Experimental, 9)

[1992Jac] Jackson, A.G., Lee, D.S., “Characterization of the Phases Present in a Ti-45 at.% Al-10 at.%

Nb Alloy”, Scr. Metall. Mater., 26, 1575-1579 (1992) (Crys. Structure, Experimental, 8)

[1992Kat1] Kattner, U.R., Boettinger, W.J., “Thermodynamic Calculation of the Ternary Ti-Al-Nb

System”, Mater. Sci. Eng., A152, 9-17 (1992) (Equi. Diagram, Thermodyn., #, 20)

[1992Kat2] Kattner, U.R., Lin, J.-C., Chang, Y.A., “Thermodynamic Assessment and Calculation of the

Ti-Al System”, Metall. Trans., 23A, 2081-2090 (1992) (Equi. Diagram, Review, Theory,

Thermodyn., #, 51)

[1992Kim] Kimura, M., Hashimoto, K., Morikawa, H., “Study on Phase Stability in Ti-Al-X Systems

at High Temperatures”, Mater. Sci. Eng., A152, 54-59 (1992) (Equi. Diagram,

Experimental, 12)

[1992Men] Menon, E.S.K., Subramanian, P.R., Dimiduk, D.M., “Phase Equilibria in Niobium Rich

Nb-Al-Ti Alloys”, Scr. Metall. Mat., 27, 265-270 (1992) (Equi. Diagram, Experimental, 22)

[1992Mur1] Muraleedharan, K., Gogia, A.K., Nandy, T.K., Banerjee, D., Lele, S., “Transformation in a

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Ti-24Al-15Nb Alloy: Part II. A Composition Invariant 0 O Transformation”, Metall.

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[1992Qua] Quatrocchi, L.S., Koss, D.A., Scarr, G., “Precipitation Hardening of Beta Titanium Alloy

by the Alpha-Two Phase”, Scr. Metall. Mater., 26, 267-272 (1992) (Experimental, 10)

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Ti-24Al-11Nb Alloys”, J. Mater. Sci., 27, 2273-2278 (1992) (Experimental, 13)

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Alloys of the Nb-Ti-Al System with 0-20% Al And Ti:Nb Ratio 1 at 1100-800°C” (in

Russian), Izv. Vyss. Uchebn. Zaved., Tsvetn. Metall., (1-2), 91-98 (1992) (Crys. Structure,


[1992Sur] Surynarayana, C., Lee, D.S., “Phase Relations in Ti-Al-Nb Alloys at 1200°C”, Scr. Metall.

Mater., 26, 919-924 (1992) (Crys. Structure, Equi. Diagram, Experimental, 16)

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and Properties of Ti-60% Nb Alloy” (in Russian), Fiz. Met. Metalloved., (11), 108-113


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Titanium Aluminides”, Mater. Sci. Eng., A153, 457-464 (1992) (Experimental, 22)

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System in the Concentration Range 25-40% Ti and 0-20% Al”, Russ. Metall., (5), 102-105

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[1992Zha] Zhang, Y., Wang, Y.-B., Chu, W.-Y., Hsiao, C.-M., Thompson, A.W., “Stress Corrosion

Cracking of Titanium Aluminide Alloys in Aqueous Solutions and Methanol”, Scr. Metall.

Mater., 26, 925-928 (1992) (Experimental, 8)

[1993Ebr] Ebrahimi, F., Hoelzer, D.T., Castillo-Gomez, J.R., “Fracture Toughness of +x

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Modelling of Nb, V, Cr and Mn Substitutions in TiAl. I: c/a Ratio and Site Preference”,

Intermetallics, 1, 99-106 (1993) (Crys. Structure, Theory, 31)

[1993Jac] Jackson, A.G., “Identification of the L60 Phase in a -Ti-Al-Nb Alloy”, Scr. Metall. Mater.,


[1993Lon] Long, M., Rack, H.J., “Thermo-Mechanical Stability of Forged Ti-25Al-11Nb (at.%)”,

Mater. Sci. Eng., A170, 215-226 (1993) (Experimental, Theory, 30)

[1993Mur] Muraleedharan, K., Banerjee, D., “Phase Transformations Involving the 2 and O Phases in

Ti-Al-Nb Alloys”, Scr. Metall. Mater., 29, 527-532 (1993) (Experimental, 16)

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[1993Nak2] Nakayama, Y., Mabuchi, H., “Formation of Ternary L12 Compounds in Al3Ti Base

Alloys”, Intermetallics, 1, 41-48 (1993) (Crys. Structure, Experimental, 40)

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[1993Oka2] Okamoto, H., “Al-Ti (Aluminium - Titanium)”, J. Phase Equilib., 14, 764 (1993) (Equi.

Diagram, Review, 5)

[1993Rub] Rubin, G., Finel, A., “Calculation of Phase Diagrams of Ternary Systems with Cluster -

Variation - Method Entropy”, J. Phys.: Condens. Matter, 5, 9105-9120 (1993) (Theory,

Thermodyn., 34)

[1993Sob] Soboyejo, W.O., “An Investigation of the Effect of the Heat Treatment on the

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[1993Zha] Zhang, W.-J., Chen, Q.-Z., Wang, Y.-D., Sun, Z.-Q., “Characteristics of Heat Resistant

Alloys Ti10Nb45Al and Ti18Nb48Al”, Scr. Metall. Mater., 28, 1113-1118 (1993) (Crys.


[1994Ban] Banerjee, D., “Is There an Ordered Tetragonal Phase in the Ti3Al-Nb System?”, Scr. Metall.

Mater., 30, 855-858 (1994) (Crys. Structure, Theory, 14)

[1994Ben1] Bendersky, L.A., Roytburd, A., Boettinger, W.J., “Phase Transformations in the (Ti, Al)3Nb

Section of the Ti-Al-Nb System. - I. Microstructural Predictions Based on a Subgroup

Relation between Phases”, Acta Metall. Mater., 42, 2323-2335 (1994) (Crys. Structure,

Theory, 36)

[1994Ben2] Bendersky, L.A., Boettinger, W.J., “Phase Transformations in the (Ti, Nb)3Al Section of the

Ti-Al-Nb System. - II. Experimental TEM Study of Microstructures”, Acta Metall. Mater.,

42, 2337-2352 (1994) (Crys. Structure, Equi. Diagram, Experimental, Thermodyn., 19)

[1994Che1] Chen, Z., Jones, I.P., Saunders, N., Small, C.J., “Characterization of Phases in Ti-42Al-8Nb

Alloy at 1200°C”, Scr. Metall. Mater., 30, 1403-1408 (1994) (Equi. Diagram,

Experimental, 9)

[1994Che2] Chen, G.L., Wang, J.G., Sun, Z.Q., Ye, H.Q., “Continuous Ordering in the TiAl+Nb

System”, Intermetallics, 2, 31-36 (1994) (Crys. Structure, Equi. Diagram,

Experimental, 24)

[1994Hou] Hou, D.H., Yang, S.S., Shyue, J., Fraser, H.L., “Investigation of B2 and Related Phases in

Ti-Modified Nb-Al Alloys”, Mater. Res. Soc. Symp. Proc., 322, 437-442 (1994) (Crys.


[1994Kum] Kumar, S.G., Reddy, R.G., Brewer, L., “Phase Equilibria in Ti3Al-Nb Alloys at 1000°C”,

J. Phase Equilib., 15, 279-284 (1994) (Equi. Diagram, Experimental, 16)

[1994Sag] Sagar, P.K., Banerjee, D., Prasad, Y.V.R.K., “Processing of an -2 Aluminide Alloy,

Ti-24Al-11Nb”, Mater. Sci. Eng., A177, 185-197 (1994) (Experimental, 21)

[1994Wan] Wang, J., Chen, G., Sun, Z., Ye, H., “Structure of a New Ordered Ternary Intermetallic

Compound in TiAl+Nb System” (in Chinese), Acta Metall. Sin. (China), 30, A525-A531


[1995Ban] Banerjee, D., “Deformation of the O and 2 Phases in the Ti-Al-Nb System”, Philos.

Mag. A., 72, 1559-1587 (1995) (Experimental, 39)

[1995Hsi] Hsing, L.M., Wadley, H.N.G., “Time-Temperature Transformation Behavior of

Ti-24Al-11Nb”, Mater. Sci. Eng., A192/193, 908-913 (1995) (Crys. Structure,

Experimental, 11)

[1995Lon] Long, M., Rack, H.J., “Phase Stability During Continuous Heating/Cooling of

TiAl-(Nb, V, Mo) Titanium Aluminide Alloys”, Mater. Sci. Technol., 11, 150-158 (1995)


[1995Mur1] Muraleedharan, K., Nandy, T.K., Banerjee, D., “Phase Stability and Ordering Behaviour of

the O phase in Ti-Al-Nb Alloys”, Intermetallics, 3, 187-199 (1995) (Crys. Structure, Equi.

Diagram, Experimental, Theory, #, 30)

[1995Mur2] Muraleedharan, K., Banerjee, D., “The 2-to-O Transformation in Ti-Al-Nb Alloys”,

Philos. Mag., 71, 1011-1036 (1995) (Crys. Structure, Experimental, Theory, 24)

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Alloy Phase Diagram Predictions”, J. Phys.: Condens. Matter, 7, 3139-3152 (1995) (Equi.

Diagram, Theory, 30)

[1995Sem] Semiatin, S.L., Smith, P.R., “Microstructural Evolution During Rolling of Ti-22Al-23Nb

Sheet”, Mater. Sci. Eng., A202, 26-35 (1995) (Experimental, 13)

[1995Xu] Xu, R., Li, D., Cui, Y., Xu, D., Li, Q., Hu, Z., “A New Phase in Rapidly Solidified

Ti3Al-Based Alloys”, Scr. Metall. Mater., 32, 305-308 (1995) (Crys. Structure,

Experimental, 6)

351


MSIT®

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[1995Zdz] Zdziobek, A., Durand-Charre, M., Driole, J., Durand, F., “Experimental Investigation of

High Temperature Phase Equilibria in the Nb-Al-Ti System”, Z. Metallkd., 86, 334-340

(1995) (Crys. Structure, Equi. Diagram, Experimental, #, 23)

[1996Che] Chen, G.L., Wang, X.T., Ni, K.Q., Hao, S.M., Cao, J.X., Ding, J.J., Zhang, X.,

“Investigation on the 1000, 1150 and 1400°C Isothermal Section of the Ti-Al-Nb System”,

Intermetallics, 4, 13-22 (1996) (Crys. Structure, Equi. Diagram, Experimental, 27)

[1996Ebr] Ebrahimi, F., Ruiz-Aparicio, J.G.L, “Diffusivity in the Nb-Ti-Al Ternary Solid Solution”,

J. Alloys Compd., 245, 1-9 (1996) (Experimental, 16)

[1996Jac] Jacob, V., Colinet, C., Desre, P., Moret, F., “Calculation of the A2/B2 Phase Boundary in

the Nb-Ti-Al System with the Cluster Variation Method” (in French), J. Phys. IV, Col. 2, 6,

C2-3-C2-10 (1996) (Equi. Diagram, Theory, 17)

[1996Men] Menon, E.S.K., Subramanian, P.R., Dimiduk, D.M., “Phase Transformations in Nb-Al-Ti

Alloys”, Metall. Mater. Trans., 27A, 1647-1659 (1996) (Crys. Structure, Equi. Diagram,

Experimental, 30)

[1996Sad] Sadi, F., Servant, C., “Transformation During Continuous Cooling of the Ti4Al3Nb Alloy”

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Experimental, 14)

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C2-15-C2-30 (1996) (Crys. Structure, Experimental, 8)

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[1996Vas] Vasudevan, V.K., Yang, J., Woodfield, A.P., “On the to B2 Ordering Temperature in a

Ti-22Al-26Nb Orthorhombic Titanium Aluminide”, Scr. Mater., 35, 1033-1039 (1996)


[1996Woo] Woodward, C., MacLaren, J.M., “Planar Fault Energies and Sessile Dislocation

Configurations in Substitutionally Disordered Ti-Al with Nb and Cr Ternary Additions”,

Philos. Mag. A, 74, 337-357 (1996) (Crys. Structure, Theory, 27)

[1997Bul] Bulanova, M., Tretyachenko, L., Golovkova, M., “Phase Equilibria in Ti-Rich Corner of the

Ti-Si-Al System”, Z. Metallkd., 88, 256-267 (1997) (Crys. Structure, Equi. Diagram,


[1997Jew] Jewett, T.J., “Comment on “Investigation on the 1000, 1150 and 1400°C Isothermal Section

of the Ti-Al-Nb System””, Intermetallics, 5, 157-159 (1997) (Crys. Structure, Equi.


[1997Nak] Naka, S., Khan, T., “Designed Novel Multiconstituent Intermetallics: Contribution of

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(1997) (Review, 17)

[1997Wan] Wang, J.G., Zhang, L.C., Chen, G.L., Ye, H.Q., “Formation of Stress-Induced 9R Structure

in a Hot-Deformed Ti-45Al-10Nb Alloy”, Scr. Mater., 37, 135-140 (1997) (Crys. Structure,

Experimental, 21)

[1998Che] Chen, G.L., Wang, J.G., Wang, X.T., Ni, X.D., Hao, S.M., Ding, J.J., “Reply to the

“Comment on Investigation on the 1000, 1150 and 1400°C Isothermal Section of the

Ti-Al-Nb System”” - Part I. Ordering of Nb in -TiAl and 1-Phase”, Intermetallics, 6,


[1998Din] Ding, J.-J., Hao, S.-M., “Reply to the “Comment on Investigation on the 1000, 1150 and

1400°C Isothermal Section of the Ti-Al-Nb System”” - Part II. Modification of 1000 and

1150°C Isothermal Sections of the Ti-Al-Nb System”, Intermetallics, 6, 329-334 (1998)


[1998Hel] Hellwig, A., Palm, M., Inden, G., “Phase Equilibria in the Al-Nb-Ti System at High

Temperatures”, Intermetallics, 6, 79-94 (1998) (Crys. Structure, Equi. Diagram,


352


MSIT®

Al–Nb–Ti

[1998Rho] Rhodes, C.G., “Order/Disorder Temperature of the bcc Phase in Ti-21Al-26Nb”, Scr.

Mater., 38, 681-685 (1998) (Equi. Diagram, Experimental, 10)

[1998Sag] Sagar, P.K., Prasad, Y.V.R.K., “Hot Deformation and Microstructural Evolution in an 2/O

Titanium Aluminide Alloy Ti-25Al-15Nb”, Z. Metallkd., 89, 433-441 (1998)

(Experimental, 25)

[1998Ser] Servant, C., Ansara, I., “Thermodynamic Assessment of the Al-Nb-Ti System”, Ber.

Bunsenges. Phys. Chem., 102, 1189-1205 (1998) (Equi. Diagram, Review,

Thermodyn., #, 76)

[1998Tak] Takeyama, M., Ohmura, Y., Kikuchi, M., Matsuo, T., “Phase Equilibria and Microstructural

Control of TiAl Based Alloys”, Intermetallics, 6, 643-646 (1998) (Equi. Diagram,

Review, 20)

[1998Wan] Wang, X.T., Chen, G.L., Ni, K.Q., Hao, S.M., “The 1400°C Isothermal Section of the

Ti-Al-Nb Ternary System”, J. Phase Equilib., 19, 200-205 (1998) (Equi. Diagram,


[1998Woo] Woodward, C., Kajihara, S., “Site Preferences and Formation Energies of Substitutional Si,

Nb, Mo, Ta, and W Solid Solution in L10 Ti-Al”, Phys. Rev. B, 57, 13459-13470 (1998)

(Crys. Structure, Theory, Thermodyn., 45)

[1998Yu] Yu, T.H., Koo, C.H., “Phase Characterization of a Hot-Rolled Ti-40Al-10Nb Alloy at 1000

to 1200°C”, Scr. Mater., 39, 915-922 (1998) (Equi. Diagram, Experimental, 9)

[1999Boe] Boehlert, C.J., “The Phase Evolution and Microstructural Stability of an Orthorhombic

Ti-23Al-27Nb Alloy”, J. Phase Equilib., 20, 101-108 (1999) (Equi. Diagram,

Experimental, 17)

[1999Cha1] Chaumat, V., Colinet, C., Moret, F., “Study of Phase Equilibria in the Nb-Ti-Al System

Theoretical Study: CVM Calculation of the Phase Diagram of bcc Nb-Ti-Al”, J. Phase

Equilib., 20, 389-398 (1999) (Equi. Diagram, Experimental, Theory, 22)

[1999Cha2] Chaumat, V., Ressouche, E., Ouladdiaf, B., Desre, P., Moret, F., “Experimental Study of

Phase Equilibria in the Nb-Ti-Al System”, Scr. Mater., 40, 905-911 (1999) (Crys. Structure,


[1999Eck] Eckert, M., Kath, D., Hilpert, K., “Thermodynamic Activities in the Alloys of the Ti-Al-Nb

System”, Metall. Mater. Trans., 30A, 1315-1326 (1999) (Equi. Diagram, Experimental,

Thermodyn., #, 44)

[1999Flo] Flower, H.M., Christodoulou, J., “Phase Equilibria and Transformation in Titanium

Aluminides”, Mater. Sci. Technol., 15, 45-52 (1999) (Equi. Diagram, Review, 46)


Elements in TiAl and Ti3Al Alloys”, Acta Mater., 47, 1129-1139 (1999) (Crys. Structure,


[1999Rav] Ravi, C., Vajeeston, P., Mathijaya, S., Asokamani, R., “Electronic Structure, Phase Stability

and Cohesive Properties of Ti2XAl (X = Nb, V, Zr)”, Phys. Rev. B, 60, 15683-15690 (1999)

(Crys. Structure, Equi. Diagram, Theory, 32)

[1999Yan] Yang, R., Hao, Y.L., “Estimation of ( + 2) Equilibrium in Two-Phase Ti-Al-X Alloys by

Means of Sublattice Site Occupancies of X in TiAl and Ti3Al”, Scr. Mater., 41, 341-346


[2000Ama] Amancherla, S., Banerjee, R., Banerjee, S., Fraser, H.L., “Ordering in Ternary B2 Alloys”,

Inter. J. Refract. Met. Hard Mater., 18, 245-252 (2000) (Equi. Diagram, Theory, 23)

[2000Che] Cheng, Z.Y., Du, X.W., Zhu, J., Cao, C.X., Sun, F.S., “ 2 Phase Transformation in

Fractured High Temperature Stress Rupture Ti-48Al-2Nb (at.%)”, J. Mater. Sci., 35,

4501-4505 (2000) (Experimental, 22)

[2000Izu] Izumi, T., Yoshioka, T., Hayashi, S., Narita, T., “Sulfidation Properties of TiAl-2 at.% X

(X = V, Fe, Co, Cu, Nb, Mo, Ag and W) Alloys at 1173 K and 1.3 Pa Sulfur Pressure in an

H2S-H2 Gas Mixture”, Intermetallics, 8, 891-901 (2000) (Experimental, 42)

353


MSIT®

Al–Nb–Ti

[2000Kai] Kainuma, R., Fujita, Y., Mitsui, H., Ishida, K., “Phase Equilibria Among (hcp), (bcc)

and (L10) Phases in Ti-Al Base Ternary Alloys”, Intermetallics, 8, 855-867 (2000) (Equi.


[2000Leo1] Leonard, K.J., Mishurda, J.C., Molloseau, B., De Graef, M., Vasudevan, V.K.,

“Identification of a New Tetragonal Phase in the Nb-Ti-Al System”, Philos. Mag. Lett., 80,


[2000Leo2] Leonard, K.J., Mishurda, J.C., Vasudevan, V.K., “Examination of Solidification Pathways

and the Liquidus Surface in the Nb-Ti-Al System”, Metall. Mater. Trans. B, 31B,

1305-1321 (2000) (Crys. Structure, Equi. Diagram, Experimental, #, 21)

[2000Mur] Murty, S.V.S., Rao, B.N., Kashyap, B.P., “Development of a Processing Map for the Hot

Working of Ti-25Al-15Nb”, Z. Metallkd., 91, 769-774 (2000) (Theory, 8)

[2000Oka] Okamoto, H., “Al - Ti (Aluminium - Titanium)”, J. Phase Equilib., 21, 311 (2000) (Equi.

Diagram, Review, 2)

[2000Sad] Sadi, F.-A., Servant, C., “In Situ Neutron Diffraction on the Alloy 50.6Ti-36.5Al-12.9Nb

(at.%)”, Z. Metallkd., 91, 504-509 (2000) (Crys. Structure, Equi. Diagram,

Experimental, 21)

[2000Wel] Weller, M., Chatterjee, A., Haneczok, G., Clemens, H., “Internal Friction of -TiAl Alloys

at High Temperature”, J. Alloys Compd., 310, 134-138 (2000) (Experimental, 15)

[2000Yan] Yang, R., Hao, Y., Song, Y., Guo, Z.X., “Site Occupancy of Alloying Additions in Titanium

Aluminides and Its Application to Phase Equilibrium Evaluation”, Z. Metallkd., 91,

296-301 (2000) (Crys. Structure, Equi. Diagram, Review, Theory, 38)


Binary System Ti-Al”, Metall. Mater. Trans., 32A, 1037-1047 (2001) (Crys. Structure,


[2001Cao] Cao, G.H., Liu, Z.G., Shen, G.J., Liu, J.-M., “Identification of a Cubic Precipitate in

-Titanium Aluminides”, J. Alloys Compd., 325, 263-268 (2001) (Crys. Structure,

Experimental, 16)

[2001Col] Colinet, C., “Applications of the Cluster Variation Method to Empirical Phase Diagram

Calculations”, Calphad, 25, 607-623 (2001) (Equi. Diagram, Review, Theory,

Thermodyn., 108)

[2001Kan] Kang, S.Y., Onodera, H., “Analyses of HCP/D019 and D019/L10 Phase Boundaries in

Ti-Al-X (X = V, Mn, Nb, Cr, Mo, Ni, and Co) Systems by the Cluster Variation Method”,

J. Phase Equilib., 22, 424-430 (2001) (Equi. Diagram, Theory, 15)

[2001Li1] Li, J., Jiang, M., Hao, S., Li, S., Zhong, Z., “Thermodynamic Calculation of the / and /

Phase Equilibria in the Ti-Al-Nb Ternary System” (in Chinese), Acta Metall. Sin. (China),

37, 1064-1068 (2001) (Equi. Diagram, Thermodyn., Theory, 13)

[2001Li2] Li, Zh., Gao, W., He, Y., “Protection of a Ti3Al-Nb Alloy by Electro-Spark Deposition

Coating”, Scr. Mater., 45, 1099-1105 (2001) (Experimental, 23)

[2001Mis] Mishurda, J.C., Vasudevan, V.K., “An Estimate of the Kinetics of the 0 to Orthorhombic

Phase Transformation in the Nb-Ti-Al System”, Scr. Mater., 45, 677-684 (2001) (Equi.


[2001Sad] Sadi, F.A., Servant, C., Cizeron, G., “Phase Transformations in Ti-29.7Al-21.8Nb and

Ti-23.4Al-31.7Nb (at.%) Alloys”, Mater. Sci. Eng., A311, 185-199 (2001) (Crys. Structure,

Equi. Diagram, Experimental, #, 20)

[2001Ser] Servant, C., Ansara, I., “Thermodynamic Modelling of the Order-Disorder Transformation

of the Orthorhombic Phase of the Al-Nb-Ti System”, Calphad, 25, 509-525 (2001) (Equi.

Diagram, Theory, Thermodyn., #, 17)

[2001Sun] Sun, F.-S., Cao, C.-X., Kim, S.-E., Lee, Y.-T., Yan, M.-G., “Alloying Mechanism of Beta

Stabilizers in a TiAl Alloy”, Metall. Mater. Trans., 32A, 1573-1589 (2001) (Crys. Structure,


354


MSIT®

Al–Nb–Ti

[2001Wan] Wang, X., Chang, H., Lei, M., “Thermodynamic Aspects of Oxidation for Nb Alloying

-TiAl Intermetallic Compounds”, Acta Metall. Sin. (China), 37, 810-814 (2001) (Theory,

Thermodyn., 20)

[2001Zha] Zhang, L.T., Ito, K., Vasudevan, V.K., Yamaguchi, M., “Beneficial Effects of O-Phase on

the Hydrogen Absorption of Ti-Al-Nb Alloys”, Intermetallics, 9, 1045-1052 (2001) (Crys.

Structure, Equi. Diagram, Experimental, Thermodyn., 13)

[2002Bou] Bououdina, M., Guo, Z.X., “Characterization of Structural Stability of

(Ti(H2)+22Al+23Nb) Powder Mixtures During Mechanical Alloying”, Mater. Sci. Eng.,

A332, 210-222 (2002) (Crys. Structure, Experimental, 20)

[2002Hod] Hodgson, A.W.E., Mueller, Y., Forster, D., Virtanen, S., “Electrochemical Characterization

of Passive Films on Ti Alloys under Simulated Biological Conditions”, Electrochim. Acta,

47, 1913-1923 (2002) (Experimental, 55)

[2002Leo1] Leonard, K.J., Mishurda, J.C., Vasudevan, V.K., “Phase Equilibria at 1100°C in the

Nb-Ti-Al System”, Mater. Sci. Eng., A329-331, 282-288 (2002) (Crys. Structure, Equi.


[2002Leo2] Leonard, K.J., Vasudevan, V.K., “Site Occupancy Preferences in the B2 Ordered Phase in

Nb-Rich Nb-Ti-Al Alloys”, Mater. Sci. Eng., A329-331, 461-467 (2002) (Crys. Structure,


[2002Wu] Wu, B., Shen, J., Chu, M., Shang, Sh., Zhang, Z., Peng, D., Liu, S., “The Ordering

Behaviour of the O Phase in Ti2AlNb-Based Alloys”, Intermetallics, 10, 979-984 (2002)

(Crys. Structure, Theory, Thermodyn., 10)

[2003Kar] Karpets, M.V., Milman, Yu.V., Barabash, O.M., Korzhova, N.P., Senkov, O.N.,

Miracle, D.B., Legkaya, T.N., Voskoboynik, I.V., “The Influence of Zr Alloying on the

Structure and Properties of Al3Ti”, Intermetallics, 11, 241-249 (2003) (Crys. Structure,

Experimental, 16)

[2003Sad] Sadi, F.A., Servant, C., “On the B2 O Phase Transformation in Ti-Al-Nb Alloys”, Mater.

Sci. Eng., A346, 19-28 (2003) (Crys. Structure, Equi. Diagram, Experimental, Theory, 28)

[2003Sch] Schmid-Fetzer, R., “Al - Ti (Aluminium - Titanium)”, MSIT Binary Evaluation Program,

in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services,


Table 1: Experimental Investigations after [1993Gam] and some Earlier Works

Achievement Sample Preparation and Characterization References

isothermal section at 1200°C 14 arc-melted alloys, annealed at 1200°C for

two weeks; diffusion couples; LOM,

SEM-EMPA, XRD

[1989Jew]

isothermal sections at 1700 and 750°C;

< 30 at.% Ti and 30 at.% Al; 2.5 - 3.0 at.%

Ti solubility in Nb3Al

36 arc-melted alloys, annealed at 1700°C for

25 h and at 750°C for 500 h; LOM, XRD

[1975Pan]

structure of Nb rich alloys alloys at 25 and 20 at.% Al, containing < 5

at.% Ti; annealed at 800°C, 500 h, water

quenched after homogenization at 1700°C for

300 h. LOM, XPD and EMPA

[1975Fed]

Nb3Al based solid solutions; Ti solubility

in Nb3Al at 700°C > 10 at.%

(Nb,Ti)3Al up to ~14 at.% Ti; arc-melted,

homogenized at 1650°C for 3 h, annealed at

700°C for 250 h; XRD, Tc (superconducting

transition)

[1975Sha]

355


MSIT®

Al–Nb–Ti

phase equilibria in the Nb rich region up to

40 mass% Al and 40 mass% Ti (~70 at.%

Al, 50 at.% Ti); partial liquidus and solidus

projections and isothermal sections at

1200, 900, 600°C; invariant equilibria L +

+ at 1950°C, L + + at

1750°C and L + + at 1470°C

arc melted alloys, step-wise annealed:

1400°C/10 h + 1200°C/50 h + 900°C/100 h +

600°C/150 h, water quenched from 1200 -

600°C; LOM, XRD, EMPA, and solidus

temperature measurements

[1992Pav1,

1992Pav2]

composition-temperature section at 6

mass% Al for 25 - 35 mass% Ti (from

Ti-55.4Nb-15.5Al to Ti-40Nb-14Al (at.%)

from solidus at ~2000°C down to 600°C


1400°C/10 h + 1200°C/50 h + 900°C/100 h +

600°C/150 h, water quenched from 1200 -

600°C; LOM, XRD, EMPA

[1991Zak]

partial isothermal sections at 1100, 900 and

800°C for the range of 0 - 20 mass% (0 to

~46 at.%) Al and from 25 to 40 mass% Ti

(~39 to 56 at.%)


(1400°C/5 h + 1300°C/30 h + 1100°C/100 h)

+ (1100°C/2 h + 900°C/300 h + 800°C/500 h),

water quenched from 1100 - 800°C; LOM,

XRD, EMPA

[1991Smi,

1992Smi,

1992Zak]

structure of Ti-60 mass% Nb alloy with 1 -

8 at.% Al (up to ~40 at.% Nb and ~12 at.%

Al) in the temperature range 1150 - 400°C;

ordering of the bcc phase and precipitation

of orthorhombic O phase

TEM, XRD of alloys quenched from 1150°C,

aged at 400 - 900°C

[1992Voz]

structure of Ti-60Nb-8Al (mass%) alloy TEM, XRD [1992Tre]

boundary of the and phases at 1650,

1200 and 1000°C in the Nb corner; the

phase ordering

5 alloys arc melted, annealed at 1650, 1200

and 1000°C for 50 h, 14 d and 30 d,

respectively; LOM, XRD, SEM, TEM, EMPA

[1992Men,

1996Men]

phase relations in Ti-Nb-15Al alloys up to

40 at.% Ti in the temperature range

> 800°C; site occupancy in the ordered

(B2) phase

plasma arc melted alloys; TEM, EMPA and

ALCHEMI

[1994Hou]

ordering and phase transformations in the

Ti3Al based alloy with ~5 at.% Nb

extruded at ~1232°C, annealed in the and

+ fields and quenched, then annealed at 700

- 1000°C and again quenched; XRD, TEM

[1977Sas]

phase transformations alloys Ti-(10-20)Nb-25Al on quenching and

low-temperature aging at 400 - 500°C; TEM

[1982Str,

1988Str]

and 2 phase boundaries up to 7.5 at.%

Nb at 1000°C

Nb containing Ti-(34-38) mass% Al alloys

aged at 1000°C for 605 ks; LOM, XRD,

EMPA

[1988Has]

phase equilibria in the region around

Ti2NbAl (transformations of the bcc phase

to the B2 and Ti4NbAl3 phases)

5 alloys, arc melted, homogenized at 1400°C

for 3 h, annealed at 1100°C for 4 d; LOM,

TEM, SAD

[1989Ben]

phase transformations (ordering of the bcc

phase, the O phase formation)

Ti-(0-30)Nb-25Al alloys both bulk and melt

spun ribbons heat treated at 700 -800°C;

LOM, XRD, SEM, TEM

[1989Kes]

compositions of the 0 and 2 phases in the

Ti-11Nb-24Al alloy in the temperature

range of 1200 - 1020°C

analytical electron microscopy technique [1989Mur]

phase transformations to Ti4NbAl3 1400 - 700°C; LOM, TEM, SEM [1990Ben1,

1990Ben2]


356


MSIT®

Al–Nb–Ti

ordering, structure of ordered phases

( 0, O)

Ti-10Nb-25Al; Ti-12.5Nb-25Al; channeling

enhanced microanalysis, convergent beam

electron diffraction (CBED)

[1987Ban,

1988Ban]

effect of heat treatment on microstructure Ti-11Nb-24Al hot-rolled sheets annealed at

1000 and 1200°C, WQ or furnace cooled (FC);

LOM, XRD, Vicker’s hardness measurements

[1990Kno]

influence of cooling rate on microstructure

and creep properties

Ti-11Nb-25Al, solution treated at 1150°C for

45 min and cooled with rates from 0.02 K s-1

to 10 K s-1 or aged at 750°C for 24 h; LOM,

SEM, creep testing

[1990Mis]

continuous cooling transformations Ti-11Nb-24Al, wedge-shaped specimens,

heated at 1230°C for 1 h, cooled in ice water;

Ti-20Nb-24Al annealed at 1250°C for 8 h, air

cooled; LOM, SEM, TEM, hardness

measurements

[1990Wey]

microstructure and compositions of the

phases ( / 0, 2, O)

Ti-12.5Nb-25Al; extruded and heat treated at

1040°C for 1 h, aged at 760°C for 1 h, creep

tested at 650°C; TEM, EMPA

[1991Akk]

behavior of ordering transformation Ti-21Nb-14Al (mass%), arc melted, forged,

rolled, annealed at 1060°C for 0.5 h, WQ or air

cooled, aged at 700°C for 1 h; TEM, XRD,

SAD

[1991Li]

partial composition-temperature section at

50 at.% Ti

arc melted alloys Ti2NbAl and Ti4Nb3Al,

homogenized at 1400°C, annealed at 700°C

for 26 d; TEM, LOM, SEM

[1991Ben]

microstructure of the Ti-20Nb-3Al alloy arc melted, homogenized at 1400°C, heat

treated in the range 1100 - 700°C; TEM, LOM

[1992Ben]

study of microstructure and evolution of

phases; reaction sequence during

isothermal aging at 650 and 850°C

including a new transition T phase; crystal

structure of the T phase and structural

relationships between T and O phases

Ti-11Nb-24Al; TEM, SAD, microdiffraction

(MD); CBED

[1992Hsi1,

1992Hsi4]

[1992Hsi2,

1992Hsi3]

transformations during aging at 450 -

850°C involving transition metastable

phases other than in [1992Hsi1, 1992Hsi4]

plasma-sprayed alloy Ti-11Nb-24Al, TEM,

XRD

[1995Hsi]

phase transformations from the to O

phase: O phase exists in two forms

Ti-15Nb-24Al alloy, various heat treatments

in the temperature range from 1200 to 650°C;

LOM, TEM, and EMPA

[1992Mur1,

1992Mur2]

phase transformations in the temperature

range from 900 to 400°C

Ti-13Nb-28.5Al alloy; TEM [1993Mur]

vertical section Ti-27.5Al up to 25 at.% Nb

from 1200 to 700°C; refined version of the

section at 27.5 at.% Al with both forms of

the O phase; formation of the O phase by

the peritectoid reaction 2 + 0 O

six Ti-(12.5-25)Nb-27.5Al alloys, arc melted,

heat treated in the range 1200 to 700°C, water

cooled; EMPA, TEM and SAD

[1995Mur1]

transformation from 2 to O at isothermal

aging at 900°C from 15 min to 200 h

alloy Ti-13Nb-28.5Al (at.%); TEM, SAD,

CBED

[1995Mur2]


357


MSIT®

Al–Nb–Ti

transformation temperatures from 600 to

1300°C

Ti-11Nb-25Al alloy; calorimetric differential

thermal analysis (CDTA); in situ high

temperature XRD

[1995Lon]

phase equilibria near Ti3Al at 1000°C four sintered alloys, equilibrated for 225 h

LOM, XRD and EMPA

[1994Kum]

possible transformation paths from high

temperature bcc/B2 to low temperature hcp

or O phase fields were predicted

alloys in the Ti3Al-Nb3Al section [1994Ben1]

schematic pseudobinary Ti3Al-Nb3Al

section up to ~35 at.% Nb

TEM study of three alloys in the Ti3Al-Nb3Al

section annealed at 1100 and 700°C

[1994Ben2]

the structure of alloys in the vicinity of

TiAl, identification of the L60 structure

Ti-20Nb-43Al at 1200°C; TEM, electron

diffraction

[1992Jac]


TiAl

alloy Ti-2.14Nb-47.2Al, plasma arc melted,

annealed at 1050°C for 96 h, LOM, XRD

[1992Kim]


TiAl ( + 0 + , 0 + + )

three alloys in a region of

Ti-(~10-20)Nb-(~40-45)Al, 1200°C for 33 h;

SEM, EMPA

[1993Nak1]


TiAl

alloys Ti-10Nb-45Al and Ti-18Nb-48Al,

1200°C/24 h; TEM, SEM, XRD, electron

diffraction

[1993Zha]

vertical sections at 10 at.% Nb, 48 at.% Al; + 0 + [1998Tak]

phase equilibria ( / 2 + / 0 + ) Ti-10Nb-40Al, from 1200°C/24 h to

1000°C/400 h, XRD, TEM, EMPA

[1998Yu]

phase equilibria between , ( 0) and at

1300 and 1250°C up to ~15 at.% Nb

LOM and EMPA [2000Kai]

effects of Nb on the microstructure and

phase constituents 2, and

in alloys (Ti52Al48-xNbx (0 x 6 at.%), arc

melted, hot isostatic pressing at 1200°C for 3

h, annealed at 1200°C for 12 h, aged at 900°C

for 8 h; LOM, SEM, EMPA, XRD, X-ray

photoelectron spectroscopy (XPS)

[2001Sun]

high temperature phase equilibria; liquidus

projection, isothermal section at 1200°C;

invariant reactions of [1990Per]

confirmed; no ternary phases at 1200°C

alloys inductively melted, homogenized at

1300°C for 20 h and annealed at 1200°C for

two weeks; SEM, XRD, EMPA and TEM;

melting temperatures measured with a

pyrometer

[1995Zdz]

isothermal sections at 1200 and 1000°C;

neither TiNbAl3 ( 1) nor T1 [1989Jew] or

T2 [1989Jew, 1990Per] were found; no

other ternary compounds;

separate area of the ordered B2 phase

detected at 1000°C; considerable solid

solubilities of the third element in most of

the binary phases

diffusion couples and bulk samples annealed

at 1200°C for 48 h, at 1000°C for 96 h and

water quenched; LOM, EMPA, TEM and

XRD

[1998Hel]

phase relations involving Ti4NbAl3 alloy Ti-20Nb-30Al, 1100°C/24 h, 900 -

700°C up to 18 d, TEM, SAD

[1992Ben]


358


MSIT®

Al–Nb–Ti

phase relations involving Ti4NbAl3 alloy Ti-10.8Nb-36.9Al, continuous cooling

from 1200°C at various rates, optical

pyrometry, DTA, dilatometry,

thermoresistometry, differential

microcalorimetry, electron microscopy, TEM,

Vicker’s hardness

[1996Sad]

phase relations involving Ti4NbAl3 alloy Ti-12.9Nb-36.5Al, arc melted, annealed

at 1300°C/6 h, SEM, in situ neutron diffraction

25 - 960°C

[2000Sad]

liquidus projection by [1995Zdz] was

proposed to be changed with respect to the

wider field of primary crystallization

without changing the nature and direction

of the liquid phase reactions; the solid state

transformations were considered ( +

, massive transformation, +

eutectoid-like transformation)

15 alloys in the range 15 to 40 at.% Al with

Nb:Ti ratios of 4:1. 2:1, 1.5:1 and 1:1.5; alloys

were arc melted, homogenized through hot

isostatic pressing (HIP) at 1425 and 1475°C

for times up to 7 h at 138 MPa followed by

water or oil quenching; LOM, XRD, DTA,

BSEI, EMPA, TEM and microhardness

measurements; data on phase equilibria in the

same 15 alloys annealed at 1100°C for 720 h

examined by BSEI, optical microscopy, XRD

and EMPA

[2000Leo2]

phase transformations alloys Ti-21.8Nb-27.9Al and

Ti-31.7Nb-23.4Al prepared by vacuum arc

melting, homogenized at 1300°C during 1

week and annealed at 1260°C for 20 and 70 h,

at 1100°C 20 and 75 h, at 900°C 140 and 1000

h, at 700°C 1500 h with ice-water quenching

after each heat treatment; microhardness

measurements, dilatometry, DTA, XRD,

SEM, TEM; continuous cooling of the alloys

from 1260°C with rates from 100 to 0.25 K·s-1

[2001Sad]

phase evolutions from B2 to the O phase alloy Ti-27Nb-23Al rolled sheet, 650 -

1090°C, up to 450 h; EMPA, SEM, TEM,

XRD, DTA

[1999Boe]

effect of cooling rate on the

transformations from B2 to the O phase

alloys Ti-37.5Nb-25Al, Ti-35Nb-30Al,

Ti-~44.5Nb-~25.6Al; DTA up to 1500°C,

SEM, TEM, electrical resistivity

measurements

[2001Mis]

phase evolutions from B2 to the O phase three alloys around Ti2NbAl, arc melted,

annealed at 1200°C for 3 h, aged at 600 -

900°C for 0.5 to 300 h; XRD, LOM, TEM,

SEM

[2001Zha]

phase evolutions from B2 to the O phase three alloys around Ti2NbAl, annealed at

1350, 900, 800 and 700°C up to 1500 h and

quenched; XRD, SEM, TEM

[2003Sad]


359


MSIT®

Al–Nb–Ti


Phase /

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments /References

(Al)

< 664.2

< 660.452

cF4

Fm3m

Cu a = 404.96

0 to 0.6 at.% Ti [1992Kat2]

[V-C2]

pure Al at 25°C [1981Kin]

, (Ti1-x-yNbxAly)

(Nb)

< 2469

( Ti)h

1670 - 882

cI2

Im3m

W

a = 330.07

a = 327.6 0.3

a = 330.65

a = 328

a = 327 5

a = 327.3 to 337.5

a = 327.3 to 328.5

a = 326

a = 330.3

0 x 1 >882°C at y = 0 [Mas2, 1987Mur]

0 y 0.448 at x = 0 [1993Gam, 1993Oka1,

2000Oka, 2003Sch]

0 y 0.46 at x = 0 [1996Tre, 1997Bul]

0 y 0.215 at x + y = 1 [Mas2]

pure Nb at 25°C [1981Kin]

for Nb-21.5 at.% Al [1980Jor]

[Mas2]

for Ti-45Nb-10Al in Ti-37.2Nb-12.2Al (at.%)

alloy annealed at 700°C for 26 d, [1991Ben]

for Ti-60.3Nb-10.8 Al alloy homogenized at

1300°C for 20 h [1999Cha1], 900°C

[1984Zak]

[1983Tro]

Ti-47Nb-6.3Al [1992Pav2]

Ti-40.8Nb-17.4Al ( 2+ + ) 900°C

[1992Pav1]

* 0 cP2

Pm3m

CsCl

a = 323.5

a = 324 3

a = 326 3

a = 326.6

a = 325.1

a = 326.9

a = 327.1

a = 327.5

a = 326.8

a = 324.4

a = 326.1

a = 323.05 0.05

a = 322.50 0.05

a = 325

a = 330

a = 328

ordered form of the high temperature

(Ti,Nb,Al) solid solutions [1989Ben,

1989Kes, 1991Ben, 1991Cha, 1992Voz,

1994Hou, 1995Mur1, 1996Men, 1996Vas,

1998Hel, 1998Rho, 1999Cha2, 1999Boe,

1999Flo, 1999Rav, 2000Leo1, 2001Sad,

2002Leo2, 2003Sad]

Ti-25Nb-25Al after rapid quench. [1991Cha]

for as cast Ti-20.6Nb-26.7Al [1999Cha2]

for Ti-54.3Nb-15.4Al, homogenized at 1300°C

for 20 h [1999Cha2]

Ti-42.5Nb-15Al, as cast [2000Leo2]


Ti-51Nb-15Al, as cast [2000Leo2]


Ti-68Nb-15Al, as cast [2000Leo2]

Ti-40.90Nb-15.44Al, annealed at 1100°C for

720 h [2002Leo1]

Ti-26.8Nb-21.8Al, annealed at 1350°C

[2003Sad]

Ti-30.2Nb-19.7Al, annealed at 1350°C

[2003Sad]

Ti-14.4Nb-30.1Al (1200°C) [1998Hel]

Ti-16.8Nb-34.6Al (1200°C) [1998Hel]

Ti-25.4Nb-25.1Al (1000°C) [1998Hel]

Ti-45Nb-25Al, as cast [1995Zdz]

Ti-45Nb-10Al [1991Ben]

and Ti-11Nb-25Al [1995Lon]

360


MSIT®

Al–Nb–Ti

( Ti1-x-yNbxAly)

< 1490

( Ti)r

< 882

hP2

P63/mmc

Mg

a = 295.03

c = 468.36

a = 291

c = 469

at x = 0 47.3 to 51.4 at.% Al at solidus

temperatures 1490 to 1462°C [1993Oka1,

1993Oka2, 2000Oka, 1993Gam, 2003Sch]

at x = 0 from ~48 at.% Al at 1520°C to 51 at.%

Al at 1485°C [1996Tre, 1997Bul]

dissolves up to 10 at.% Nb at 1200°C

[1998Hel]

pure Ti at 25°C [Mas2, V-C2, 1981Kin]

dissolves up to ~2 at.% Nb in the Nb-Ti system

[Mas2]

Ti-5Nb-40Al annealed at 1400°C for 6 h, WQ

[1996Che]

(Ti1-xNbx)3Al, 2

Ti3Al

< 1164

hP8

P63/mmc

Ni3Sn

a = 580.6

c = 465.5

a = 574.6

c = 462.4

a = 576.7 0.4

c = 465.4 0.7

a = 580 10

c = 480 10

a = 580 10

c = 460 10

a = 580

c = 466

a = 574.3

c = 498.4

a = 572.4 to 574.3

c = 498.4

~20 to 38.2 at.% Al

D019 ordered phase (“ 2Ti3Al”);

maximum at 30.9 at.% Al [1992Kat2, 1993Oka1,

1993Oka2]

< 1180°C [1993Gam]

maximum at 32.5 at.% Al and 1200°C

[1996Tre, 1997Bul]

at 22 at.% Al [L-B]

at 38 at.% Al [L-B] [V-C]

Ti-12.4Nb-30.9Al (1000°C) [1998Hel]

Ti-11Nb-24Al [1990Wey]

in thin films Ti-11Nb-24Al [1992Hsi1]

in alloy Ti-11Nb-25Al [1995Lon]

Ti-13.8Nb-13.4Al (mass%) [1984Zak]

[1983Tro]

Phase /

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]


361


MSIT®

Al–Nb–Ti

TiAl,

< 1463

tP4

P4/mmm

AuCu

a = 400.5

c = 407.0

a = 400.0 0.1

c = 407.5 0.1

a = 398.4 0.1

c = 406.0 0.1

a = 399

c = 408

a = 399.4

c = 409.6

a = 399.3

c = 410.4

a = 397.9

c = 412.6

a = 398

c = 419

a = 399

c = 407

L10 ordered phase (“ TiAl”)

46.7 to 66.5 at.% Al [1992Kat2, 1993Oka1]

~52 to 65 at.% Al at solidus temperatures,

~50 to 60 at.% Al at 1000°C [1996Tre,

1997Bul]

50 to 62 at.% Al at 1200°C [2001Bra]

at 50 at.% Al [2001Bra]

at 62 at.% Al [2001Bra]

at 55.4 and 61.8 at.% Al [1998Hel]

(Ti0.70Nb0.30)Al [1991Smi, 1992Zak]

Ti-19Nb-53Al ( in a + alloy) [1997Jew]

Ti-18.9Nb-55.6Al (1200°C) [1998Hel]

Ti-10Nb-50Al, annealed at 1400°C for 6 h,

WQ [1996Che]


WQ [1996Che]

Phase /

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]


362


MSIT®

Al–Nb–Ti

TiAl2,

< 1199

oC12

Cmmm

ZrGa2

tP4

P4/mmm

AuCu

tI24

I41/amd

HfGa2

tP32

P4/mbm

Ti3Al5

a = 1208.84

b = 394.61

c = 402.95

a = 403.0

c = 395.5

a = 397.0

c = 2430.9

a = 396.7

c = 2429.68

a = 397

c = 2430

a = 394.89

c = 412.36

a = 397.16

c = 405.92

a = 1129.3

c = 403.8

chosen stoichiometry [1992Kat2] summarizing

several phases [2003Sch]:

metastable modification of TiAl2, only

observed in as-cast alloys [2001Bra];

listed as TiAl2(h) (66 to 67 at.% Al,

1433-1214°C) by [1990Sch]

Ti1-xAl1+x; 63 to 65 at.% Al at 1300°C, stable

range 1445 - 1170°C [2001Bra];

listed as orthorhombic, Pmmm, with

pseudotetragonal cell by [1990Sch]

(range ~1445 - 1424°C)

for Ti36Al64 at 1300°C [2001Bra]

stable structure of TiAl2 <1216°C [2001Bra];

66 to 67 at.% Al at 1000°C [2001Bra];

listed as TiAl2(r) by [1990Sch];

< 1210°C [1996Tre, 1997Bul]

[2001Bra]

[1990Sch]

stable between 65.5 and 66.9 at.% Al, dissolves

~5 at.% Nb at 1200°C [1998Hel]

Ti-4.1Nb-64.6Al ( in a + alloy at 1200°C)

(for the CuAu type subcell) [1998Hel]

Ti-3.8Nb-65.8Al ( in a + + alloy at

1000°C (for the CuAu type subcell) [1998Hel]


Phase /

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]


363


MSIT®

Al–Nb–Ti

tetragonal

superstructure of

AuCu type

[2001Bra]

tI16

I4/mmm

ZrAl3

tP28

P4/mmm

Ti2Al5

a* = 395.3

c* = 410.4

a* = 391.8

c* = 415.4

a=398.81-392.3

c=1649.69-1653.49

a = 393

c = 1654

a = 390.53

c = 2919.63

summarizes several phases [2003Sch]:

Ti5Al11

stable range 1416 - 995°C [2001Bra]

66 to 71 at.% Al at 1300°C [2001Bra]

(including the stoichiometry Ti2Al5);

at 66 at.% Al [2001Bra]

* AuCu subcell only

at 71 at.% Al [2001Bra]

* AuCu subcell only

D023 type [V-C]

68.5 to 70.9 at.% Al, 1416 - 1206°C [1990Sch]

69-71 at.%Al, 1450-990°C [1996Tre,

1997Bul]

for 69.4 at.% Al, accepted as Ti2Al5, stable

between 69.4 and 71.8 at.% Al at 1200°C,

dissolves ~2 at.% Nb [1998Hel]

“Ti2Al5”; 1416 - 990°C [1992Kat2]

~1215 - 985°C [1990Sch];

included in hom. region of Ti5Al11 [2001Bra]

(Ti1-xNbx)Al3,

TiAl3 (h)

< 1393

NbAl3 < 1680

tI8

I4/mmm

TiAl3

a = 384.9

c = 860.9

a = 385.3

c = 858.7

a = 384.1

c = 860.9

a = 385.2

c = 859.9

a = 384.6

c = 862.0

a = 384

c = 865

a = 384.6

c = 860.9

a = 384.2

c = 861.6

a = 385.9

c = 857.6

a = 385.6

c = 858.6

a = 386.2

c = 859.0

0 x 1 [1989Jew, 1989Kal, 1990Per,

1992Pav2, 1995Zdz, 1996Che, 1998Din,

1998Hel, 1998Wan]

74.2 to 75.0 at.% Al [2003Sch]

D022 ordered phase, 1387 - 735°C,

74.5 to 75 at.% Al at 1200°C [2001Bra]

homogeneity range 74.4 - 75.3 at.% Al

[1998Hel], 74.3 - 75.6 at.% Al [Mas2]

[2003Kar]

[1980Jor]

Ti-7.4Nb-73.6Al (in a + alloy) [1998Hel]

Ti-12.0Nb-74.9Al [1998Hel]


WQ [1996Che]

Ti-18.8Nb-74.6Al [1998Hel]

Ti-19.0Nb-75.1Al [1998Hel]

Ti-21.2Nb-72.4Al (in a + alloy) [1998Hel]

Ti-23.2Nb-72.5Al (in a + + alloy)

[1998Hel]

Ti12Nb16Al72 [1991Spa] in alloy

Ti-12Nb-63Al

Phase /

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]


364


MSIT®

Al–Nb–Ti

TiAl3 (l)

< 950 (Ti rich)

tI32

I4/mmm

TiAl3 (l)

a = 387.7

c = 3382.8

74.5 to 75 at.% Al [2001Bra]

TiAl3 cP4

Pm3m

AuCu3

a = 397.2 metastable, obtained at 85 at.% Al from splat

cooling [2001Bra]

Nb2Al,

< 1940

tP30

P42/mnm

CrFe a = 995.2 to 986.6

c = 516.8 to 518.7

a = 990.1

c = 517.0

a = 992.64

c = 515.54

a = 991.5

c = 517.3

32 to 42 at.% Al at solidus temperatures, 32 to

35 at.% Al at 1300°C [Mas2, V-C]

[1980Jor]

Ti-46.9Nb-41.2Al (in a + + alloy,

1200°C) [1998Hel]

Ti-48.9Nb-36.2Al (1000°C) [1998Hel]

Ti-54Nb-35.9Al (900°C) [1992Pav1]

Nb3Al,

< 2060

cP8

Pm3n

Cr3Si

a = 519.7 to 518.0

a = 516.97

a = 517.26

a = 519.0

18.6 to 25 at.% Al [Mas2]

at 19 to 25 at.% Al [1980Jor]

Ti-45.2Nb-24.1Al (in a + alloy, 1000°C)

[1998Hel]

Ti-52.7Nb-22.2Al (in a + + alloy, 1200°C)

[1998Hel]

in the Ti-65.8Nb-26.7Al ( + ) alloy annealed

at 900°C for 100 h

Phase /

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]


365


MSIT®

Al–Nb–Ti

* Ti2NbAl, O

980

oC16

Cmcm

NaHg

a = 608.93 0.02

b = 956.94 0.04

c = 466.66 0.02

a = 605

b = 961

c = 465

a = 610.6 0.3

b = 955.7 0.3

c = 463.1 0.6

a = 609.5 0.3

b = 956.9 0.3

c = 466.0 0.6

a = 610 10

b = 980 10

c = 470 10

a = 608

b = 962

c = 466

a = 612

b = 956

c = 466

a = 615

b = 953

c = 466

a = 596

b = 986

c = 467

a = 604

b = 971

c = 464

[1988Ban, 1990Moz]

exists in two forms, O1(h) (~980 to 900°C) and

O2 (r) (below ~900°C) with different site

occupancies [1992Mur2, 1995Mur1,

1995Mur2, 2002Wu]

Ti-25Nb-25Al, annealed at 700°C for 228 h

(neutron powder diffraction, Rietveld

refinement) [1990Moz];

ordered, distorted Ni3Sn type

for Ti2NbAl in the Ti-37.2Nb-12.2Al alloy

annealed at 700°C, 26 d [1991Ben]

in the Ti-40.1Nb-18.4Al alloy, 1150°C, WQ +

700°C/5 h [1992Tre]

in the same alloy, 1150°C, WQ + 700°C/5 h +

400°C/5 h [1992Tre]

in Ti-20Nb-25Al, cooled from field

[1990Wey] and in thin films Ti-11Nb-24Al

[1992Hsi4]

in Ti-20Nb-25Al, 800°C [1989Kes]

in Ti-25Nb-25Al, 800°C [1989Kes]

in Ti-30Nb-25Al, 800°C [1989Kes]

for O1 and O2 in Ti-15Nb-24Al [1990Mur]

in Ti-11Nb-25Al [1995Lon]

Phase /

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]


366


MSIT®

Al–Nb–Ti

* Ti4NbAl3,

900

hP6

P63/mmc

Ni2In

a = 458.0 0.3

c = 552.0 0.4

a = 455.5

c = 554.2

a = 457.6

c = 552.4

[1990Ben1, 1990Ben2, 1992Ben, 2000Sad]

in the Ti-12.5Nb-37.5Al alloy aged at 700°C

for 26 d (Ti4NbAl3+ 0+( 2)) [1990Ben1]

in situ neutron diffraction at RT for the

Ti-12.9Nb-36.5Al alloy ( 0+ 2+Ti4NbAl3)

[2000Sad]

the same alloy at 805°C [2000Sad]

´ hP2

P63/mmc

Mg a = 580 10

c = 480 10

disordered martensite phase in Ti-Nb-25Al at

Nb content < 5 at.% [1988Str]

[1990Wey]

´´ oP4

P2221

a = 296.5

b = 492.8

c = 464.6

metastable phase, in rapidly solidified

Ti3Al-Nb alloys containing < 2 at.% Nb

[1995Xu] [1990Wey]

T hP3

P6/mmm

TiCr

metastable phase, in Ti-5Nb-25Al, aged at 350

- 550°C [1978Zak]

[1990Ben1, 1991Li, 1992Sur, 1992Voz,

1996Men, 2000Leo2, 2000Sad, 2001Sad]

´´ trigonal

P3ml

a = 455.54 0.10

c = 554.15 0.14

a = 460

c = 580

a = 457.5

c = 560.4

metastable phase [1990Ben1, 1990Sho,

1992Sur, 1994Che1, 2000Leo2, 2000Sad,

2001Sad]; T - the idealized version of the same

phase [1990Ben1]

Ti3Nb0.75Al2.25 at 23°C [1990Sho]

[1989Ben] Ti4NbAl3

[1991Cha] Ti-25Nb-25Al

Phase /

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]


367


MSIT®

Al–Nb–Ti

10 20

1200

1300

1400

1500

1600

1700

1800

Ti 25.00Nb 0.00Al 75.00

Ti 0.00Nb 25.00Al 75.00Nb, at.%

Te

mp

era

ture

, °C

1680+-5°C

L

ε

L+ε

L+ζ+ε

ζ+L

1393°C

Fig. 1: Al-Nb-Ti.

The TiAl3-NbAl3section

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb


Axes: at.%

ζ

α

β

γ

δ

σ

ε

Fig. 2: Al-Nb-Ti.

Liquidus surface

projection [1995Zdz].

Dotted lines is the

limit of primary

crystallization field

found by [2000Leo1]

368


MSIT®

Al–Nb–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb


Axes: at.%

L

α

β

δ

σ

β0

L+σ

σ+δβ+σ

β+δ

L+α

β+β0

β0+σ

α+β

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb


Axes: at.%

ε

γ

α

β

σ

δ

ζ

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb


Axes: at.%

L

α

β

δ

σ

β0

L+σ

σ+δβ+σ

β+δ

L+α

β+β0

β0+σ

α+β

Fig. 3: Al-Nb-Ti.

Calculated liquidus

surface projection

[1998Ser]

Fig. 4: Al-Nb-Ti.

Calculated isothermal

section at 1650°C

[1998Ser]

369


MSIT®

Al–Nb–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb


Axes: at.%L

γ

α

β

ε

σ

δ

ζ

ε+γ

γ+σ

β0

β+σ σ+δ

ε+σα+γ

β+δ

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb


Axes: at.%L

α

β

ζε

σ

δ

γγ

1

ε+σ

σ+δ

β+δβ+σ

γ+σβ+α

α+γ

β+γ

ε+γ

L+ζ+ε

Fig. 6: Al-Nb-Ti.


section at 1400°C

[1998Ser]

Fig. 5: Al-Nb-Ti.

Experimental


1400°C [1996Che,

1998Wan]

370


MSIT®

Al–Nb–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb


Axes: at.%L

ε

ζ

γ

α σ

δ

β

γ+ε

ε+σ

σ+δ

γ+σ

β0+σ

α+γ

L+ε

β0

β+δ

50

60

10 20

40

50

Ti 70.00Nb 0.00Al 30.00

Ti 40.00Nb 30.00Al 30.00

Ti 40.00Nb 0.00Al 60.00 Data / Grid: at.%

Axes: at.%

α+β

γ+β

γ+α

γ

α

β

Fig. 7: Al-Nb-Ti.


section at 1300°C

[2001Sad]

Fig. 8: Al-Nb-Ti.

Partial isothermal

section at 1300°C

[2000Kai]

371


MSIT®

Al–Nb–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb


Axes: at.%

εζ

η

γ

α σ

δ

β

ε+σ

σ+δ

β+δ

γ+σα+γ

β+σ

L

L+ε

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb


Axes: at.%L

ε

ζ

γ

α σ

δ

β

β0

ε+γ

ε+σ

σ+δ

β+δ

γ+σ

β0+σ

α+γ

L+ε

Fig. 9: Al-Nb-Ti.

Experimental


1200°C [1995Zdz]

Fig. 10: Al-Nb-Ti.


section at 1200°C

[1998Ser]

372


MSIT®

Al–Nb–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb


Axes: at.%

εζη

γ

α

σ

δ

β

α2

α

γ+ε

γ+α

β0

β0+σ

ε+σ

β+δ

σ+δ

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb


Axes: at.%

εζη

γ

α2

σ

δ

β

α

ε+σ

σ+δ

β+δ

β0

β0+σ

α2+γ γ+σ

ε+γ

Fig. 11: Al-Nb-Ti.


section at 1100°C

[1998Ser]

Fig. 12: Al-Nb-Ti.


section at 1100°C

[2001Sad]

373


MSIT®

Al–Nb–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb


Axes: at.%

ε

β0

η

γ

α

σ

δ

β

α2

ζ

ε+σ

δ+σ

β+δ

β0

γ+ε

γ+σα2+γ

β0+σ

α2+β

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb


Axes: at.%L

β0

ε

α2

η

γ

α

σ

δ

β

ζ

β+δ

σ+δ

ε+σγ+σ

ε+γ

γ+α2

α2+β

α2+σ

Fig. 13: Al-Nb-Ti.


section at 1020°C

[1998Ser]

Fig. 14: Al-Nb-Ti.

Experimental


1000°C mainly based

on [1998Hel]

374


MSIT®

Al–Nb–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb


Axes: at.%

εη

α

σ

δ

β/β0

α2

γε+σ

σ+δ

γ+σα2+γ

γ+ε

O1

β0

Fig. 15: Al-Nb-Ti.


section at 990°C

[2001Ser]

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb


Axes: at.%

ε

η

γ

α

σ

δ

β

α2 σ+δ

β+δ

ε+σγ+σ

ε+γ

ε+η

τ

O

ε+γ+σ

Fig. 16: Al-Nb-Ti.


section at 900°C

[1998Ser]

375


MSIT®

Al–Nb–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb


Axes: at.%

ε

η

γ

σ

δ

β

α2

α

ε+σ

σ+δ

β+δ

τ

γ+ε

γ+ε+σ

O2

α2+β O2+β

γ+τ γ+σ

O2+σ

η+ε

Fig. 17: Al-Nb-Ti.


section at 800°C

[1998Ser]

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb


Axes: at.%

ε

τ

η

γ

α

σ

δ

β

α2

O2

γ+ε

ε+σ

σ+δ

δ+βO2+β

α+β

γ+τ

γ+ε+σ

α2+O2

O2+σ

τ+γ+σ

Fig. 18: Al-Nb-Ti.


section at 700°C

[2001Ser]

376


MSIT®

Al–Nb–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb


Axes: at.%

ε

σ

δ

β

O2

τ+O2+σ

O2+β

τ+σ+ε σ+ε

γ+τ+εγ+ε

γ+η+ε

O2+σσ+δO2+δ+σ

O2+β+δ

Fig. 19: Al-Nb-Ti.

Calculated partial


600°C [1998Ser]

10 20 30

700

800

900

1000

1100

1200

Ti 72.50Nb 0.00Al 27.50

Ti 37.50Nb 35.00Al 27.50Nb, at.%

Te

mp

era

ture

, °C

α β0

α2 O1

O2

Fig. 20: Al-Nb-Ti.

The partial isopleth

along 27.5 at.% Al

377


MSIT®

Al–Nb–Ti

10 20 30

600

700

800

900

1000

1100

1200

1300

Ti 72.50Nb 0.00Al 27.50

Ti 37.50Nb 35.00Al 27.50Nb, at.%

Te

mp

era

ture

, °C

β0

β0+σ

σ+O1

σ+O2

O2+τα2+O2+τ

α2

α2+O2

α2+O1

α2+β0+σ

α+β0

β+αβ

α

O2+τ+σ

α2+σ

α+α2

α2+O2+σ

α2+O1+σ

β0+σ+O1α2+β0

β+α2

Fig. 21: Al-Nb-Ti.

The calculated partial

isopleth along

27.5 at.% Al

[1998Ser]

0.15

0.1

0.05

0

0 0.05 0.1 0.15 0.2

Ti

Al

α

x

Fig. 22: Al-Nb-Ti.

Thermodynamic

activities of Ti and Al

in the alloys

(Ti0.38Al0.62)1-xNbx

at 1200°C [1996Eck]

378


MSIT®

Al–Nb–Ti

-150

-130

-110

-90

-70

-50

-30

-10

10

30

50

0 0.05 0.1 0.15 0.2

∆ mix

-1H

,kJ·m

ol

x

Ti

Al

Fig. 23: Al-Nb-Ti.

Partial enthalpies of

mixing of Ti and Al in

the alloys


at 1200°C [1999Eck]

-70

-50

-30

-10

10

30

50

0 0.05 0.1 0.15 0.2

∆ mix

-1-1

S,

J·m

ol

K·

x

Ti

Al

Fig. 24: Al-Nb-Ti.

Partial entropies of

mixing of Ti and Al in

the alloys


at 1200°C [1999Eck]

379


MSIT®

Al–Nb–Ti

10-1 100 101 102 103

Time, s

Te

mp

era

ture

,°C

200

400

600

800

1000

1200

Thermodynamic calculation1378

1107

992

915

845

�

�0+�

�0

�0+ +O1�

�2+ +O2

�0��

�0 2�

� 0 2 O+ ��0�Om

�0�O

�0�� 0�

��

100K/s50K/s 20K/s

10K/s

5K/s

3K/s2K/s 1K/s

0.5K/s

0.25K/s

15K/s

�+ 2+O2

Fig. 25: Al-Nb-Ti.

CCT diagram for the

Ti-27.9Al-21.8Nb

alloy (cooling rate

in K s-1) [2001Sad]

10-1 100 101 102 103

Time, s

Te

mp

era

ture

,°C

200

400

600

800

1000

1200

Thermodynamic calculation

12541217

1017985

940915895

�� +

�0+�

�0+ +O1��+O1O1

O1+O2

O 2

�0��

�0 2�

� 0 2 O+ �

�0�Om

�0�O

�0��

��

100K/s

50K/s

20K/s10K/s

5K/s

4K/s

3K/s

2K/s

1K/s

0.5K/s

0.25K/s

30K/s

Fig. 26: Al-Nb-Ti.

CCT diagram for the

Ti-23.4Al-31.7Nb

alloy [2001Sad]

380


MSIT®

Al–Nb–Zr

Aluminium – Niobium – Zirconium

Lazar Rokhlin, Natalia Bochvar

Literature Data

In his investigations along the section ZrAl3-NbAl3 [1962Poe] found that at 660°C a significant amount

(about 12.5 at.%) of Zr can be dissolved in NbAl3.

Nb-rich alloys containing up to 26 at.% Zr and 30 at.% Al were investigated first by [1967Yam] using X-ray

diffraction methods and metallography, later Nb rich alloys were investigated in the same way by

[1974Fed1, 1974Fed2]. From the latter works a partial isothermal section results at 800°C. The work of

[1967Yam] allowed him to construct partial isothermal sections of the phase diagram at 1300 and 1100°C.

Detailed microscopy studies of Zr-rich quenched samples allowed [1968Tre] to construct partial isothermal

sections at 1000, 900, 800 and 700°C for the Zr corner of the phase diagram. The description presented later

by [1970Ali] is merely based of the on [1968Tre].

[1970Han] studied the Al-Nb-Zr system in almost the entire concentration range employing X-ray

diffraction method to determine the crystal structures and lattice constants of the phases involved. The

conclusions of these experiments were presented as isothermal section at 925°C which confirms the

significant solubility of Zr in NbAl3 found earlier by [1962Poe].

In [1977Ale] the solubility of Zr in Nb3Al was established once more by X-ray measurements of lattice

parameters, this time after annealing at 700°C.

Aiming to establish the boundaries between the phase areas more precisely, [1990Per] studied the Zr corner

of the phase diagram once more. These authors employed light metallography, quantitative metallography,

X-ray diffractometry and electron microprobe X-ray analysis on samples with controlled oxygen and

nitrogen contents. The resulting isothermal sections at 800, 771 and 730°C showed the same phase fields,

but compared with [1968Tre] and [1990Per] shifted phase boundaries.

[1989Sub] discussed the phase stability of NbAl3 depending on the solubility of Zr using the solubility data

from [1970Han].

[1993Hub] presented a review on the Al-Nb-Zr phase diagram basing on investigations conducted by

[1970Han, 1968Tre, 1974Fed1, 1974Fed2]. [1993Hub] gave the isothermal section of the phase diagram at

925°C which was constructed according to [1970Han] with addition of the phase areas in Zr corner

according to [1968Tre].

The most recent investigation of the Al-Nb-Zr system were done in the Al corner by [1994Sok] using X-ray

phase analysis and light metallography. The partial isothermal section at 500°C, constructed by [1994Sok],

shows a significantly larger solubility of Nb in ZrAl3 and significantly smaller solubility of Zr in NbAl3 than

the earlier work by [1970Han] does. This fact is difficult to explain even if the lower temperature of

[1994Sok] is taken into account. In this evaluation the data on solid solubilities in the compounds by

[1970Han] are preferred because they were obtained on more exact measurements of the lattice constants.

Binary Systems

The edge binary systems Al-Zr and Al-Nb are accepted as recently evaluated by [2003Sch] and [2003Vel],

respectively. Phase relations in Nb-Zr are accepted as drawn by [1992Oka].

Solid Phases

Two ternary compounds exist in the Al-Nb-Zr system according to [1970Han]. The ternary compound 1

has a homogeneity range limited by 12 to 25 at.% Nb and 46 to 54 at.% Al. The 1 homogeneity field in the

isothermal section at 925°C has the shape of a deformed ellipse. [1970Han] gave for the 1 the formulae

Zr5Nb2Al6-Zr3Nb3Al7. In the assessment of [1993Hub] for the 1 compound the generalized formulae

Zr5-2xNb2+xAl6+x was assumed with 0 x 1.

The composition of 2 is about Zr35Nb30Al35. The crystal structure of 2 has not been described yet.

381


MSIT®

Al–Nb–Zr

Analyzing X-ray diffraction patterns [1970Han] indicates that a third ternary compound with cubic crystal

structure of the CsCl type may exist in the middle part of the Zr3Al-Nb3Al section, at elevated temperatures.

The homogeneity ranges of the binary compounds ZrAl2, Zr4Al3, Zr5Al3 and Zr2Al extend substantially

into the ternary system, up to 19 at.% Nb. ZrAl3 dissolves up to 2 at.% Nb at 925°C. These solubility values

from [1970Han] are accepted here, although [1994Sok] reports that at 500°C as much as 8 at.% Nb would

be dissolved in ZrAl3. Other Al-Zr compounds, ZrAl, Zr3Al2, Zr3Al, dissolve only insignificant amounts of

Nb.

The phase (Nb,Zr)2Al has a large homogeneity range in the binary system and extends substantially into

the ternary system dissolving up to 15 at.% Zr [1970Han]. For (Nb,Zr)3Al investigations at different

temperatures by [1970Han, 1974Fed1, 1977Ale] give a consistent trend for the amount of Zr that can be

dissolved in this phase: about 10 at.% at 925°C, 5 at.% at 800°C and 4 at.% at 700°C. The data reported by

[1967Yam] i.e. 3 at.% Zr dissolved at 1300 and 1100°C has to be taken with care. Details on crystal

structure data of the solid phases are presented in Table 1.

Isothermal Sections

Figure 1 displays the isothermal section of the Al-Nb-Zr phase diagram at 925°C. It is constructed mainly

after [1970Han] with additions of two supposed three-phase fields, Zr3Al2+(Zr,Nb)5Al3+(Zr,Nb)4Al3,

Zr3Al+(Zr,Nb)2Al+( Zr1-x-yNbx-yAly) and (Nb,Zr)Al3+(Nb,Zr)Al2+Nb3Al2 which should exist according

to the phase rule. The boundaries of the miscibility gap in the (Nb,Zr) continuous solid solution shown at

the Nb-Zr side take the miscibility gap in the binary Nb-Zr system into account. The section shows the

compound (Zr,Nb)5Al3(h) at 925°C established firmly by [1970Han]. This does not contradict the binary

Al-Zr phase diagram by [2003Sch] because the lower limit of existence for that phase is shown by

[1970Han] at about 1000°C only tentatively. Unlike in [1970Han] the homogeneity ranges of the Al-Zr

compounds are shown as line-compounds taking into consideration that they are very narrow in the binary

Al-Zr system. Figures 2, 3, 4 display the partial isothermal sections of the Zr corner of the phase diagram

at 800, 771 and 730°C after [1990Per]. The estimated solubility of Al in ( Zr) had to be shifted to meet the

binary Al-Zr after [2003Sch].


Figure 5 shows vertical sections of the surface between ( Zr1-x-yNbx-yAly) solid solution and

( Zr)+( Zr1-x-yNbx-yAly) phase areas. The sections correspond to the constant Al contents of 0, 3.3 and 6.7

at.%. The sections were constructed mainly after [1990Per] with some corrections to be consistent with the

accepted Al-Zr binary phase diagram.


Additions of Al and Nb lower the corrosion resistance of Zr in water at elevated temperatures and high

pressure [1968Tre]. Adding Nb additive has a favorable effect on high temperature hardness and creep

resistance of Zr3Al, as found by [2003Tew].

The superconductivity of the compound Nb3Al with Zr additives was studied in [1975Fed, 1975Sha,

1977Ale]. Addition of Zr to the compound decreased temperature of the superconductivity transition Tc.

Miscellaneous

[1985Zak] studied structural transformations during decomposition of the ( Zr) based solid solution in

Zr-rich alloys containing Al and Nb, and described the sequential formation of a number of metastable

phases.

Similarly [1999Tew] studied the structure transformations in the Zr3Al alloys containing up to 10 mass%

Nb. The alloys were rapidly quenched from liquid state and annealed then. Sequence of the solid phase

formations was established.

382


MSIT®

Al–Nb–Zr

References

[1962Poe] Poetzschke, M., Schubert, K., “On the Constitution of Some Systems Homologous or

Quasihomologous to T4-B3. II. The System Ti-Al, Zr-Al, Hf-Al, Mo-Al and Some Ternary

Systems” (in German), Z. Metallkd., 53(8), 548-561 (1962) (Equi. Diagram, Crys.


[1967Yam] Yamamoto, A.S., “The Determination of the Niobium-Rich Region of the Ternary Phase

Diagram, Niobium-Aluminium-Zirconium. Phase Equilibria of the Niobium - Tungsten -

Hafnium and Niobium - Tungsten - Zirconium Alloy Systems”, U.S. At. Energy Comm.

Publ., 1-82 (1967) (Equi. Diagram, Experimental, Mechan. Prop., 26)

[1968Tre] Tregubov, I.A., Kudryavzev, D.L., “The Zr Corner of the Phase Diagram and Properties of

Alloys of the Zr-Al-Nb System” (in Russian), in “Fiziko-Khimiya Splavov Zirkoniya”,

14-17 (1968) (Equi. Diagram, Experimental, Corrosion, 3)

[1970Ali] Alisova, S.P., Budberg, P.B., “Aluminium-Niobium-Zirconium” (in Russian), Diagrammy

Sostoyaniya Met. Sistem, 14, 133-133a (1970) (Equi. Diagram, Review, 1)

[1970Han] Hansen, R.G., Raman, A., “Alloy Chemistry of ( -U)-Related Phases. III. -Phases with

Non-Transition Elements”, Z. Metallkd., 61, 115-120 (Equi. Diagram, Crys. Structure,


[1974Fed1] Fedorova, M.A., Burnashova, V.V., Turchinskaya, M.I., Sokolovskaya, E.M., “Phase

Composition and Superconductivity in Alloys of the System Nb-Al-Ti {Zr, Hf}” (in

Russian), Moskovskii Universitet, Moscow, 2137-74, (1974) (Experimental, 10) (quoted in

Alisova, S.P., Budberg, P.B., Diagrammy Sostoyaniya Met. Sistem, 20, 133-134 (1974)

(Equi. Diagram, Review, 1)

[1974Fed2] Fedorova, M.A., Burnashova, V.V., Sokolovskaya, E.M., Kripyakevich, P.I., “Ternary

Compounds in (Ti, Zr, Hf)-(Nb, Ta)-Al Systems” (in Russian), Tezisy Dokl. Vses. Konf.

Kristallokhim. Intermetall. Soedin., 2nd, Lvov, 18 (1974) (Crys. Structure, 0)

[1975Fed] Fedorova, M.A., Turchinskaya, M.I., Sokolovskaya, E.M., “Influence of Group IVb

Elements on the Structure and Superconductive Properties of the Intermetallic Compound

Nb3Al”, Phys. Met. Metallogr., 30, 86-87 (1975), translated from Vest. Mosk. Univ., Ser. 2:

Khim., 30, 238-240 (1975) (Experimental, 4)

[1975Sha] Shamrai, V.F., Postnikov, A.M., “Study og Some Ternary Solid Solutions Based on the

Compound Nb3Al” (in Russian), Dokl. Akad. Nauk SSSR, 224, 1130-1133 (1975)

(Experimental, 8)

[1977Ale] Alekseevskii, N.Yu., Ageev, N.V., Shamrai, V.V., “Superconductivity of Some

Three-Component Solid Solutions Based on the Compound Nb3Al”, Phys. Met. Metallogr.,

43(1), 29-35 (1977), translated from Fiz. Met. Metalloved., 43(1), 38-44 (1977)

(Experimental, 14)

[1985Zak] Zakharova, M.I., Badaev, O.P., “Influence of Aluminium on Structure Transformations of

the Solid Solution in Alloy Zr-Nb-Al, Phys. Met. Metallogr., 60(1), 188-200 (1985),

translated from Fiz. Met. Metalloved., 60(1), 199-201 (1985) (Experimental, 0)

[1989Sub] Subramanian, P.R., Simmons, J.P., Mendiratta, M.G., Dimiduk, D.M., “Effect of Solutes on

Phase Stability in Al3Nb”, Mat. Res. Soc. Symp. Proc., 133(3), 51-56 (1989) (Equi.

Diagram, Expermental, 12)

[1990Per] Peruzzi, A., Bolcich, J., “Experimental Determination of the Phase Relationships in

Zr/2.5-8.0 at.% Nb/0-6.7 at.% Al Alloys with 750 at. ppm O and 250 at. ppm N Between

730-900°C”, J. Nucl. Mater., 174, 1-15 (1990) (Equi. Diagram, Experimental, #, 18)

[1992Oka] Okamoto, H., “Nb-Zr (Niobium-Zirconium)”, J. Phase Equilib., 13(5), 577 (1992) (Equi.

Diagram, Review, 8)

[1993Bar] Barth, E.P., Sanchez, J.M., “Obersevation of a New Phase in the Niobium-Alumionium

System”; Scr. Metall. Mater., 28, 1347-1352 (1993) (Crys. Structure, Equi. Diagram,

Experimental, 9)

383


MSIT®

Al–Nb–Zr

[1993Hub] Hubert-Protopopescu, M., Lukas, H.L., Ran, Q., “Aluminium-Niobium-Zirconium”, MSIT

Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials

Science International Services GmbH, Stuttgart; Document ID: 10.16071.1.20, (1993)

(Crys. Structure, Equi. Diagram, Assessment, 13)

[1994Sok] Sokolovskaya, E.M., Kazakova, E.F., Podd'yakova, E.I., Portnoi, V.K.,

Tolmachiova, N.Yu., “Isothermal Section of Al-Nb-Zr System at 770 K” (in Russian), Vest.

Mosk. Univ., Ser. 2: Khim., 35(4), 342-344 (1994) (Equi. Diagram, Experimental, 6)

[1999Tew] Tewari, R., Mukhopadhyay, P., Banerjee, S., Bendersky, L.A., “Evolution of Ordered

Phases in (Zr3Al)-Nb Alloys”, Acta Mater., 47(4), 1307-1323 (1999) (Crys. Structure,

Experimental, 48)

[2000Tew] Tewari, R., Dey, G.K., Ravi, K., Kutty, T.R.G., Banerjee, S., “Hot Hardness and

Indepentation Creep Behaviour of Zr3Al-Nb Alloys”, Trans. Indian Inst. Met., 53(3),

381-389 (2000) (Experimental, Mechan. Prop., 22)



GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram,

Assessment, 168)

[2003Vel] Velikanova, T., Ilyenko, S., “Al-Nb (Aluminium-Niobium)”, MSIT Binary Evaluation



Assessment, 84)


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

( Zr1-x-yNbx-yAly)

( Zr)

1855 - 863

(Nb)

< 2477

cI2

Im3m

W a = 360.90

a = 330.4

at 0 x 1 and 0 y 0.1 at 925°C

[1970Han]

at x = 0, y = 0,


[2003Sch]

at x = 1, y = 0,


[2003Vel]

( Zr)

< 863

hP2

P63/mmc

Mg

a = 323.16

c = 514.75

pure Zr at 25°C [Mas2]


[2003Sch]

(Al)

< 660.452

cF4

Fm3m

Cu

a = 404.96 pure Al at 25°C [Mas2]

Zr3Al

< 1019

cP4

Pm3m

Cu3Au

a = 437.2 0.3 [2003Sch], dissolves small amount of

Nb [1970Han]

384


MSIT®

Al–Nb–Zr

(Zr,Nb)2Al

< 1215

hP6

P63/mmc

Ni2In a = 461.0

c = 591.3

a = 489.39 0.05

c = 592.83 0.05

a = 489.4

c = 592.8

dissolves up to 19 at.% Nb at 925°C

[1970Han],

for Zr59Nb20Al21 [1970Han]

for Zr2Al [2003Sch]

for Zr2Al [1970Han]

(Zr,Nb)5Al3(h)

1400 - ?

tI32

I4/mcm

W5Si3 a = 1087

c = 529.6

a = 1104.4

c = 539.1

a = 1105

c = 539.6

dissolves about 16 at.% Nb at 925°C

[1970Han],


for Zr5Al3 (h) [2003Sch]

for Zr3Al5 (h) [1970Han]

Zr5Al3(r)

?

hP16

P63/mcm

Mn5Si3

a = 817.4

c = 569.8

[2003Sch]

Zr3Al2< 1480

tP20

P42/mnm

Zr3Al2

a = 763.0 0.1

c = 699.8 0.1

[2003Sch], dissolves small amount of

Nb [1970Han]

(Zr,Nb)4Al3

Zr4Al3 1030

hP7

P6/mmm

Zr4Al3 a = 536.8

c = 533.3

a = 543.3 0.5

c = 539.0 0.5


[1970Han],


for Zr4Al3 [2003Sch, 1970Han]

Zr5Al41550 - ~1000

hP18

P63/mcm

Ti5Ga4

a = 844.8

c = 580.5

[2003Sch]

ZrAl

< 1275 25

oC8

Cmcm

CrB

a = 335.9 0.1

b = 1088.7 0.3

c = 427.4 0.1

[2003Sch], dissolves small amount of

Nb [1970Has]

Zr2Al3< 1590

oF40

Fdd2

Zr2Al3

a = 960.1 0.2

b = 1390.6 0.2

c = 557.4 0.02

[2003Sch], dissolves up to 1 at.% Nb at

500°C [1994Sok]

(Zr,Nb)Al2< 1660

hP12

P63/mmc

MgZn2 a = 525.4

c = 869.0

a = 528.24 0.05

c = 874.82 0.05

a = 528.2

c = 874.8


[1970Han],


for ZrAl2 [2003Sch]

for ZrAl2 [1970Han]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

385


MSIT®

Al–Nb–Zr

(Zr,Nb)Al3< 1580

tI16

I4/mmm

ZrAl3 a = 400.5

c = 1727

a = 399.93 0.05

c = 1728.3 0.2

a = 401.0

c = 1731.5


[1993Hub],


for ZrAl3 [2003Sch]

for ZrAl3 [1962Poe]

(Nb,Zr)3Al

Nb3Al

< 2060

cP8

Pm3m

Cr3Si a = 519.7

a = 518.6

a = 518.7

dissolves about 10 at.% Zr at 925°C

[1970Han]

for Zr2.6Nb72.4Al25 [1977Ale]

for Nb3Al [2003Vel]

for Nb3Al [1970Han]

Nb3Al2 1590

tP20

P42/mnm

Al2Zr3

a = 707 8

c/a 0.05

[1993Bar]

42.4 at.% Al, equilibria to be checked

(Nb,Zr)2Al

( phase)

< 1940

tP30

P42/mnm

CrFe a = 995.2

c = 517.4

a = 988.7

c = 516.2

a = 994.3

c = 518.6

dissolves about 15 at.% Zr at 925°C

[1970Han],

for Zr13.5Nb53.5Al33 [1970Han]


for Nb2Al [1970Han]

(Nb1-xZrx)Al3

NbAl3 < 1680

tI8

I4/mmm

TiAl3

a = 387.9

c = 877.1

a = 389

c = 876

a = 387

c = 874

a = 384.1

c = 860.9

a = 384.5

c = 860.1

a = 384

c = 858

at 0 x 0.68 at 925°C [1970Han],

at x = 0.5 [1970Han]

at x = 0.5, annealed at 660°C, two-phase

alloy [1962Poe]

at x = 0.24, annealed at 660°C,

two-phase alloy [1962Poe]

at x = 0 [2003Vel]

at x = 0 [1970Han]

at x = 0 [1962Poe]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

386


MSIT®

Al–Nb–Zr

* 1,

Zr5-2xNb2+xAl6+x

hR13

R3m

W6Fe7

a = 522.7

c = 2830

a = 528.2

c = 2858

a = 529.6

c = 2873

at 0 x 1

at x = 1

[1970Han]

at x = 0.5

[1970Han]

at x = 0.1

[1970Han]

* 2, Zr35Nb30Al35 - - [1970Han]

* 3 cP2

Pm3m

CsCl

Assumed to be stable in the as cast

condition, between 25 to 38 at.% Nb

along section Zr3Al-Nb3Al [1970Has]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

20

40

60

80

20 40 60 80

20

40

60

80

Zr Nb


Axes: at.%

(Zr,Nb)Al3

(Zr,Nb)Al2

Zr2Al

3

ZrAl

Zr3Al

2

(Zr,Nb)2Al

(βZr1-x-y

Nbx-y

Aly)

(Nb,Zr)3Al

(Nb,Zr)2Al

(Nb,Zr)Al3

τ1

τ2

(Zr,Nb)5Al

3

Zr3Al

(Zr,Nb)4Al

3

(Al)

Nb3Al

2

Fig. 1: Al-Nb-Zr.


925°C

387


MSIT®

Al–Nb–Zr

90

10

10

Zr Zr 86.00Nb 14.00Al 0.00

Zr 86.00Nb 0.00Al 14.00 Data / Grid: at.%

Axes: at.%

(αZr)

(αZr)+Zr3Al

(αZr)+(βZr1-x-y

Nbx-y

Aly)+Zr

3Al

(αZr)+(βZr1-x-y

Nbx-y

Aly)

(βZr1-x-y

Nbx-y

Aly)

Fig. 2: Al-Nb-Zr.

Partial isothermal

section at 800°C

90

10

10

Zr Zr 86.00Nb 14.00Al 0.00


Axes: at.%

(αZr)

(αZr)+(βZr1-x-y

Nbx-y

Aly)

(αZr)+(βZr1-x-y

Nbx-y

Aly)+Zr

3Al

(αZr)+(βZr1-x-y

Nbx-y

Aly)

(βZr1-x-y

Nbx-y

Aly)

Fig. 3: Al-Nb-Zr.

Partial isothermal

section at 771°C

388


MSIT®

Al–Nb–Zr

90

10

10

Zr Zr 86.00Nb 14.00Al 0.00


Axes: at.%

(αZr)

(αZr)+Zr3Al

(αZr)+(βZr1-x-y

Nbx-y

Aly)+Zr

3Al

(βZr1-x-y

Nbx-y

Aly)

(αZr)+(βZr1-x-y

Nbx-y

Aly)

4 8 12 20160

1000

Nb, at.%

Tem

pera

ture

,°C

900

800

700

600

1

3( Zr) + ( Zr Nb Al )1- - -x y x y yα β

( Zr Nb Al )β 1- - -x y x y y

2

3

2

1

( Zr Nb Al )β 1-x-y x-y y

( Zr)+( Zr Nb Al )α β 1-x-y x-y y

Fig. 4: Al-Nb-Zr.

Partial isothermal

section at 730°C

Fig. 5: Al-Nb-Zr.

Partial vertical

sections at 0 at.% Al

(1), 3.3.at.% Al (2),

6.7 at.% Al (3)

389


MSIT®

Al–Ni–Ru

Aluminium – Nickel – Ruthenium

Benjamin Grushko

Literature Data

The previous assessment of this system was made in the MSIT Evaluation Program [1993Tre] based on the

data from [1980Tsu, 1985Sok] for 550ºC, partial isothermal sections at 1000 and 1250°C in the Al-poor

region and on a liquidus projection published by [1985Cha1, 1985Cha2, 1986Cha]. No ternary phases and

no significant solubility of the third element in the binary compounds were revealed apart from about 5 at.%

Ru that solve in Ni2Al3 [1980Tsu]. Even NiAl and RuAl, both of the CsCl type structure and alike lattice

parameters, were found to form a wide miscibility gap at 1250°C [1986Cha] which, according to [1980Tsu],

becomes significantly wider at 550°C. However, in [1985Sok] and then in [1998Hor] a continuous range of

solid solutions between these phases was concluded.

A number of the ternary alloys were investigated in [1997Hor1, 1997Hor2]. But the data obtained are of

limited use in the determination of isothermal sections and the homogeneity ranges for the phases because

the experimental technique employed resulted in as-cast, thus not equilibrated samples.

In [1997Poh, 2000Sun1, 2000Sun2, 2001Sun, 2002Sun, 2002Hir] the formation of the several

quasicrystalline phases was observed in the high-Al range of the Al-Ni-Ru system.

Only for the decagonal phase the stability was confirmed by [2003Mi1, 2003Mi2, 2004Mi1] with a

periodicity of 1.6 nm (D4 phase) together with four stable crystalline phases which were determined in the

Al-rich region. The phase equilibria in the temperature range of 700 to 1100°C can be described by partial

isothermal sections built on [2003Mi2]. Additional data are reported by [1997Hor1] from investigations on

alloys of low Al contents.

Binary Systems

The description of the Al-Ni phase equilibria has been accepted from [2003Sal].

According to the recent work [2004Mi2] the Al-Ru system contains six intermediate phases: RuAl6,

Ru4Al13, RuAl2, Ru2Al5, Ru2Al3 and RuAl, that have been previously reported in [1996Bon1, 1996Bon2].

RuAl12 reported in [1963Obr] was not confirmed. Different crystal structures were associated with RuAl2and Ru2Al3. The data accepted in Table 1 originate from [2004Mi2]. Apart from the stable phases a

metastable icosahedral phase (I) and decagonal D1 were reported in this system [1990Wan]. Only RuAl

exhibits a significant compositional range. By heating up to 2100°C the melting point of RuAl was yet not

reached [2004Mi2]. Ru dissolves up to 4 at.% Al [Mas2].

The Ni-Ru system does not contain stable intermediate phases [L-B]. At 1550°C Ni dissolves up to 34.5

at.% Ru while Ru up to 50 at.% Ni. Variations of the lattice parameters of these solid solutions with the

compositions are compiled in [L-B]. A metastable phase was reported in [1979Var] in the range of 30 to

40 at.% Ru.

Solid Phases

At 1500-1600°C the congruent RuAl and NiAl phases form a continuous range of solid solution (Ru,Ni)Al

[1998Hor] which naturally separates the high-Al and low-Al ranges. Considering the very high melting

temperatures [2001Liu] the constitution of the corresponding alloys in equilibrium is unclear even at

1000°C.

At 1100°C the Ru4Al13, RuAl6, and RuAl2 phases dissolve up to 7.0, 0.5 and 0.7 at.% Ni, while Ni2Al3 and

NiAl3 dissolve about 0.7 and < 0.5 at.% Ru, respectively [2003Mi2].

The ternary m phase ((Ru,Ni)2Al9) is isostructural to Co2Al9 and forms at almost constant 82 at.% Al

between 4.5 and 7.0 at.% Ru, the hexagonal H-phase is located in a small range around Ru8.5Ni16.0Al75.5,

and the decagonal D4 phase forms in the vicinity of the H phase around Ru11Ni16Al73 [2003Mi2]. The

orthorhombic O, (Ru,Ni)4Al13 phase (O, Co4Al13 type) is observed in a small compositional range around

390


MSIT®

Al–Ni–Ru

Ru16.0Ni8.0Al76.0 while the C phase (C, Rh2Al5 type) forms in a compositional range of about 9.0-12.0 at.%

Ni and 72.0-73.0 at.% Al [2003Mi2].

The diffraction pattern of the D4 phase is characterized in [2000Sun1, 2000Sun2, 2003Man]. The

experiments made by [2004Mi1] point to a compositional dependence of the D range on temperature from

the melting to at least 700°C. The D1 phase was characterized by [2000Sun1, 2003Man] and an I phase by

[2001Sun]. These phases were not observed in annealed samples [2000Sun1, 2001Sun, 2003Mi1, 2003Mi2]

and are considered to be metastable. The crystallographic data of the ternary phases and their stability

against temperature are listed in Table 1.

It was argued in [2004Mi1] that the stable ternary D4 phase is an extension of a metastable Al-Ru D4 phase.

Pseudobinary Sections

The RuAl-NiAl part of the phase diagram is suggested to be a pseudobinary section, between congruent

RuAl and NiAl, at least at temperatures >1500°C.


A reaction scheme of the Al-Ni-Ru system has been presented in [1993Tre]. However, it had to be revised

for the high-Al range because recently [2003Mi1, 2003Mi2] reported the formation of five ternary phases

by incongruent reactions at the temperatures given in Table 1. The type of the reactions was not established.

The presently accepted reaction scheme for the low-Al part is presented in Fig. 1.

Liquidus Surface

A tentative liquidus projection of the Al-Ni-Ru system was proposed in the previous evaluation published

in [1993Tre]. Considering the ternary phases which now are supposed to form from the liquid, the high Al

part of the liquidus surface can not be accepted anymore.

For the Al-poor part of the system [1997Hor1] published an alternative version of a liquidus projection.

However, considering the solubility data of Ni in (Ru), Ru in (Ni), and Al in (Ni) and (Ru), the indicated

location of the eutectic is improbable because the liquid phase which takes part in the reaction

L Ni3Al+(Ni)+(Ru) does not lie inside the respective tie triangle.

For the same reason, the liquidus projection for the Al-rich part of the diagram constructed in [2000Hoh]

by using data obtained from the as-cast samples can not be accepted. Measured phase equilibria at the

subsolidus temperatures are needed to support a decision.

The Al-poor region of the liquidus projection still applies as described by [1993Tre], see Fig. 2.

Isothermal Sections

In the section of the Al-rich part of the Al-Ni-Ru system at 1600°C given by [1997Hor2] some of the

samples investigated were already liquid. The partial isothermal section at 1250°C is presented in Fig. 3

according to [1993Tre], and the 1100°C isothermal section in Fig. 4 according to [2003Mi2]. The 1000°C

isothermal section (Fig. 5) is combined from the data in [2003Mi2] for high Al compositions and from

[1993Tre] for low-Al compositions considering the continuous b range. The isothermal sections at 900°C

(Fig. 6), 800°C (Fig. 7) and 700°C (Fig. 8) are based on [2003Mi2], however the compositional limits of

the investigated ranges are shifted to higher Al-concentrations in order to reach at lower temperatures

equilibrium.


Potential applications for (Ru,Ni)Al at high temperatures may arise from its high strength, reasonable high

temperature toughness and good oxidation resistance as mentioned by [1997Wol1, 1997Wol2]. Synthesized

by mechanical alloying, the (Ru,Ni)Al alloys with grain sizes of 20-40 nm show a high stability at elevated

temperature [2001Liu].

391


MSIT®

Al–Ni–Ru

Miscellaneous

The difficult mixing of Al and Ru in high-Ru binary and ternary alloys with small fractions of Ni requires

repeated melting and prolonged annealing in order to obtain homogeneous materials.

References

[1963Obr] Obrowski, W., “On the Alloys of Ruthenium with Boron, Berylium and Aluminium” (in

German), Metallwissenschaft und Technik (Berlin), 17(2), 108-112 (1963) (Equi. Diagram,

Crys. Structure, Experimental, 10)

[1965Eds] Edshammar, L-E., “The Crystal Structure of Ru4Al13”, Acta Chem. Scand., 19, 2124-2130


[1966Eds] Edshammar, L-E., “An X-Ray Investigation of Ruthenium-Aluminium Alloys”, Acta

Chem. Scand., 20, 427-431 (1966) (Crys. Structure, Experimental, 3)

[1968Eds] Edshammar, L-E., “The Crystal Structure of RuAl6”, Acta Chem. Scand., 22, 2374-2400


[1979Var] Varich, N.I., Petrunina, A.N., Russ. Metall., 90-91 (1979) (Crys. Structure, Experimental, 3)

[1980Tsu] Tsurikov, V.F., Sokolovskaya, G.M., Kazakova, E.F., “Interaction of Nickel and

Aluminium with Ruthenium” (in Russian), Vestn. Mosk. Univ., Khim., 21(5), 512-514


[1985Cha1] Chakravorty, S., West, D.R.F., “Phase Equilibria Between NiAl and RuAl in the Ni-Al-Ru

System”, Scr. Metall., 19(11), 1355-1360 (1985) (Equi. Diagram, Crys. Structure,

Experimental, 10)

[1985Cha2] Chakravorty, S., Hashim, H., West, D.R.F., “The Ni3Al-Ni3Cr-Ni3Ru Section of the

Ni-Cr-Al-Ru System”, J. Less-Common Met., 20, 2313-2322 (1985) (Equi. Diagram, Crys.


[1985Sok] Sokolovskaya, E.M., Tsurikov, V.F., Orybenkov, S.B., Makanov, U.M., “Phase Diagrams

in Some Systems Containing Aluminum” (in Russian), Stable and Metastable Phase

Equilibria in Metallic Systems, 86(6:72), 79-83 (1985) (Equi. Diagram, 11)

[1986Cha] Chakravorty, S., West, D.R.F., “The Constitution of the Ni-Al-Ru System”, J. Mater. Sci.,

21(8), 2721-2730 (1986) (Equi. Diagram, Crys. Structure, Experimental, #, *, 23)

[1990Wan] Wang, Z.M., Gao, Y.Q., Kuo, K.H., “Quasicrystals of Rapidly Solidified Alloys of Al-Pt

Group Metals – II. Quasicrystals in Rapidly Solidified of Al-Ru and Al-Os Alloys”.

J. Less-Common Met., 163 (1990) (Experimental, Crys. Structure)

[1993Fle] Fleischer, R.L., “Boron and off-Stoichiometry Effects on the Strength and Quality of

AlRu”, Metall. Trans. A, 24A, 227-230, (1993) (Experimental)

[1993Tre] Tretyachenko, L., Sheftel, E., Ibe, G., Grieb, B., Rogl, P., “Aluminum-Nickel-Ruthenium”,

MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI,

Materials Science International Services GmbH, Stuttgart; Document ID: 10.16434.1.20,

(1993) (Crys. Structure, Equi. Diagram, Assessment, 9)

[1996Bon1] Bonniface, T.D., Cornish, L.A., “An Investigation of the High Aluminium end of the Al-Ru

Phase Diagram”, J. Alloys Compd., 233, 241-245 (1996) (Experimental, 11)

[1996Bon2] Bonniface, T.D., Cornish, L.A., “An Investigation of the Al-Ru Phase Diagram above 25

at.% Al”, J. Alloys Compd., 234, 275-279 (1996) (Equi. Diagram, Experimental, 15)

[1997Hor1] Horner, L.J., Cornish, L.A., Witcomb, M.J., “A Study of the Al-Ni-Ru Ternary System

Below 50 at.% Aluminium”, J. Alloys Compd., 256, 213-220 (1997) (Equi. Diagram,

Experimental, 14)

[1997Hor2] Horner, L.J., Cornish, L.A., Witcomb, M.J, “Constitution of the Al-Ni-Ru Ternary System

Above 50 at.% Aluminium”, J. Alloys Compd., 256, 221-227 (1997) (Equi. Diagram,

Experimental, 14)

[1997Poh] Pohla, C., Ryder, P.L., “Crystalline and Quasicrystalline Phases in Rapidly Solidified Al-Ni

Alloys”, Acta Mater., 45, 2155-2166 (1997) (Crys. Structure, Experimental, 48)

392


MSIT®

Al–Ni–Ru

[1997Wol1] Wolff, I.M., “Towards a Better Understanding of Ruthenium Aluminide”, JOM, (1), 34-39

(1997) (Review, 58)

[1997Wol2] Wolff, I.M., Sauthoff, G., Cornish, L.A., Steyn, H. de V., Coetzee, R.,

“Structure-Property-Application Relationships in Ruthenium Aluminide RuAl”, Structural

Intermetallics, 1997, The Minerals, Metals & Materials Society, 815-823 (1997) (Crys.

Structure, Electr. Prop., Experimental, Mechan. Prop., 41)

[1998Hor] Horner, I.J., Hall, N., Cornish, L.A., Witcomb, M.J., Cortie, M.B., Boniface, T.D., “An

Investigation of the B2 Phase Between AlRu and AlNi in the Al-Ni-Ru Ternary System”,

J. Alloys Compd., 264, 173-179 (1998) (Equi. Diagram, Experimental, 23)

[2000Hoh] Hohls, J., Cornish, L.A., Ellis, P., Witcomb, M.J., “Solidification Phases and Liquidus

Surface of the Al-Ni-Ru System Above 50 at.% Aluminium”, J. Alloys Compd., 308,


[2000Sun1] Sun, W., Hiraga, K., “Formation and Structures of Decagonal Quasi-Crystals in the

Al-Ni-Ru System”, Mater. Sci. Eng. A, 294-296, 147-151 (2000) (Crys. Structure,

Experimental, 12)

[2000Sun2] Sun, W., Hiraga, K., “A New Highly Ordered Al-Ni-Ru Decagonal Quasicrystal with 1.6

nm Periodicity”, Philos. Mag. Lett., 80(3), 157-164 (2000) (Crys. Structure,

Experimental, 29)

[2001Liu] Liu, K.W., Muecklich, F., Pitschke, W., Birringer, R., Wetzig, K., “Formation of

Nanocrystalline B2-Structured (Ru,Ni)Al in the Ternary Ru-Al-Ni System by Mechanical

Alloying and its Thermal Stability”, Mater. Sci. Eng. A, 313, 187-197 (2001) (Crys.


[2001Sun] Sun, W., Hiraga, K., “Structural Study of a Superlattice Al-Ni-Ru Decagonal Quasicrystal

Using High-Resolution Electron Microscopy and a High-Angle Annual Dark-Field

Technique”, Philos. Mag. Lett., 81(3), 187-195 (2001) (Crys. Structure, Experimental, 17)

[2002Hir] Hiraga, K., “The Structure of Quasicrystals Studied by Atomic-Scale Observations of

Transmission Electron Microscopy”, Adv. Imag. Electr. Phys., 122, 1-86 (2002) (Review,


[2002Sun] Sun, W., Hiraga, K., “A New Icosahedral Quasicrystal Coexisting with Decagonal

Quasicrystals in the Al-Ni-Ru System”, J. Alloys Compd., 347, 110-114 (2002) (Crys.


[2003Man] Mandal, P., Hashimoto, T., Suzuki, K, Hosono, K., Kamimura, Y., Edagawa, K.,

“Formation of Decagonal and Approximant Phases in the Al-Ni-Ru System”. Philos. Mag.

Lett., 85, 315-323 (2003) (Experimental, Crys. Structure, 24)

[2003Mi1] Mi, S., Grushko, B., Dong, C., Urban, K., “Ternary Al-Ni-Ru Phases”, J. Alloys Compd.,

351, L1-L5 (2003) (Equi. Diagram, Crys. Structure, Experimental, 10)

[2003Mi2] Mi, S., Grushko, B., Dong, C., Urban, K., “Isothermal Sections of the Al-Rich Part of the

Al-Ni-Ru Phase Diagram”, J. Alloys Compd., 359, 193-197 (2003) (Equi. Diagram,

Experimental, 21)





[2004Mi1] Mi, S., Grushko, B., Dong, C., Urban, K., “Phase Equilibrium in the Vicinity of the

Al-Ni-Ru Decagonal Phase”, J. Non-Cryst. Solids., 334-335, 214-217 (2004) (Equi.


[2004Mi2] Mi, S., Balanetskyy, S., Grushko, B., “A Study of the Al-Rich Part of the Al-Ru Alloy

System”, Intermetallics, 11(7), 643-649 (2004) (Equi. Diagram, Crys. Structure,

Experimental, 18)

393


MSIT®

Al–Ni–Ru


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Al)

< 660.452

cF4

Fm3m

Cu

a = 404.96 dissolves 0.01 at.% Ni at 639.9°C

[2003Sal]

(Ru)

< 2334

hP2

P63/mmc

Mg

a = 270.53

c = 428.20

at 25°C [V-C]

dissolves 4 at.% Al at 1920°C [Mas2]

dissolves 50 at.% Ni at ~1500°C [L-B]

(Ni)

< 1455

cF4

Fm3m

Cu

a = 352.40 [Mas2]


[2003Sal]

dissolves 34.5 at.% Ru at ~1500°C

[L-B]

RuAl6< 734

oC28

Cmcm

Al6Mn

a = 748.8

b = 655.6

c = 896.1

[1968Eds]

[2004Mi2]

dissolves < 0.5 at.% Ni [2003Mi2]

Ru4Al13

< 1420

mC102

C2/m

Fe4Al13

a = 1586.2

b = 818.8

c = 1273.6

= 107.88°

[1965Eds], dissolves 7 at.% Ni at

1000°C [2003Mi2]

Ru2Al51340 - 1492

oC*

Cmcm

Fe5Al2

a = 780

b = 660

c = 420

[2004Mi2]

RuAl2< 1805

oF24

Fddd

TiS2

a = 801.2

b = 471.7

c = 878.5

[1966Eds]

[2004Mi2]

Ru2Al3< 1675

tI10

I4/mmm

Os2Al3

a = 307.9

c = 1433

[1966Eds]

[2004Mi2]

, (Ru1-xNix)yAl1-y

RuAl

< at least 2100

NiAl

< 1651

cP2

Pm3m

CsCl

a = 293.9

a = 303

a = 293

a = 299.16

a = 299.16

a = 287

a = 288.72 0.02

a = 287.98 0.02

a = 288.64

0 x 1 [1985Sok]

in Ru29Ni24Al47 annealed at 1600°C

[1998Hor]

[V-C]

in 56 at.% Ru [1986Cha]

[2001Liu]

42 to 69.2 at.% Ni [2003Sal]

[1993Fle]

at 63 at.% Ni

at 50 at.% Ni

at 54 at.% Ni

[2001Liu]

Ni2Al hP3

P3m1

CdI2

a = 407

c = 499

Metastable [2003Sal]

394


MSIT®

Al–Ni–Ru

Ni2Al

< 1372

cP4

Pm3m

AuCu3

a = 358.9

a = 356.32

a = 357.92

a = 356.77

at 63 at.% Ni

disordered

ordered

73 to 76 at.% Ni [2003Sal]

Ni5Al3< 723

oC16

Cmmm

Pt5Ga3

a = 753

b = 661

c = 376

63 to 68 at.% Ni

at 63 at.% Ni [2003Sal]

Ni3Al4< 702

cI112

Ia3d

Ni3Ga4

a = 1140.8 [2003Sal]

Ni2Al3< 1138

hP5

P3m1

Ni2Al3

a = 402.8

c = 489.1

36.8 to 40.5 at.% Ni [2003Sal]

NiAl3< 856

oP16

Pnma

NiAl3

a = 661.3

b = 736.7

c = 481.1

[2003Sal]

NixAl1-x tP4

P4/mmm

AuCu

m**

a = 379.5

c = 325.6

a = 379.5

c = 325.6

a = 375.1

c = 330.7

a = 379.9 to 380.4

c = 322.6 to 323.3

a = 371.7 to 376.8

c = 335.3 to 339.9

a = 418

b = 271

c = 1448

= 93.4°

0.60 < x < 0.68

Martensite, metastable [2003Sal]

at 62.5 at.% Ni

at 63.5 at.% Ni

at 66.0 at.% Ni

at 64 at.% Ni

at 65 at.% Ni [2003Sal]

[2003Sal]

Ni2Al9 mP22

P21/a

Co2Al9

a = 868.5

b = 623.2

c = 618.5

= 96.50º

Metastable [2003Sal]

D1 -

P105mc or

P105/mmc

a = 373.3

c = 407.3

Decagonal in 24-30 at.% Ni [2003Sal]

D4 a = ?

c 1600

Decagonal [2003Sal], contained some Si

, (NiRu) t**

-

-

a = 451.1

c = 362.0

Metastable [1979Var]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

395


MSIT®

Al–Ni–Ru

* (Ru,Ni)2Al9< 783

mP22

P21/a

Co2Al9

a = 863.6

b = 633.3

c = 627.3

= 95.12º

in Ru5,6Ni12Al82 [2003Mi1, 2003Mi2]

* O, Ru4Al13

< at least 1000

oP102

Pmn21

O-Co4Al13

a = 1496.0

b = 825.3

c = 1266.8

Around Ru16.0Ni8Al76 [2003Mi1,

2003Mi2]

*C, (Ru,Ni)2Al51233 to at least 1100 Pm3 or P23

C-Rh2Al5

a = 767.4 High-temperature phase

9.0-12.0 at.% Ni and 72.0-73.0 at.% Al

[2003Mi1, 2003Mi2]

* H

< 930

h** a = 1213.2

c = 2702.0

[2003Mi1, 2003Mi2]

* D4

< 1057 a = 248

c = 1670

Decagonal [2002Sun, 2003Mi1,

2003Mi2]

[2003Man]

* D1 Decagonal Metastable (?) [2000Sun1]

* I Icosahedral Metastable (?) [2002Sun]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

Fig. 1: Al-Ni-Ru. Reaction scheme of the Al-poor part [1993Tre]

l + (Ni) Ni3Al

1372 p4

l + (Ru) (Ni)

1550 p2

L + Ni3Al β + (Ni)T

3<1369 U

2

l β + (Ru)

1920 e1

(Ni) + β + Ni3Al

l Ni3Al + β

1369 e2

L β + (Ru) + (Ni)1250<T4<T

3E

1

L+β+(Ni)

β + (Ru) + (Ni)

L+(Ru)+(Ni)

L+(Ru)+β

Al-Ni-RuNi-Ru Al-Ni Al-Ru

396


MSIT®

Al–Ni–Ru

20

40

60

80

20 40 60 80

20

40

60

80

Ru Ni


Axes: at.%

U2

Ni3Al

E1

β

(Ru)

(Ni)

p2

e1

e2p

4

20

40

60

80

20 40 60 80

20

40

60

80

Ru Ni


Axes: at.%

(Ni)

(Ru)

Ni3Al+(Ni)+β

(Ni)+(Ru)+β

Ni3Al

β

(Ru)+β

(Ru)+(Ni)

β+(Ni)

Fig. 2: Al-Ni-Ru.

Liquidus projection of

the Al-poor part

Fig. 3: Al-Ni-Ru.

Partial isothermal

section at 1250°C

397


MSIT®

Al–Ni–Ru

20

40

60

80

20 40 60 80

20

40

60

80

Ru Ni


Axes: at.%

(Ni)

(Ru)

(Ru)+(Ni)+β

(Ni)+Ni3 Al+β

Ni3Al

β

D

M O

L

Ni2Al

3

(Ru)+β

β+Ni3Al

(Ru)+(Ni)

10

20

30

40

10 20 30 40

60

70

80

90

Ru 45.00Ni 0.00Al 55.00

Ru 0.00Ni 45.00Al 55.00


Axes: at.%

β0

C

LM

Ni2Al

3

RuAl2

Fig. 5: Al-Ni-Ru.

Partial isothermal

section at 1000°C

Fig. 4: Al-Ni-Ru.

Partial isothermal

section at 1100°C

398


MSIT®

Al–Ni–Ru

10

20

30

10 20 30

70

80

90

Ru 40.00Ni 0.00Al 60.00

Ru 0.00Ni 40.00Al 60.00


Axes: at.%

L

Ni2Al

3

HD

M

O

10

20

30

10 20 30

70

80

90

Ru 40.00Ni 0.00Al 60.00

Ru 0.00Ni 40.00Al 60.00


Axes: at.%

L

Ni2Al

3

H

D

MO

NiAl3

Fig. 6: Al-Ni-Ru.

Partial isothermal

section at 900ºC

Fig. 7: Al-Ni-Ru.

Partial isothermal

section at 800ºC

399


MSIT®

Al–Ni–Ru

10

20

10 20

80

90

Ru 30.00Ni 0.00Al 70.00

Ru 0.00Ni 30.00Al 70.00


Axes: at.%L

NiAl3

M

m

RuAl6

Fig. 8: Al-Ni-Ru.

Partial isothermal

section at 700ºC

400


MSIT®

Al–Ni–Si

Aluminium – Nickel – Silicon

Olga Fabrichnaya, Georg Beuers, Christian Bätzner and Hans Leo Lukas

Literature Data

The Al-rich corner was studied several times using thermal and microscopic analyses [1926His, 1930Ota,

1934Fus, 1939Wei, 1942Phi]. The phase relations at Ni contents up to 33.3 at.% have been recently studied

by [2002Ric, 2003Ric1]. A ternary eutectic exists between (Al), NiAl3 and Si. The values given for

temperature and concentration of the eutectic melt are between 560 and 568°C, 3.0 and 5.2 mass% (1.4 and

2.5 at.%) Ni, 11.0 and 11.8 mass% (10.8 and 11.7 at.%) Si. According to the measurements of [2003Ric1]

the temperature of ternary eutectic is 565°C and the composition of the liquid is 2 at.% Ni and 11 at.% Si.

Isopleths are reported for 2 [1942Phi], 6 and 14 mass% Si [1930Ota] and for 2 [1939Wei, 1942Phi], 2.5

[1930Ota], 3 [1990Kuz], 4 [1930Ota, 1939Wei], 5 [1959Phi], 7.5 and 12.5 [1930Ota] mass% Ni. The

isopleths agree well though only [1939Wei] gives a Si solubility in (Al) in agreement with the binary Al-Si

system. [1930Ota] ignores that totally and [1942Phi] gives a much lower value. Recently [2002Ric,

2004Ric] experimentally obtained isoplethal sections for 10, 20, 30, 33.3, 40, 45, 50, 55, 60 and 66.7 at.%

Ni. [1934Fus] gave the Al-rich liquidus surface indicating two more invariant reactions. However the

ternary phase Ni3(Al1-xSix)7 was not taken into account by [1934Fus]. Recently new data on the liquidus

surface were reported by [2003Ric1] at compositions up to 33.3 at.% Ni and between 33.3 and 66.7 at.% Ni

by combination of differential thermal analysis (DTA), powder X-ray diffraction (XRD), metallography and

electron probe microanalysis (EPMA).

The Ni-rich part with more than 50 at.% Ni was investigated by [1959Gua1]. Alloys were melted from

carbonyl-Ni (99.9%), Al of 99.99% and Si of 99.98% purity, annealed at 1100 and 900°C and examined by

metallography and X-ray diffraction. Solid solubility of Al in , Ni2Si, was studied by [1993Bos] and it was

shown that , Ni2Si, could dissolve up to 21 at.% Al. This result has been confirmed by [2002Ric, 2004Ric].

[2004Ric] has reported lattice parameters for the solid solution of Al in ,Ni2Si, as function of composition.

NiAl is reported to dissolve about 15 at.% Si [1959Gua1]. The Si solubility of more than 10 at.% Si in NiAl

is confirmed by [1977Lit], by [2002Ric] (15 % of Si) and by [2004Ric] (20 % of Si). A partial isothermal

section at 750°C with less than 50 at.% Ni content was given in [1969Pan]. The Si solubility in the phase

Ni2Al3 was determined by [1969Pan, 1981Ger, 2003Ric1]. According to [1969Pan] approximately 17 %

Al may be substituted by Si at 750°C, according to [1981Ger] it is 25 % at 600°C. According to recent

measurements of [2003Ric1] 19.2 % Al can be substituted by Si at 550°C that corresponds to 11.5 at.%

solubility of Si in the Ni2Al3 phase. [2004Ric] reported solid solubility of Si in the Ni2Al3 phase to be 18

at.% at 800 and 1000°C. The solubility of Si in NiAl3 was reported to be about 0.6 mass% Si by [1951Pra]

and 0.7 at.% Si by [2003Ric1]. In NiSi2 33 % Si may be substituted by Al [1969Pan, 1981Ger]. According

to [2003Ric1] maximum solubility of Al in NiSi2 at 550°C is 25.7 at.% that means that 38.5 % Si can be

substituted by Al. The large ternary solubilities in NiAl, Ni2Al3 and NiSi2 are compatible with the lattice

parameter data of [1962Wit], although these data do not give exact ranges of homogeneity. Lattice

parameters for NiSi2-xAlx in the whole homogeneity range up to 25.7 at.% Al have been measured by

[2003Ric1].

Binary Systems

The Al-Ni and Al-Si binaries are accepted from [2003Sal, 2003Luk]. The phase diagram for Ni-Si systems

is accepted from [1999Du], but homogeneity ranges for 2 and 3 phase and phase relations involving and

´ phases being adopted from [1987Nas].

Solid Phases

The solid phases are given in Table 1. [1962Wit] mentioned the possibility that NiAl and NiSi2 may have

a common range of homogeneity, regarding the CaF2 structure to be an ordered modification of the CsCl

401


MSIT®

Al–Ni–Si

structure with 50 % vacancies on the Ni sublattice. [1981Ger], however, gave clearly separated fields for

these two phases. These phases have been considered as different phases by [2002Ric, 2003Ric1]. It has

been shown by [2003Ric1] that Ni2Si could dissolve up to 25 at.% Al that corresponds to x = 0.77 for

chemical formula NiSi2-xAlx. According to [2002Ric] NiAl could dissolve Si. Solid solutions containing

~15 at.% Si has been synthesized by [2002Ric] and lattice parameters for these solid solutions has been

measured.

[2004Ric] reported lattice parameters of ternary solid solutions of Si in NiAl at 45, 50 and 55 at.% Ni as

function of composition in the range between 5 and 20 at.% Al.

Some controversy exists regarding the mutual solid solubilities of the isostructural binary phases Ni3Al and

Ni3Si. [1959Gua1] reports that Ni3Al at 1100°C may replace 2/3 of Al by Si. [1959Gua2], however, in

comparing solubilities of different 3rd elements in Ni3Al claimed Si substitution of 50 % of the Al at

1150°C. [1981Ger, 1981Zar] on the other hand, reported a 600°C isothermal section showing really no Si

solubility for Ni3Al. According to [1983Och, 1984Och1, 1984Och2] a continuous solid solution

Ni3Al1-xSix with a linear decrease of the lattice parameter was reported for alloys annealed at 1000°C and

quenched.

The solubility of Al in Ni3Si2 and NiSi was found by [2004Ric] to be very small: 1.0 and 1.5 at.%,

respectively.

A ternary phase Ni2AlSi ( 1) was first reported by [1956Sch, 1957Ess] and confirmed by [1962Wit,

1969Pan, 1981Ger] to have the FeSi structure type. In [1959Gua1] a phase close to this composition was

also mentioned. Lattice parameters for Ni2AlSi phase with different Al and Si contents have been recently

measured by [2002Ric]. Another ternary phase Ni3(Al1-xSix)7 ( 2) (x 0.17) of the Ir3Ge7 type was first

reported by [1962Wit] and confirmed by [1969Pan, 1981Ger, 2003Ric1]. The EPMA results of [2003Ric1]

show that 2 phase exists in a small composition range from 9 to 11.4 at.% Si. The lattice parameters of 2

for compositions of 9 and 11.4 at.% Si are given in [2003Ric1].

A phase ' which is a superstructure of , Ni2-xSi, was reported by [1994Bos] and a formula Ni8-xAlySi4-y

was designated to this phase. The stability of ´ phase has been confirmed by [2002Ric] and crystal

structure has been carefully studied. The formula Ni13 xAlySi9-y and name 3 has been designated to this

phase by [2002Ric].

At 1000°C the extension of the homogeneity range of 3 was found to be much larger than at 800°C

[2004Ric]. Based on experimental results of [2004Ric] there is no evidence for two separate phase fields

for , Ni2-xSi, and 3. Since the structure of , Ni2-xSi, is not completely clear and structure determination

of 3 from quenched samples is only possible in a small part of the homogeneity range, a detailed high

temperature XRD study would be necessary to clarify if one single phase forms or closely related

superstructures.

A phase of approximate composition Ni4AlSi (Ni66Al17Si17) was first mentioned by [1959Gua1] and also

reported by [1981Ger]. The X-ray pattern of this phase was complex and no structural analysis was made.

Later it has been shown by [1993Bos, 2002Ric] that this phase is a part of , Ni2-xSi1-yAly, solid solution.

Richter [2002Ric] has found a new ternary phase ( 4) stable at temperature 550°C, but not at 800°C. The

composition of this phase is Ni61Al4Si35. The observed reflections could be indexed with an orthorhombic

unit cell [2002Ric]. The space group for this phase is reported by [2004Ric].


The invariant eutectic near the Al corner is well established. A partial reaction scheme, based on [1934Fus]

has been recently changed by [2003Ric1] taking into account the ternary phase Ni3(Al1-xSix)7. The partial

reaction scheme at Ni content up to 33.3 at.% based on [2003Ric1] data is presented in Fig. 1a. The partial

reaction scheme for solid state reactions involving 3 and 4 phases is presented in Fig. 1b. The temperatures

and compositions of phases taking part in invariant equilibria involving liquid phase are presented in

Table 2.

402


MSIT®

Al–Ni–Si

Liquidus Surface

The part of the liquidus surface for compositions up to 66.7 at.% Ni is given in Fig. 2 based on works of

[2003Ric1, 2004Ric] but made compatible with the accepted binary systems. The invariant equilibria

containing the Ni3(Al1-xSix)7 ternary phase has been experimentally studied by [2003Ric1]. The invariant

reactions involving the 1 and 3 phases have been studied by [2004Ric].

A systematic investigation was carried out to determine the solvus in Ni-Al-X ternary systems, with X

being transition metal or subgroup B-elements, using the differential thermal analysis (DTA) in [1991Mis].

Solvus isotherms were presented for X = Si, Ga and Ge. In these systems a continuous solid solution was

formed between Ni3Al and Ni3Si. However, in this work the solvus is not reproduced, because there was

inconsistency between figure captions and figures.

Isothermal Sections

An isothermal section of the Ni-rich part (>50 at.% Ni) at 1100°C is given by [1959Gua1]. The (Ni) solvus

is also given for 900°C. However, the isothermal sections based on these data and presented by [1993Beu]

at 900 and 1100°C seem to be inconsistent with new findings of [1993Bos, 1994Bos, 2002Ric] that Ni4AlSi

is a part of solid solution , Ni2-xSi1-yAly, and that there is a field of stability of 3 phase. The Al-rich part

of the 600°C isothermal section presented by [1993Beu] is based on [1939Wei, 1941Han, 1942Phi,

1959Phi], the Al-poor part is based on [1981Ger] with the solubility of Si in Ni3Al changed according to

[1959Gua1]. It should be noted that, isothermal section at 600°C presented by [1993Beu] is also

inconsistent with data of [1993Bos, 2002Ric] concerning the Ni4AlSi phase and the existence of the new

4 phase. The isothermal section at 550°C combined from data [1981Ger, 1993Bos, 2002Ric, 2003Ric1,

2004Ric] is presented in Fig. 3a. According to the accepted Al-Ni binary diagram the Ni3Al4 phase is stable

up to 710°C. This phase was not found in the ternary system by [1981Ger]. Tie lines between Ni3Al4,

Ni1+xAl1-ySiy and Ni2(Al1-xSix)3 are shown tentatively in Fig. 3a. The phase relations at Ni contents up to

33 at.% at 550°C [2003Ric1] are the same as at 600°C [1981Ger]. The only difference is the appearance of

a narrow stability field of the liquid phase in the Al-Si binary at 600°C. The phase relations at higher Ni

content are assumed to be the same at 550 and 600°C because there is no change in phase stability in this

temperature range. This part of phase diagram is accepted from [1993Beu] with corrections made according

to data of [1993Bos, 2002Ric, 2004Ric]. Some modifications have been also made to comply the ternary

phase diagram with the accepted binaries.

The partition of Si between (Ni) and Ni3Al at 1000-1300°C and between Ni3Al and NiAl at 900-1300°C

was investigated using diffusion couples by [1994Jia]. Partition coefficients

and

were determined. It was shown that for the equilibrium between Ni3Al and (Ni) phases partition coefficient

is slightly more than one and decreases with increasing temperature. For the equilibrium between Ni3Al and

NiAl the partition coefficient is more than one at 900-1100°C and less than one at 1300°C.

Isothermal sections at 800 and 1000°C from the experimental study of [2004Ric] are presented in Figs. 3b

and 3c. They are based on XRD and EPMA data. The results of [2003Ric1] obtained at 800°C and Ni

content between 0 and 33.3 at.% were taken into account by [2004Ric]. Besides the liquid phase which is

present in the Al-rich corner of the phase diagram as well as in area adjacent to binary compound NiSi, the

section at 1000°C is dominated by extended solid solution phase fields. As it is mentioned above, the

experimental results by [2004Ric] could not distinguish between the phase fields of 3 and , Ni2Si.

NiAl

Si

AlNi

Si

NiAlAlNi

Si xxK /33 / =

)()/(/33 Ni

Si

AlNi

Si

AlAlNi

Si xxK =

403


MSIT®

Al–Ni–Si


Isoplethal sections at 10, 20, 30 and 33.3 at.% Ni from [2003Ric1] and at 40, 45, 50, 55, 60 and 66.7 at.%

Ni from [2004Ric] are presented in Figs. 4 a-j, slightly modified for consistency with the accepted binary

diagrams.

Thermodynamics

[1984Mar] measured the enthalpy of melting of the ternary eutectic L NiAl3+(Al)+(Si), to be

12.22 kJ (mol-1 of atoms).

The partial enthalpy of Ni at infinite dilution in Al-Si melts was measured by [1985Eml] ranging from

-139.1 kJ mol-1 in pure Al to -140.3 kJ mol-1 in Al+45 at.% Si at 1547°C.

[2000Wit] determined partial and integral enthalpies of mixing of liquid Al-Ni-Si alloys by

high-temperature isoperibolic calorimetry for three sections with constant concentration ratio of Ni and Si

at 1302°C. The results of [2000Wit] are shown in Fig. 5 (partial enthalpies of mixing) and Fig. 6 (integral

enthalpies of mixing). The integral enthalpy of mixing of liquid Al-Ni-Si alloys exhibits a highly negative

and strongly asymmetric dependence on composition with a minimum near Al0.26Ni0.56Al0.18, which gives

evidence of short-range ordering. Using a regular associate model entropy and Gibbs energy of mixing for

liquid Al-Ni-Si alloys have been calculated at 1302°C by [2000Wit]. The contribution of the ternary excess

term is essential and the regular associate model description of enthalpy of mixing of liquid corresponds to

the experimental data only if a ternary associate with the stoichiometry Ni2AlSi is assumed.

The chemical potential of Al in Al-Ni-Si melt was derived from EMF measurements at 900°C and

compositions with different ratio xNi/xSi = 0.066, 0.215 and 1.02. These data are presented in Fig. 7. It shows

that the chemical potential of Al increases at high Al content (xAl > 0.75) and in contrast, decreases when

xNi/xSi increases at low Al content. The derived activity of Al shows negative deviation from ideality.

Addition of Ni to Al-Si alloys increases the deviation from ideality.

The heat capacity of Ni3(Al1-xSix) alloys for x = 0, 0.05, 0.08 and 0.15 from 1.4 to 25 K obtained using

semiadiabatic heat pulse method is presented in Fig. 8.

Calculations of the ternary system have been performed by [1985Kau], however, without taking into

account the ternary phases.


Mechanical properties of Ni3(Al,Si) xSi = 0.025 single crystal with stress axes parallel to crystallographic

orientation near [001] were investigated by both compressive creep and compression tests at temperature of

900°C by [1991Miu]. Magnetic properties of Ni3(Al,Si) at x = 0-0.1 were measured at temperatures

1.8-400 K by [1993Ful]. It was shown that when Si is substituted for Al, the Curie temperature decreases

and goes to 0 K at a critical concentration of about 10 % Si. The electrical resistivity of NiSi2-xAlx phase

was measured at 4.2-300 K at xAl = 0.15, 0.26 and 0.3 by [2003Ric2]. The studied solid solution is a

promising materials for silicon epitaxy as it shows perfect lattice match to Si at composition xAl = 0.26.

The conditions for precipitation of fine ductile (Ni) particles in the Ni3Al matrix were established by

[1998Mer]. This could improve mechanical properties of Ni3Al alloy.

Miscellaneous

The Al-Ni2Si reactions were studied in lateral diffusion couples containing Al islands on Ni-Si multiple

layers by [1990Liu]. The samples were first in situ annealed in transmission electron microscope at

temperatures of 370°C to form Ni2Si phase in the multiple-layer area. Then they were in situ annealed at

temperatures in the range of 498-545°C. During the second-step anneal a sequential formation of NiAl3,

Ni2Al3 and Ni3Si2 was observed. The lateral growth of NiAl3 and Ni2Al3 is a result of Al diffusion in Al-Ni

silicide reaction, the lateral growth of Ni3Si2 is caused by the diffusion of Si atoms dissociated from the

silicides.

Diffusion of Si in the Ni3Al phase has been studied from 900 to 1325°C using the diffusion couple (Ni-24.2

Al (at.%), Ni-22.3Al-3.14Si (at.%)) by [1994Min]. The diffusion profiles in the annealed diffusion couple

404


MSIT®

Al–Ni–Si

were measured by electron probe microanalysis. The diffusion coefficient of Si was derived from the

diffusion profiles and activation energies were calculated.

The effect of alloying elements on the morphological stability of the interface between Ni3Al and NiAl

phases was investigated using ternary diffusion couples annealed at temperatures in the range of

900-1300°C by [2001Kai]. Planar stable interfaces were found in couples with Si.

The structure and thermal stability of rapidly solidified Al-Ni-Si alloys have been investigated using X-ray

diffraction and thermal analysis measurements by [1986Dun]. Series of alloys Ni14Al86-xSix showed a

region of stoichiometry that yields icosahedral symmetry and a region that yields an amorphous phase.

References

[1926His] Hisatsure, C., Suiyókuai Shi, 5, 52 (1926) (Experimental, Equi. Diagram)

[1930Ota] Otani, B., “Silumin and its Structure” (in Japanese), Kinzuku no Kenkyu, 7, 666-686 (1930)


[1934Fus] Fuss, V., “Metallography of Aluminium and its Alloys” (in German), Springer Verlag,

Berlin, 143-145 (1934) (Equi. Diagram, Review, 1)

[1939Wei] Weisse, E., “The Al Corner of the Ternary Al-Ni-Si System” (in German), Aluminium

Archiv, 26, 5-25 (1939) (Experimental, Equi. Diagram, 16)

[1941Han] Hanemann, H., Schrader, A., “On the Ternary Systems of Al” (in German), Z. Metallkd., 33,


[1942Phi] Phillips, H.W.L., “The Constitution of the Aluminium-Rich Alloys of the

Aluminium-Nickel-Iron and Aluminium-Nickel-Silicon Systems”, J. Inst. Met., 68, 27-46

(1942) (Experimental, Equi. Diagram, 15)

[1951Pra] Pratt, J.N., Raynor, G.V., “The Intermetallic Compounds in the Alloys of Aluminium and

Silicon with Chromium, Manganese, Iron, Cobalt and Nickel”, J. Inst. Met., 79, 211-232

(1951) (Experimental, Equi. Diagram, 32)

[1956Sch] Schubert, K., Burkhardt, W., Esslinger, P., Günzel, E., Meissner, H.G., Schütt, W.,

Wegst, J., Wilkens, M., “Some Structural Results on Metallic Phases” (in German),

Naturwissenschaften., 43, 248-249 (1956) (Crys. Structure, 17)

[1957Ess] Esslinger, P., Schubert, K., “On the Systematics of the Structure Family NiAs” (in German),

Z. Metallkd., 48, 126-136 (1957) (Experimental, Review, Crys. Structure, 19)

[1959Gua1] Guard, R.W., Smith, E.A., “Constitution of Nickel-Base Ternary Alloys. III: Ni-Al-Si

System”, J. Inst. Met., 88, 369-374 (1959) (Experimental, Equi. Diagram, #, 5)

[1959Gua2] Guard, R.W., Westbrook, J.H., “Alloying Behavior of Ni3Al ( ' Phase)”, Trans. Met. Soc.

AIME, 215, 807-814 (1959) (Experimental, Equi. Diagram, 27)

[1959Phi] Phillips, H.W.L., “Annotated Equilibrium Diagram of Some Aluminium Alloys Systems”,

Inst. Metall, London, 84-86 (1959) (Equi. Diagram, Review, #, 6)

[1962Wit] Wittmann, A., Burger, K.O., Nowotny, H., “Investigations in the Ternary System, Ni-Al-Si

as Well as of Mono- and Disilicides of Some Transition Metals” (in German), Monatsh.

Chem., 93, 674-680 (1962) (Experimental, Crys. Structure, 20)

[1969Pan] Panday, P.K., Schubert, K., “Structure Investigations in Some Mixtures T-B3-B4 (T = Mn,

Fe, Co, Ir, Ni, Pd; B3 = Al, Ga, Tl; B4 = Si, Ge)” (in German), J. Less-Common Met., 18,


[1977Lit] Litvinov, V.S., Lesnikova, Ye.G., “ Phase Stability in Ni-Al-Si Alloys”, Phys. Met.

Metallogr., 44, 150-153, translated from Fiz. Met. Metalloved., 44, 1297-1299 (1977)

(Experimental, 7)

[1978Bha] Bhan, S., Kudielka, H., “Ordered bcc Phases at High Temperature in Alloys of Transition

Metals and B-Subgroup Elements”, Z. Metallkd., 66, 333-336 (1978) (Experimental, Crys.

Structure, 18)

[1979Ell] Ellner, M., Heinrich, S., Bhargava, M.K., Schubert, K., “Structure Study of the Ni-Si

System” (in German), J. Less-Common Met., 66, 163-173 (1979) (Experimental, Equi.

Diagram, Crys. Structure, 22)

405


MSIT®

Al–Ni–Si

[1981Ger] German, N.V., “Ternary Systems Ni-Si-Al and Co-Si-Al” (in Russian), Vestn. Lvov. Univ.

Ser. Khim., 23, 61-64 (1981) (Experimental, Equi. Diagram, 6)

[1981Zar] Zarechnyuk, O.S., German, N.V., Yanson, T.I., Rychal, R.M., Muravyeava, A.A., “Some

Phase Diagrams of Aluminium with Transition Metals, Rare Earth Metals and Silicon” (in

Russian), Fazovye Ravnovesiya v Metallicheskych Splavach, Nauka, Moscow, 69-73 (1981)


[1983Och] Ochiau, S., Oya, Y., Suzuki, T., “Solubility Data in Ni3Al with Ternary Addition”, Bull.

P.M.E. (T.I.T.), 52, 1-16 (1983) (Experimental, Equi. Diagram, 7)

[1984Mar] Martynova, N.M., Rodionova, E.K., Tishura, T.A., Cherneeva, L.I., “Enthalpy of Melting

of Metallic Eutectics”, Russ. J. Phys. Chem., 58, 616-617 (1984), translated from Zh. Fiz.

Khim., 58, 1009-1010 (1984) (Thermodyn., Experimental, 6)

[1984Och1] Ochiau, S., Mishima, Y., Suzuki, T., “Lattice Parameter Data of Ni( ), Ni3Al( ') and

Ni3Ga( ') Solid Solutions”, Bull. P.M.E. (T.I.T.), 53, 15-28 (1984) (Crys. Structure,

Experimental, 66)

[1984Och2] Ochiau, S., Oya, Y., Suzuki, T., “Alloying Behaviour of Ni3Al, Ni3Ga, Ni3Si and Ni3Ge”,

Acta Metall., 32, 289-298 (1984) (Experimental, Theory, Thermodyn., 90)

[1985Eml] Emlin, B.I., Gizenko, N.V., “Investigation of Melts of Aluminium and Silicon With

3d-Metals and the Improvement of the Process of Production of Cast Alloys” (in Russian),

Fiz. Khim. Issled. Malootkhod. Prots. Electrotkh., Nauka, Moscow, USSR, 186-194 (1985)

(Experimental, Thermodyn., 10)

[1985Kau] Kaufman, L., “Application of Computer Methods for Calculation of Multicomponent Phase

Diagrams of High Temperature Structure Ceramics”, AFOSR-TR-84-0972, 7-11 (1984)

(Theory, 0)

[1986Dun] Dunlap, R.A., Dini, K., “Amorphization of Rapidly Quenched Quasicrystalline

Al-Transition Metal Alloys by the Addition of Si”, J. Mater. Res., 1(3), 415-419 (1986)


[1987Nas] Nash, P., Nash, A., “The Ni-Si (Nickel-Silicon) System”, Bull. Alloy Phase Diagrams, 8,

6-14 (1987) (Review, Equi. Diagram, 59)

[1987Hil] Hilpert, K., Kobertz, D., Venugopal, V., Miller, M., Gerads, H., Bremer, F.T., Nickel, H.,



[1990Kuz] Kuznetsov, G.M., Kalulova, L.M., Mamzurin, O.B., “Phase Equilibria in the Al-Cu-Ni,

Al-Cu-Si, Al-Ni-Si and Al-Cu-Ni-Si System Alloys”, Izv. Vyss. Uchebn. Zaved., Tsvetn.

Metall., 2, 94-100 (1990) (Equi. Diagram, Experimental, Thermodyn., 7)

[1990Liu] Liu, J.C., Mayer, J.W., “Aluminum and Ni-Silicide Lateral Reactions”, J. Mater. Res., 5(2),

334-340 (1990) (Experimental, Equi. Diagram, Phys. Prop., 19)


Ternary Systems”, Mater. Sci. Eng. A, 146, 123-130(1991) (Assessment, Experimental,

Equi. Diagram, 5)

[1991Miu] Miura, S., Hayashi, T., Takekawa, M., Mishima, Y., Suzuki,T., “The Compression Creep

Behavior of Ni3Al-X Single Crystals?”, High-Temp.Ordered Intermetallic Alloys IV,

Mater. Res. Soc. Symp. Proc, 213, 623-628 (1991) (Experimental, Phys. Prop., 9)

[1991Ver] Verhoeven, J.D., Lee, J.H., Laabs, F.C., Jones, L.L., “The Phase Equilibria of Ni3Al

Evaluated by Directional Solidification and Diffusion Couple Experiments”, J. Phase

Equilib., 12, 15-22 (1991) (Experimental, Equi. Diagram, #, 10)

[1993Beu] Beuers, G., Bätzner, C., Lukas. H.L., “Aluminium-Nickel-Silicon”, in MSIT Ternary




[1993Bos] Bosselet, F., Viala, J.C., Colin, C., Mentzen, B.F., Bouix, J., “Solid State Solubility of

Aluminum in the -Ni2Si Nickel Silicate”, J. Mat. Sci. Eng., A167, 147-154 (1993) (Crys.


406


MSIT®

Al–Ni–Si

[1993Ful] Fuller, C.J., Lin, C.L., Mihalisin, T., “Thermodynamic and Magnetic Properties of

(Ni1-xMx)3Al with M=Cu and Pd and Ni3(Al1-xSix)”, J. Appl.Phys., 73(10), 5338-5340

(1993) (Crys. Structure, Experimental, Phys. Prop., 13)

[1994Bos] Bosselet, F., Viala, J.C., Mentzen, B.F., Bouix, J., Colin, C., “ '-Ni8-xSi4-yAly: A New

Ternary Phase Deriving from -Ni2Si in the Al-Ni-Si System”, J. Mat. Sci. Lett., 13,


[1994Jia] Jia, C.C., Ishida, K., Nishizawa, T., “Partition of Alloying Elements Between (A1),

´(L12) and (B2) Phases in Ni-Al Base Systems”, Metall. Mater. Trans. A, 25, 473-485

(1994) (Crys. Structure, Experimental, Equi. Diagram, 25)

[1994Min] Minamino, Y., Yamane, T., Saji, S., Hirao, K., Jung, S.B., Kohira, T., “Diffusion of Cu, Fe

and Si in L1(2)-Type Intermetallic Compound Ni3Al” (in Japanese), J. Jpn. Inst. Met.,

58(4),397-403 (1994) (Crys. Structure, Experimental, Kinetics, 28)

[1998Mer] Merabtine, R., Devaud-Rzepwski, J., Bertrandt, C. Dallas, J.-P., Trichet M.-F., Cornet, M.,

“Ductile Phase Precipitation in the L12 Ternary Intermetallic Alloy Ni3(AlSi)”, J. Alloys

Compd., 278, 75-77 (1998) (Crys. Structure, Experimental, 11)

[1999Du] Du, Y., Schuster, J.C., “Experimental Investigations and Thermodynamic Description of the

Ni-Si and C-Ni-Si Systems”, Met. Trans. A, 88A, 2409-2418 (1999) (Equi. Diagram,


[2000Wit] Witusiewicz, V.T., Arpshofen, I., Seifert, H.J., Sommer, F., Aldinger, F., “Thermodynamics

of Liquid and Undercooled Liquid Al-Ni-Si Alloys”, J. Alloys Comp., 305, 151-171 (2000)


[2001Kai] Kainuma, R., Ichinose, M., Ohnuma, I., Ishida, K., “Formation of ´/ Interface

Morphologies in Ni-Al-X Ternary Diffusion Couples”, Mater. Sci. Eng. A, 312,168-175


[2002Ric] Richter, K.W., “Crystal Structure and Phase Relations of Ni13 xAlySi9-y”, J. Alloys Comp.,

338, 43-50 (Crys. Structure, Equi. Diagram, Experimental, 16)




[2003Ric1] Richter, K.W., Isper, H., “The Al-Ni-Si Phase Diagram Between 0 and 33.3 at.% Ni”,

Intermetallics, 11, 101-109 (2003) (Crys. Structure, Equi. Diagram, Experimental, 10)

[2003Ric2] Richter, K.W., Hiebl, K., “NiSi1.74Al0.26 and NiSi1.83Ga0.17: Two Materials with Perfect

Lattice Match to Si”, Appl. Phys. Lett., 23(3), 497-499 (2003) (Crys. Structure, Electr. Prop.,

Experimental, 13)





[2004Ric] Richter, K.W., Chandrasekaran, K., Ipser, H., “The Al-Ni-Si Phase Diagram. Part II: Phase

Equilibria between 33.3 and 66.7 at.% Ni”, Intermetallics, 12(5), 545-554 (2004) (Crys.

Structure, Experimental, Equi. Diagram, 24)

407


MSIT®

Al–Ni–Si


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Ni)

< 1455

cF4

Fm3m

Cu

a = 352.40 at 25°C [Mas2]

(Al)

< 660.45

cF4

Fm3m

Cu

a = 404.96 at 25°C [Mas2]

(Si)

< 1414

cF8

Fm3m

C-diamond

a = 543.06 at 25°C [Mas2]

Ni3Al1-xSix

Ni3Al

< 1372

1, Ni3Si

< 1035

cP4

Pm3m

Cu3Au

a = 356.55

a = 356.9

a = 350

a = 351

a = 354

0 x 1.0 [1984Och1, 1984Och2]

24.5 to 26 at.% Al at 700°C

[1987Hil] 23.8 to 26.3 at.% Al

at 1200°C [1991Ver]

at x = 0.0 [V-C]

at x = 0 [1993Bos]

at x = 1.0 [1987Nas]

at x = 1.0 [1984Och1]

at x = 0.5 [1959Gua1]

Ni5Al3 700

oC16

Cmmm

Pt5Ga3

a = 744

b = 668

c = 372

32 to 36 at.% Al [Mas, V-C]

Ni1+xAl1-ySiy

NiAl

< 1638

cP2

Pm3m

CsCl

a = 281.6

a = 288.64

a = 286.21

a = 286.32

a = 287.07

a = 286.89

a = 286.85

a = 285.7

a = 285.91 to 282.8

a = 287.85 to 284.8

a = 286.96

-0.35 x 0.55 [Mas]

0 y 0.5 [1962Wit]

30.8 to 58 at.% Al [Mas]

at x = 0; y = 0.5 [1962Wit]

at x = 0; y = 0 [V-C]

at x = 0.2020; y = 0.3303 [2002Ric]

at x = 0.2020; y = 0.3193 [2002Ric]

at x = 0.1739; y = 0.1913 [1993Bos]

at x = 0.2173; y = 0.1729 [1993Bos]

at x = 0.3419; y = 0.1686 [1993Bos]

at x = 0.276; y = 0.1479 [1993Bos]

at x = -0.1818; y = 0.091-0.3636

[2004Ric]

at x = 0; y = 0.1-0.4 [2004Ric]

at x = 0.2222; y = 0.1111 [2004Ric]

Ni2(Al1-xSix)3

Ni2Al3 < 1133

hP5

P3m1

Ni2Al3

a = 400.0

c = 479.1

a = 403.63

c = 490.04

a = 403.65

c = 490.03

a = 401.51

c = 482.31

0 x 0.25 [1962Wit, 1981Ger]

at x = 0.25 [1962Wit]

59.5 to 63.2 at.% Al [Mas]

at x = 0 [V-C] at x = 0 [2002Ric]

at x = 0.19167 [2002Ric]

408


MSIT®

Al–Ni–Si

Ni3Al4 cI112

Ia3d

Ni3Ga4

a = 1140.8 0.1 [2003Sal]

NiAl3< 854

oP16

Pnma

NiAl3

a = 661.14

b = 736.62

c = 481.12

[V-C, Mas] max. solubility

of Si = 0.6 % [1951Pra]

3, Ni3Si(h2)

1200 - 1125

cP2

Pm3m

CsCl

a = 280.08 at 1153°C [1978Bha, V-C]

2, Ni25Si9(h1)

1265 - 975

hR34

hP34

a = 669.8

c = 2885.5

a = 669.8

c = 961.8

90 % of quenched sample [1979Ell]

stacking variant,

10 % present in quenched sample

[1979Ell]

, Ni31Si12

< 1242

hP43

P321

a = 667.1

c = 1228.8

a = 667.9

c = 1222.9

[V-C]

[1993Bos]

, Ni2-xSi(h)

1306 - 825

hP6

P63/mmc

Ni2Si

a = 383.6 to 380.2

c = 494.8 to 486.3

0.37 x 0.68 [1979Ell]

33.4 to 41 at.% Si [Mas2]

parameters of splat cooled

samples [1979Ell]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

409


MSIT®

Al–Ni–Si

, Ni2-xAlySi1-y(r)

Ni2Si

< 1255

oP12

Pbnm

Co2Si

a = 502.2

b = 374.1

c = 708.8

a = 493.2

b = 374.9

c = 716.9

a = 499.5

b = 373.6

c = 707

a = 499.24

b = 374.9

c = 708.51

a = 498.24

b = 374.73

c = 711.04

a = 497.3

b = 374.9

c = 709

a = 497.75

b = 375.16

c = 712.09

a = 497.1

b = 375.61

c = 713.78

a = 496.6

b = 375.89

c = 715.05

a = 495.87

b = 376.92

c = 721.1

a = 492

b = 378.9

c = 732

a = 498

b = 375

c = 711.8

a = 496

b = 376

c = 717.5

a = 495

b = 376.8

c = 722.2

a = 495.8

b = 377.2

c = 723

a = 495.8

b = 378

c = 725.8

at x = 0; y = 0

at x = 0, y = 0.39 [V-C]

at x = 0.0606; y = 0.02939 [2002Ric]

at x = 0.1671; y = 0.1275 [1993Bos]

at x = 0.1751; y = 0.2345 [1993Bos]

at x = 0.2752; y = 0.1260 [1993Bos]

at x = 0.2826; y = 0.337 [1993Bos]

at x = 0.2452; y = 0.3636 [1993Bos]

at x = 0.2376; y = 0.3867 [1993Bos]

at x = 0.1751; y = 0.4689 [1993Bos]

at x = 0.0674; y = 0.6129 [1993Bos]

at x = 0; y = 0.05 [2004Ric]

at x = 0; y = 0.1 [2004Ric]

at x = 0; y = 0.15 [2004Ric]

at x = 0; y = 0.17 [2004Ric]

at x = 0; y = 0.2 [2004Ric]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

410


MSIT®

Al–Ni–Si

´, Ni3Si2(h)

845 - 800

- - [Mas]

, Ni3Si2(r)

< 830

oC80

Cmc21

Ni3Si2

a = 1222.9

b = 1080.5

c = 692.4

a = 1225

b = 1082

c = 693

[V-C]

[1993Bos]

NiSi

< 992

oP8

Pnma

MnP

a = 518

b = 334

c = 562

a = 510.3

b = 333.3

c = 562.8

[V-C]

xAl = 0.015, xSi = 0.485

NiSi2(h)

993 - 981

- - [Mas]

NiAlxSi2-x

NiSi2(r)

< 981

cF12

Fm3m

CaF2

a = 551

a = 540.6

a = 541.5

a = 542.2

a = 542.5

a = 542.5

a = 543.0

a = 543.2

a = 543.8

a = 544.9

a = 546

a = 546.8

a = 547.9

a = 548.2

a = 540.6

0 x 0.77

x = 0.5 [1962Wit]

x = 0 [V-C] [2003Ric1]

x = 0.07[2003Ric1]

x = 0.12 [2003Ric1]

x = 0.15 [2003Ric1]

x = 0.17 [2003Ric1]

x = 0.23 [2003Ric1]

x = 0.3 [2003Ric1]

x = 0.36 [2003Ric1]

x = 0.5 [2003Ric1]

x = 0.53 [2003Ric1]

x = 0.6 [2003Ric1]

x = 0.72 [2003Ric1]

x = 0.75 [2003Ric1]

x = 0 [V-C]

* 1, Ni2AlSi cP8

P213

FeSi

a = 455.9

a = 453.1 to 455.3

a = 453.7

a = 452.99

a = 455.16

[1956Sch, 1957Ess]

[1962Wit]

[1981Ger]

xAl = 0.165, xSi = 0.32 [2003Ric1]

xAl = 0.26, xSi = 0.235 [2003Ric1]

* 2, Ni3(Al1-xSix)7 cI40

Im3m

Ir3Ge7

a = 829.1

a = 829.1

a = 831.59

a = 830.53

x 0.17 [1962Wit]

[1981Ger]

x = 0.1286 [2003Ric1]

x = 0.1629 [2003Ric1]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

411


MSIT®

Al–Ni–Si


* 3, Ni13 xAlySi9-y hP66

P3121

Ga3Ge6Ni13

(designated before

as GaGe2Ni4)

a = 766.3

c = 1467

a = 765.3

c = 1466.5

a = 770.2

c = 1472

a = 770.4

c = 1474

a = 770.2

c = 1474

a = 771.2

c = 1473.2

x = -0.5714; y = 1.0714 [2002Ric]

x = -0.4998; y = 0.9 [2003Ric1]

x = 0.5; y = 1.9125 [2003Ric1]

x = 0.78481; y = 1.93671 [2003Ric1]

x = 1.0769; y = 1.615385 [2003Ric1]

x = 0.5; y = 2.25 [1994Bos]

* 4, Ni61Al4Si35 oC104

Cmcm

Ni16AlSi9

a = 1213.7

b = 1126.5

c = 853.3

[2003Ric1, 2004Ric]

Reaction T [°C] Type Phase Composition at.%

Al Ni Si

L + Ni1+xAl1-ySiy Ni2(Al1-xSix)3 1155 e1(max) L

Ni1+xAl1-ySiyNi2(Al1-xSix)3

60

49

52

29

44

40

11

7

8

L (Si) + NiAlxSi2-x 1085 e2(max) L

(Si)

NiAlxSi2-x

17

0

20

30

0

33

53

100

47

L Ni1+xAl1-ySiy + NiAlxSi2-x 1080 e3(max) L

Ni1+xAl1-ySiyNiAlxSi2-x

25.5

34

21

37.5

34

21

37

21

45

L + Ni1+xAl1-ySiy Ni2(Al1-xSix)3 +

NiAlxSi2-x

1071 U1 L

Ni1+xAl1-ySiyNi2(Al1-xSix)3

NiAlxSi2-x

33

34

44

22

34

45

40

34

33

21

16

44

L + 3 + Ni1+xAl1-ySiy 1 998 P1 L

3

Ni1+xAl1-ySiy

1

13

8

29

25

49

59

50

50

38

33

21

25

L + Ni1+xAl1-ySiy 1 + NiAlxSi2-x 969 U2 L

Ni1+xAl1-ySiy

1

NiAlxSi2-x

13

35

25

19

47

44

50

34

40

21

25

47

L + NiAlxSi2-x NiSi + 1 928 U3 L

NiAlxSi2-x

NiSi

1

6

17

1

23

51

34

50

50

43

49

49

27

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

412


MSIT®

Al–Ni–Si

L NiSi + 3 + 1 925 E1 L

NiSi

3

1

6

1

3.5

20

52

50

57.5

50.5

42

49

39

29.5

L + NiAlxSi2-x Ni2(Al1-xSix)3 + (Si) 839 U4 L

NiSi2-xAlxNi2(Al1-xSix)3

(Si)

56

29

45

0

16

33

40

0

28

38

15

100

L + NiAl3 + Ni2(Al1-xSix)3 2 778 P2 L

NiAl3Ni2(Al1-xSix)3

2

68

75

50

60

12

25

40

30

20

0

10

10

L + Ni2(Al1-xSi)3 2 + (Si) 775 U5 L

Ni2(Al1-xSix)3

2

(Si)

66

50

60

0

12

40

30

0

22

10

10

100

L + 2 NiAl3 + (Si) 659 U6 L

NiAl3(Si)

76

59

74

0

8

30

25

0

16

11

1

100

L (Al) + (Si) + NiAl3 565 E2 L

(Al)

(Si)

NiAl3

87

100

0

74

2

0

0

25

11

0

100

1

Reaction T [°C] Type Phase Composition at.%

Al Ni Si

413


MSIT®

Al–Ni–Si

Fig

. 1a:

Al-

Ni-

Si.

Rea

ctio

n s

chem

e

Al-

Ni

Ni-

Si

Al-

Ni-

Si

Al-

Si

l +

Ni 2

Al 3

NiA

l 3

86

2p

2

L(S

i)+

NiA

l xS

i 2-x

10

85

e 2m

ax

L+

NiA

l xS

i 2-x

Ni 2

Al 3

+(S

i)8

39

U4

L +

(S

i)

NiS

i 2

97

0p

1

l N

iAl 3

+ (

Al)

64

0e 6

L (

Al)

+ (

Si)

57

7e 7

L+

NiA

l 3+

Ni 2

(Al 1

-xS

i x) 3

τ 27

78

P1

L+

Ni 2

(Al 1

-xS

i x) 3

τ 2+

(Si)

77

5U

5

L +

τ2

ΝiA

l 3+

(S

i)6

59

U6

LN

iAl 3

+ (

Al)

+ (

Si)

56

5E

2

NiA

l 3+

(Al)

+(S

i)

NiA

l 3+

Ni 2

(Al 1

-xS

i x) 3

τ 2+

NiA

l 3+

(Si)

NiS

i 2-x

Al x

+N

i 2A

l 3+

(Si)

Ni 2

(Al 1xS

i x) 3

+τ 2

+(S

i)

414


MSIT®

Al–Ni–Si

τ3(θ)+Ni

3Al

1-xSi

xδ+Ni

1+xAl

1-ySi

y930 U

τ3(θ)+Ni

3Al

1-xSi

x+Ni

1+xAl

1-ySi

yτ

3(θ)+Ni

3Al

1-xSi

x+δ

Ni1+x

Al1-y

Siy+Ni

3Al

1-xSi

x+δ

τ3(θ) + Ni

1+xAl

1-ySi

yδ + τ

1U

τ3(θ)+τ

1+Ni

1+xAl

1-ySi

y

Ni1+x

Al1-y

Siy+τ

1+δ

τ3(θ) + δ + Ni

3Si

2(ε/ε') τ

4786 P

δ+τ4+Ni

3Si

2(ε/ε')

τ3(θ) + δ τ

4 + τ

1U

Ni2(Al

1-xSi

x)3+τ

2+(Si)

τ3(θ) + τ

4τ

1+ Ni

3Si

2(ε/ε') U

τ4+τ

1+Ni

3Si

2(ε/ε')

τ3(θ) τ

1+ Ni

3Si

2(ε/ε') + NiSi770 E

τ1+NiSi+Ni

3Si

2(ε/ε')

τ3(θ)+τ

1+NiSi

Al-Ni-Si Ni-Si

NiSi + θ Ni3Si

2(ε/ε')

845 p

θ δ + Ni3Si

2(ε/ε')

825 e

Fig. 1b: Al-Ni-Si. Proposed ternary reaction scheme for solid state reactions according to [2004Ric].

No difference is assumed for ε and ε ' in the Ni-Si binary and for θ and τ3

in the Al-Ni-Si ternary systems.

Temperatures of p and e reactions in the Ni-Si binary system are corrected according to [1987Nas]

20

40

60

80

20 40 60 80

20

40

60

80

Ni Al


Axes: at.%1400°C

1300°C

1200°C

1100°C

1000°C

900°C

800°C

NiAl

p1

e2max

NiAlxSi

2-x

Ni2 Al

3

P2

U4

U5 τ

2

U6 600°C

e7

E2

(Al)

(Si)

p3

p2

e6

NiAl3

e1max

e3maxE

1

U3 U

2

NiSi

τ1 P

1θ/τ3

U1

e5

e4

Fig. 2: Al-Ni-Si.

Partial liquidus

surface projection

including fields of

primary

crystallization

415


MSIT®

Al–Ni–Si

20

40

60

80

20 40 60 80

20

40

60

80

Ni Al


Axes: at.%

ε

τ1

τ4

(Ni)

τ2

NiSi

NiAlxSi

2-x

δγ

NiAl3

Ni5Al

3

(Si)

Ni1+x

Al1-y

Siy

Ni3Al

4 Ni2(Al

1-xSi

x)3

Ni3Al

1-xSi

x(Al)

Ni3Si

Fig. 3a: Al-Ni-Si.


550°C

20

40

60

80

20 40 60 80

20

40

60

80

Ni Al


Axes: at.%

τ1

L

NiSi

NiAlxSi

2-x

(Al)Ni1-x

Al1-y

Siy

Ni2(Al

1-xSi

x)3

NiAl3

εδ

τ3(θ)

(Si)Fig. 3b: Al-Ni-Si.


800°C

416


MSIT®

Al–Ni–Si

20 40 60 80

500

750

1000

1250

Ni 10.00Al 90.00Si 0.00

Ni 10.00Al 0.00Si 90.00Si, at.%

Te

mp

era

ture

, °C

L+NiAl3

L+τ2

L+NiAl3+(Si)

U6

(Al)+NiAl3+(Si)

E2

L+(Si)

L+(Si)+NiAlxSi2-x U4

e2max

970°C

P2

L+τ2+(Si)

L+Ni2(Ai1-xSix)3+(Si)

τ2+(Si)+Ni2(Al1-xSix)3

(Si)+NiAlxSi2-x

NiAl3+τ2+(Si)

L

20

40

60

80

20 40 60 80

20

40

60

80

Ni Al


Axes: at.%

τ4(θ)

Ni1+x

Al1-y

Siy

Ni2(Al

1-xSi

x)3

L

Ni2-x

AlySi

1-y

L

NiAlxSi

2-x

δ

(Si)

Fig. 4a: Al-Ni-Si.


at.% Ni

Fig. 3c: Al-Ni-Si.


1000°C

417


MSIT®

Al–Ni–Si

20 40 60

500

750

1000

1250

Ni 30.00Al 70.00Si 0.00

Ni 30.00Al 0.00Si 70.00Si, at.%

Te

mp

era

ture

, °C

L+NiAlxSi2-x

e2max

L+(Si)

L+(Si)+NiAlxSi2-x

(Si)+NiAlxSi2-x

Ni2(Al1-xSix)3+

(Si)+NiAlxSi2-x

L+NiAl

L

L+Ni2(Al1-xSix)3

L+Ni2(Al1-xSix)3+NiAlxSi2-x

L+Ni2(Al1-xSix)3+(Si)

Ni2(Al1-x)3+τ2+NiAl3

NiAl3+Ni2(Al1-xSix)3

Ni2(Al1-xSix)3+τ2+(Si)

P2 U5

U4

970°C

τ2

Fig. 4c: Al-Ni-Si.


at.% Ni

20 40 60

500

750

1000

1250

Ni 20.00Al 80.00Si 0.00

Ni 20.00Al 0.00Si 80.00Si, at.%

Te

mp

era

ture

, °C

L

L+NiAlxSi2-x

e2max

L+(Si)

970°C

(Si)+NiAlxSi2-x

L+τ2+(Si)

NiAl3+τ2+(Si)

L+τ2+NiAl3

L+(Si)+NiAl3

NiAl3+(Al)+(Si)

L+(Si)+NiAlxSi2-x

L+(Si)+Ni2(Al1-xSix)3

L+Ni2(Al1-xSix)3

L+NiAl3+Ni2(Al1-xSix)3

P2

L+NiAl3 Ni2(Al1-xSix)3+

(Si)+NiAlxSi2-x

Ni2(Al1-xSix)3 +τ2+(Si)

U4

U6

E2

U5

Fig. 4b: Al-Ni-Si.


at.% Ni

418


MSIT®

Al–Ni–Si

20 40

500

750

1000

1250

1500

Ni 40.00Al 60.00Si 0.00

Ni 40.00Al 0.00Si 60.00Si, at.%

Te

mp

era

ture

, °C

L+(Si)

L+NiSi2+

L+NiSi+NiSi2

L

L+Ni1+xAl1-ySiy

L+NiSi2+τ1

L+NiSi2

NiSi+NiSi2

NiS

i 2+τ 1+

NiS

i

NiS

i 2+τ 1

NiSi2+τ1+Ni1+xAl1-ySiy

Ni 1+xA

l 1-y

Siy+

NiS

i2

NiSi2+Ni2(Al1-xSix)3+

Ni1+xAl1-ySiy

Ni2(Al1-xSix)3+NiSi2

(Ni2(Al1-xSix)3)

L+Ni2(Al1-xSix)3+Ni1+xAl1-ySiy

L+Ni1+xAl1-ySiy+NiSi2

(Si)

Fig. 4e: Al-Ni-Si.


at.% Ni

40 50 60

800

900

1000

1100

1200

Ni 33.00Al 32.00Si 35.00

Ni 33.00Al 0.00Si 67.00Si, at.%

Te

mp

era

ture

, °C L+NiAlxSi2-x

L+(Si)+NiAlxSi2-x

L+(Si)

970°C

NiAlxSi2-x

L+Ni2(Al1-xSix)3+NiAlxSi2-x

(Si)+Ni2(Al1-xSix)3+NiAlxSi2-x

U4

Fig. 4d: Al-Ni-Si.


at.% Ni

419


MSIT®

Al–Ni–Si

20 40

500

750

1000

1250

1500

Ni 45.00Al 55.00Si 0.00

Ni 45.00Al 0.00Si 55.00Si, at.%

Te

mp

era

ture

, °C

NiSi+NiSi2

NiSi2+τ1+NiSiNiSi2+τ1+Ni1+xAl1-ySiy

Ni1+xAl1-ySiy+NiSi2

Ni1+xAl1-ySiy

L+(Ni1+xAl1-ySiy)

L+Ni1+xAl1-ySiy+NiSi2

L+NiSi2+τ1

L+NiSi2 L+NiSi

L+NiSi+NiSi2

L+Ni1+xAl1-ySiy+τ1

NiSi2+τ1

L

Fig. 4f: Al-Ni-Si.


at.% Ni

20 40

500

750

1000

1250

1500

Ni 50.00Al 50.00Si 0.00

Ni 50.00Al 0.00Si 50.00Si, at.%

Te

mp

era

ture

, °C

L

L+τ1+NiSi2

NiSi+τ1

NiSi

L+NiSi+NiSi2

L+NiSi

L+τ3(θ)

L+τ1

L+τ1+τ3(θ)

τ1

L+(Ni1+xAl1-ySiy)

L+Ni1+xAl1-ySiy+τ3(θ)

Ni1+xAl1-ySiy

Ni1+xAl1-ySiy+τ1

L+NiSi2

U3P1

Fig. 4g: Al-Ni-Si.


at.% Ni

420


MSIT®

Al–Ni–Si

10 20 30 40

500

750

1000

1250

1500

Ni 55.00Al 45.00Si 0.00

Ni 55.00Al 0.00Si 45.00Si, at.%

Te

mp

era

ture

, °C

Ni1+xAl1-ySiy+

Ni1+xAl1-ySiy+τ3(θ)+τ1

τ1+τ4+Ni3Si2(ε)

Ni2-xAlySi1-y(δ)+Ni1+xAl1-ySiy+τ3(θ)

Ni1+xAl1-ySiy+τ3(θ)

τ1+τ4+Ni2-xAlySi1-y(δ)

τ3(θ)

τ3(θ)+NiSi

τ1+τ3(θ)+NiSi

L+τ3(θ)

L+τ3(θ)+τ1

L+τ3(θ)+NiSi

Ni3Si2(ε)+NiSi+τ1

τ3(θ)+τ1

Ni3Si2(ε)+NiSi+

Ni2-xAlySi1-y(δ)+Ni1+xAl1-ySiy+τ1

Ni2-xAlySi1-y(δ)

L+Ni1+xAl1-ySiy

Ni1+xAl1-ySiy


L

P1

E1

Fig. 4h: Al-Ni-Si.


at.% Ni

10 20 30

500

750

1000

1250

1500

Ni 60.00Al 40.00Si 0.00

Ni 60.00Al 0.00Si 40.00Si, at.%

Te

mp

era

ture

, °C

L


L+τ3(θ)

τ3(θ)


L+Ni1+xAl1-ySiy

Ni1+xAl1-ySiy+τ3(θ)+Ni3Al

Ni1+xAl1-ySiy

Ni1+xAl1-ySiy+Ni3Al

Ni1+xAl1-ySiy+Ni2-xAlySi1-y(δ)

Ni1+xAl1-ySiy+Ni2-xAlySi1-y(δ)+τ3(θ)

Ni1+xAl1-ySiy+Ni2-xAlySi1-y(δ)+Ni3Al

Ni2-x AlySi1-y(δ)+Ni1+xAl1-ySiy+τ1

Ni2-xAlySi1-y(δ)+τ1+τ4

Ni3Si2(ε)+τ1+τ4

Ni3Si2(ε)

Fig. 4i: Al-Ni-Si.


at.% Ni

421


MSIT®

Al–Ni–Si

10 20 30

500

750

1000

1250

1500

Ni 66.70Al 33.30Si 0.00

Ni 66.70Al 0.00Si 33.30Si, at.%

Te

mp

era

ture

, °C


L+(Ni1+xAl1-ySiy)

Ni1+xAl1-ySiy

τ3(θ)

Ni2-xAlySi1-y(δ)


Ni1+xAl1-ySiy+Ni2-xAlySi1-y(δ)+τ3(θ)Ni3Al+Ni1+xAl1-ySiy

Ni3Al+Ni1+xAl1-ySiy+Ni2-xAlySi1-y(δ)

L+τ3(θ)


L

?

Ni3Al+

Fig. 4j: Al-Ni-Si.

Vertical section at

66.7 at.% Ni

Al, at.%

∆HN

i,kJ

mo

l·

-1

0 20

-40

0

Ni Si1-y y

-80

-120

-160

40 60 80 100

Al

Al-NiAl-Ni SiAl-Ni SiAl-Ni Si

0.8

0.5

0.2

0.2

0.5

0.8

Al, at.%

∆HN

i,kJ

mo

l·

-1

0 20

-40

0

Ni Si1-y y

-80

-120

-160

40 60 80 100

Al


0.8

0.5

0.2

0.2

0.5

0.8

Al, at.%

∆HN

i,kJ

mo

l·

-1

0 20

-40

0

Ni Si1-y y

-80

-120

-160

40 60 80 100

Al


0.8

0.5

0.2

0.2

0.5

0.8

Fig. 5a: Al-Ni-Si.

Partial enthalpy of

mixing of nickel of

ternary liquid and

undercooled liquid

Al-Ni-Si alloys at

1302 3°C. Standard

states: Al(l), Ni(l) and

Si(l) [2000Wit]

422


MSIT®

Al–Ni–Si

Al, at.%

∆HA

l,kJ

mo

l·

-1

0 20

-40

0

Ni Si1-y y

-120

40 60 80 100

Al


0.8

0.5

0.2

0.2

0.5

0.8Al-Si

-80

-160

Fig. 5c: Al-Ni-Si.

Partial enthalpy of

mixing of aluminum

of ternary liquid and

undercooled liquid

Al-Ni-Si alloys at

1302 3°C. Standard


Si(l) [2000Wit]

Al, at.%

∆HS

i,kJ

mo

l·

-1

0 20

-50

0

Ni Si1-y y

-200

40 60 80 100

Al

Al-SiAl-Ni SiAl-Ni SiAl-Ni Si

0.8

0.5

0.2

0.2

0.5

0.8

-150

-100

-250

Fig. 5b: Al-Ni-Si.

Partial enthalpy of

mixing of silicon of

ternary liquid and

undercooled liquid

Al-Ni-Si alloys at

1302 3°C. Standard


Si(l) [2000Wit]

423


MSIT®

Al–Ni–Si

20

40

60

80

20 40 60 80

20

40

60

80

Ni Al


Axes: at.%

-10-20

-30-40

-50

-10

-20

-30

-40

-50

-60

-69

Fig. 6b: Al-Ni-Si.

Isolines for integral

enthalpy of mixing

based on

experimental data of

[2000Wit]

Al, at.%

∆H,

kJ

mo

l·

-1

0 20

-10

0

Ni Si1-y y

40 60 80 100

Al

-20

-30

-40

-50

-60

-70


0.8

0.5

0.2

0.2

0.5

0.8Al-Si

Fig. 6a: Al-Ni-Si.

Integral enthalpy of

mixing of liquid and

undercooled liquid

Al-Ni-Si alloys at

1302 3°C. Standard


Si(l) [2000Wit]

424


MSIT®

Al–Ni–Si

25

50

45

40

35

30

55

10025 75

T K2 2,

C/T

,m

Jm

ol

K-1

-2·

·

Ni (Al Si )3 1-x x

x=0.05

x=0.15

x=0.08

0 50

0

-2

-4

-8

-6

0.80.6 1.0

XAl

µ Al

-1,

kJ

mo

l·

1

2

3

Fig. 7: Al-Ni-Si.

Partial molar free

enthalpy of Al in

Al-Ni-Si melts at

900°C with respect to

mole fraction of Al

and p=0.066 (1),

0.215 (2) and 1.020

(3)

Fig. 8: Al-Ni-Si.

C/T vs T2 for

Ni3(Al1-xSi) with

x=0.05, 0.08 and 0.15

425


MSIT®

Al–Ni–Ta

Aluminium – Nickel – Tantalum

Viktor Kuznetsov

Literature Data

Phase equilibria and intermetallic phase formation has been reviewed by [1990Kum]. However, this was

followed by a thorough assessment of the data published up to 1991 by [1993Zak]. They presented

graphically the Ni3Al-TaNi3 section, an assessed Scheil reaction scheme, liquidus and solidus projections,

the solvus of the ´ (Ni3Al) based phase and two partial isothermal sections for 1000 and 1250°C. The

existence of six ternary phases, TaNiAl, TaNi2Al, Ta0.5Ni3Al0.5, Ta5Ni2Al3, Ta~55Ni~10Al~35 and TaNiAl2was accepted. However, some earlier work was not mentioned in [1993Zak]. [1965Ram] had indicated that

in addition to the TaNiAl and TaNi2Al phases, which had been established initially by [1964Mar], a phase

with a structure “closely resembling” that of NiTi2 (in Table 1 of [1965Ram] denoted as NiTi2) was present

in alloys of gross compositions Ta25Ni38Al37 and Ta25Ni25Al50. The phase was found in both the as cast

state and after annealing for 7 days at 900°C, but with an amount significantly less after the heat treatment.

Moreover, it was found in the as cast sample with a composition of Ta25Ni50Al25, but later transformed

almost entirely to TaNi2Al after annealing. Unfortunately, no compositional data for the phase was given.

Later, [1974Ali] performed a DTA study of 5 alloys in the Ni3Al-TaNi3 section in the course of a study of

the Ni3Al-Ni3Ta-Ni3Nb pseudoternary system.

In more recent years the phase equilibria in this system have been investigated in much detail. [1994Joh]

studied five arc-melted alloys with compositions close to NiAl+15 at.% Ta (on the eutectic line) in the as

cast and directionally solidified state by using scanning electron microscopy with EDS to measure phase

composition. From the results, a fragment of the liquidus projection (for NiAl-Ni2TaAl-NiTaAl

composition region) was constructed suggesting a peritectic formation for the Ni2AlTa ternary phase.

[2001Miu] used DTA to determine liquidus and solidus temperatures of alloys made by arc-melting Al, Ni

and Ta of purities 99.99, 99.95 and 99.9 mass%, respectively, followed by a homogenization treatment of

1000°C for 24 h. [1991Mis] determined the solvus line of the phase at temperatures between 827 and

1327°C using DTA. Energy-dispersive X-ray spectroscopy was used to confirm the phase constitution of

the alloys. [1994Jia] studied the partition of Ta between and ´, as well as between the ´ and phases

using a diffusion couple technique. The results are presented in tabular form with phase composition and

partition coefficients for 1300, 1200, 1100, 1000 and 800°C and also rendered graphically as partial sections

for some selected temperatures. [1996Pal] re-investigated two partial isothermal sections for Ta contents of

< 50 at.% for 1000 and 1250°C in order to confirm the work of [1993Zak]. 32 compositions were prepared

from components of purities of 99.95 mass% Ni, 99.99 mass% Al and 99.97 mass% Ta using levitation

melting. Heat treatment at 1000°C was performed in Ar filled silica ampoules for 168 h for alloys in the

NiAl+TaNiAl composition region and for 500 h for alloys of all other compositions. Water quenching

followed the heat treatment. At 1250°C, the heat treatment was carried out in a box made from Ta sheets;

each specimen was wrapped into Ta foil, and the box was filled with Ti-filings. The heat treatment was

carried out in an Ar atmosphere for between 100 and 20 h with subsequent cooling under flowing gas.

Samples were examined by metallography, X-ray diffraction and electron microprobe. The results show

significant differences from the assessed data of [1993Zak].

[1999Sun] studied the partition of Al and Ta between the liquid and fcc phases in samples quenched from

the two phase liquid + fcc state. The compositions of the phases were measured by EPMA. Equilibrium

conditions were confirmed by the measure of homogeneity of the solid phase. In addition, they performed

a simultaneous regression analysis of their own data, the published data of [1993Zak] and data for the

Ni-Cr-Al-Ta quaternary. Good agreement (within approx. 1%, i.e. 7 to 10 K) between the different sets was

found. Data for liquid compositions and partition coefficients for Al and Ta were tabulated.

Very little work has been carried out on the thermodynamic properties in this system. [1999Roc] measured

the low-temperature (3.2 to 10.3 K) heat capacity of the TaNi2Al phase and calculated its electron structure

by the LMTO technique. Combining the results of both, the electron-phonon interaction constant was

426


MSIT®

Al–Ni–Ta

derived. Some phase boundaries have been calculated using CALPHAD and ab initio techniques.

[1991Kau] performed an approximate CALPHAD calculation of the phase equilibria. However, ternary

phases were not taken into account, although dissolution of Al in the TaNi binary compound was allowed

in the calculation. [1991Eno] calculated the equilibria between the and ´ phases at 1000°C, using the

cluster variation method based on empirical Lennard-Jones type interatomic pair potentials. Good

agreement with experimental data was obtained.

A number of investigations of mechanical properties have been made. [1991Sas] noticed the precipitation

of Ta enriched phase whilst studying the mechanical properties of (NiAl)0.95Ta0.05. [1996Mac] measured

lattice spacing and mechanical properties of the TaNiAl ternary phase. Mechanical properties were also

studied by [1991Bon], [1991Hay], [1991Mas], [1991Sas].

Binary Systems

The Ni-Ta system is taken from [Mas2], [1991Nas]. For the Al-Ni binary, the latest version [2003Sal]

evaluated within the MSIT Binary Evaluation Program is accepted; it does not differ significantly from that

of [1987Hil, 1988Bre], which was used by [1993Zak]. The Al-Ta system is taken from [2003Cor], who

accepted results of the thermodynamic assessment of the system performed by [1996Du].

Solid Phases

[1993Zak] accepted the existence of six ternary phases, TaNiAl, TaNi2Al, Ta0.5Ni3Al0.5, Ta5Ni2Al3,

Ta~55Ni~10Al~35 and TaNiAl2. The TaNiAl phase has a wide solubility range for Al (11 to 50 at.%), but

restricted for Ta (32.5 to 37.5 at.%). [1996Pal] noted, that in comparing calculated and observed intensities

of X-ray diffractions lines, the suggestion is that Al substitutes for Ni on two different crystallographic sites

which exist in the MgZn2 structure to a similar extent. The lattice constants of that phase seem to depend

on cooling rate; the reason for this is unclear, but because no peak broadening was observed, it is not likely

to be due to stacking faults introduced by thermal stresses on cooling [1996Pal].

The true composition of the TaNi2Al phase was found to be off-stoichiometric: 51 to 55% Al and 22.5 to

25% Ta at 1000°C; 52 to 58% Al and 17.5 to 24% Ta at 1250°C [1996Pal].

[1996Pal] did not find any trace of the Ta5Ni2Al3, Ta~55Ni~10Al~35 and TaNiAl2 phases as presented by

[1993Zak], nor the NiTi2 phase reported by [1965Ram]. The existence of the first three was explicitly

rejected; the latter was not considered anyhow by [1996Pal], but no such phase was detected in the

composition range studied by [1965Ram]. As [1993Zak] noted weak support for the existence of

Ta5Ni2Al3, Ta~55Ni~10Al~35 and TaNiAl2, these phases are considered in present review to be non-existent.

The ternary phase proposed by [1965Ram] seems to be metastable; also, phases with that structure are often

stabilized by impurities such as C, N or O.

Crystallographic information for the solid phases, including the probably metastable ternary phase, is

summarized in Table 1. Detailed data for the concentration dependence of the lattice spacing of TaNiAl

[1996Pal] are given on Fig. 1. For 20 < xNi < 50 (xNi in at.%) that dependence is essentially linear:

a (pm) = 487.6+0.413(50-xNi), c (pm) = 791.5+0.476(50-xNi), though marked deviations from that can be

seen for less Ni [1996Pal]. For lattice spacing of the phase, linearity holds for all compositions studied:

a (pm) = 487.5+0.386(50-xNi), c (pm) = 2653+2.90(50-xNi) (also for 20 < xNi < 50 at.%).


No pseudobinary sections have been found in the system, though some authors suggested such behavior for

the Ni3Al-TaNi3 section, see section “Temperature – Composition Sections”.


Data for the invariant equilibria and Scheil reaction scheme (Fig. 2) were assessed by [1993Zak], and are

accepted here with some alterations. Table 2 is based on [1993Zak], but with a corrected error in the

temperature of the U4 reaction, noted by [1996Pal]. The reaction U5, presented by [1993Zak] has been

omitted as it was shown to be unlikely by [1996Pal]. The eutectic e1(min) L +TaNiAl, is added from

427


MSIT®

Al–Ni–Ta

[1994Joh]. It is not possible to include the ternary peritectic reaction +TaNiAl+L TaNi2Al and

surrounding univariant eutectic L +TaNiAl and peritectics +L TaNi2Al and TaNiAl+L TaNi2Al

suggested by [1994Joh] in either Table 2 or the reaction scheme as neither temperatures nor phase

compositions were determined. (See however discussion of liquidus below).


The liquidus data from [2001Miu] are in good agreement with [1993Zak] for both edge systems, but differ

markedly for intermediate compositions. The data of [2001Miu] are preferred as they result from detailed

work and seem to be more reliable. On the other hand, [2001Miu] presents mono- and invariant equilibria

lines taken from [1986Wil] which is the main source for [1993Zak]. The liquidus surface is presented here

in Fig. 3. It is a composite of the liquidus taken from [2001Miu] and the liquidus surfaces of ´, 3 and

phases taken from [1993Zak]. The partial liquidus projection from [1994Joh] is added tentatively, although

its connection with other parts of liquidus surface remains rather unclear.

Figure 4 provides isotherms of the solidus from [2001Miu]. Figures 5 and 6 present the data of [2001Miu]

showing the dependence of the liquidus and solidus temperatures on Al variation at parametric Ta content,

and on Ta variation at parametric Al content, respectively. These data give more detailed representation than

is possible in Figs. 3 and 4. Figure 7 presents the isotherms of the /( + ´) solvus surface as determined by

[1991Mis].

Isothermal Sections

Isothermal sections at 1273 and 1000°C are presented in Figs. 8 and 9, respectively, generally accepted from

[1996Pal]. The results differ significantly from those of the earlier assessment of [1993Zak]. On the other

hand, the data given in the original work disagree with the accepted Al-Ta binary system (and even with the

binary accepted by the authors [1996Pal] themselves). To maintain consistency, it was necessary to change

the region adjacent to Al-Ta system, which in any event is based on just two alloys. In particular, the

homogeneity range of TaAl3 phase is removed, and that of Ta2Al3 is split into stoichiometric Ta5Al7 and

Ta39Al69 phases at 1250°C (Fig. 8) and into Ta5Al7 and Ta2Al3 at 1000°C (Fig. 9). These changes were

suggested by [1996Du] who analyzed the results of [1996Pal] during their assessment of the Al-Ta binary

and is accepted here. Also, the position of the phase corners of +TaNiAl+TaAl3 and L+TaNiAl+TaAl3tie triangles had to be shifted somewhat to make them compatible with the accepted version of the Al-Ni

binary.

The data of [1994Jia] for - ´ and ´- equilibria, presented in tabular form, are reproduced in Tables 3

and 4.


[1993Zak] suggested the section Ni3Al-Ni3Ta to be “partly pseudobinary” and mentioned some

experiments on directional growth of a “pseudobinary eutectic” [1972Hub, 1974Mol]; the reported

composition of the latter is indeed in very good agreement with the composition of the e2 reaction of

[1993Zak]. The DTA study of [1974Ali] is also in agreement, though the authors themselves interpreted

their results as indication of a simple pseudobinary section with a single eutectic. As indicated by

[1993Zak], the Ni3Al-Ni3Ta section cannot be pseudobinary due to the incongruent formation of Ni3Al.

Moreover, in the presently accepted version of the Al-Ni binary, the Ni3Al phase becomes

off-stoichiometric starting from approx. 1347°C up to the melting point [1987Hil, 1988Bre]. Also, the phase

boundaries of ternary TaNi2Al phase as determined for 1000°C by [1996Pal] are not crossed by the

Ni3Al-Ni3Ta join. No account of these phenomena was taken by [1993Zak]. On the other hand, the assessed

liquidus-solidus region of that section is indeed independently confirmed by the results of [1974Ali] and by

directional solidification experiments, reported by [1993Zak]. So, this fragmentary section is reproduced

from [1993Zak] with minor corrections and given as Fig. 10, though the true phase relations should be much

more complicated both in a region closer to the Ni3Al side and at lower temperatures.

428


MSIT®

Al–Ni–Ta

Thermodynamics

No thermodynamic studies have been carried out except for low-temperature (3.2 to 10.3 K) measurements

of the heat capacity of the TaNi2Al phase performed by [1999Dar]. Their results, when treated in the

standard way (Cp(T) = elT + CD( /T)), give el = 10.01 0.14 mJ mol K-2, D = 299 1.9 K. This

equation is valid only below approximately 7 K.


The influence of Ta additions on mechanical properties of NiAl was studied in [1991Mas, 1991Sas]. Such

properties of Ta alloyed single crystals of ´ Ni3Al can be found in [1991Bon]; creep behavior of that phase

was studied by [1991Hay]. Some mechanical properties of the Laves phase TaNiAl were measured by

[1996Mac].

Miscellaneous

[2001Ter] suggested the usage of thermal conductivity measurements for determination of site preferences

in the ´ Ni3Al phase. The results are in broad agreement with the phase diagram determinations, which

suggest that Ta substitutes for Al in Ni3Al.

[2001Kai] investigated the morphological stability of the interface between ´(L12) and (B2) phases in

diffusion couples. In addition, the results of an unpublished calculation of thermodynamic properties are

cited and used in the discussion of the results.

References


of the MnCu2Al and MgZn2 Types Containing Aluminium and Gallium”, Sov.

Phys.-Crystallogr., 9, 619-620 (1964), translated from Kristallografiya, 9, 737-738 (1964)

(Crys. Structure, 4)

[1965Gie] Giessen, B.C., Grant, N.J., “New Intermediate Phases in Transition Metal Systems. II”, Acta

Crystallogr., 18, 99 (1965) (Crys. Structure, 4)

[1965Ram] Raman, A., Schubert, K., “On the Crystal Structure of Some Alloy Phases Related to TiAl3.

III. Investigations in Several T-Ni-Al and T-Cu-Al Systems” (in German), Z. Metallkd., 56,


[1968Hun] Hunt, C.R., Raman, A., “Alloy Chemistry of ( U)-Related Phases. I. Extens Ion of - and

Occurrence of ´-Phases in the Ternary Systems Nb(Ta)-X-Al (X = Fe, Co, Ni, Cu, Cr,

Mo)”, Z. Metallkd., 59(9), 701-707 (1968) (Crys. Structure, Equi. Diagram, 14)

[1972Hub] Hubert, J.-C., Kurz, W., Lux, B., “Growth by Directed Solidification of the Ni3Al-Ni3Ta

Quasibinary Eutectic” (in French), J. Cryst. Growth, 13-14, 757-764 (1972) (Equi.

Diagram, 15)

[1972Min] Mints, R.S., D´yakonova, N.P., Umansky, Ya.S., Bondarenko, Yu.A., Bondarenko, T.A.,

“Interaction of the Phase Ni3Al with Ni3Ta”, Sov. Physics Doklady, 17(9), 904-906 (1973)

translated from Dokl. Akad. Nauk SSSR, 206(1), 87-88 (1972) (Crys. Structure,

Experimental, 5)

[1974Ali] Alikhanov, V.A., Pyatnitskii, V.N., Sokolovskaya, E.M., “Phase Diagram of the System

Ni3Al-Ni3Nb-Ni3Ta” (in Russian), Vestn. Mosk. Univ., Ser. 2:Khim., 15, 698-701 (1974)


[1974Mol] Mollard, F., Lux, B., Hubert, J.C., “Directionally Solidified Composites Based on the

Ternary Eutectic Ni-Ni3Al-Ni3Ta ( / ´ - )”, Z. Metallkd., 65, 461-468 (1974) (Equi.


[1974Var] Varli, K.V., D’yakonova, N.P., Umansky, Ya.S., Bondarenko, Yu.A., Putman, A.M.,

“Crystal Structure of the Ternary Phase of the Ni-Ta-Al System”, Vses. Konf. Kristallokhim.

Intermet., 2nd, Tezisy Dokl., Lvov Gos. Univ.: Lvov, USSR, 49 (1974) (Crys. Structure, 0)

429


MSIT®

Al–Ni–Ta

[1979Nas] Nash, P., West, D.T.F., “Phase Equilibria in the Ni-Ta-Al System”, Met. Sci., 13(12),


[1984Och] Ochiai, S., Mishima, Y., Suzuki, T., “Lattice Parameter Data of Ni ( ), Ni3Al ( ´) and

Ni3Ga ( ´) Solid Solutions”, Bull. P. M. E.,(T. I. T.), 53, 15-28 (1984) (Crys. Structure,

Experimental, Rewiew, 56)

[1984Wil] Willemin, P., Dugue, O., Durand-Charre, M., Davidson, J., “High-Temperature Phase

Equilibria in the Ni-Al-Ta System”, Superall. 1984 Champ., MS/AIME, Conf: Pa. USA,


[1985Mis] Mishima, Y., Ochiai, S., Suzuki, T., “Lattice Parameters of Ni( ), Ni3Al( ´) and Ni3Ga( ´)

Solid Solutions with Additions of Transition and B-Subgroup Elements”, Acta Metall.,

33(6), 1161-1169 (1985) (Crys. Structure, Review, 64)

[1986Hua] Huang, S.C., Briant, C.L., Chang, K.-M., Taub, A.I., Hall, E.L., “Carbon Effects in Rapidly

Solidified Ni3Al”, J. Mater. Res., 1(1), 60-67 (1986) (Experimental, Mechan. Prop., 27)

[1986Wil1] Willemin, P., Dugue, O., Durand-Charre, M. J., Davidson, H., “Experimental

Determination of Nickel-Rich Corner of Ni-Al-Ta Phase Diagram”, Mater. Sci. Technol.,

2(4), 344-348 (1986) (Equi. Diagram, 13)

[1986Wil2] Willemin, P., Durand-Charre,, M., Ansara, I., “Liquid-Solid Equilibria in the System

Ni3Al-Ni3Ta and Ni3Al-Ni3Ti”, High Temp. Alloys Cas Turbines Other Appl., Pt.2, Comm.

Euro. Communicates, Rep. EUR 10567, 955-964 (1986) (Equi. Diagram, Thermodyn., 8)

[1987Hil] Hilpert, K., Kobertz, D., Venugopal, V., Miller, M., Gerads, H., Bremer, F.T., Nickel, H.,



[1987Kha] Khadkikar, P.S., Vedula, K., “An Investigation of the Ni5Al3 Phase”, J. Mater. Res., 2(2),


[1988Bre] Bremer, F.J., Beyss, M., Karthaus, E., Hellwig, A., Schober, T., Welter, J.-M., Wenzl, H.,

“Experimental Analysis of the Ni-Al Phase Diagram”, J. Cryst. Growth, 87, 185-192 (1988)



Cr, Mn, Fe, Co, Ni, Cu, Zn)”, Int. Mater. Rev., 35, 293-327 (1990) (Crys. Structure, Equi.


[1991Bon] Bonneville, J., Martin, J.L., “The Strain Rate Sensitivity of Ni3(Al,Ta) Single Crystals”,

High-Temp. Ordered Intermetallic Alloys IV, Mater. Res. Soc. Symp. Proc., 213, 629-634

(1991) (Mechan. Prop., Experimental, 17)


Compositions in Nickel-Base Superalloys by Cluster Variation Method”, Calphad, 15,

143-158 (1991) (Equi. Diagram, Calculation, 34)

[1991Hay] Hayashi, T., Shinoda, T., Mishima, Y., Suzuki, T., “Effect of Off-Stoichiometry on the

Creep Behavior of Binary And Ternary Ni3Al”, High-Temp. Ordered Intermetallic Alloys

IV, Mater. Res. Soc. Symp. Proc., 213, 617-622 (1991) (Mechan. Prop., Experimental, 7)

[1991Kau] Kaufman, L., “Calculation of the Multicomponent Tantalum Based Phase Diagrams”,

Calphad, 15, 261-282 (1991) (Equi. Diagram, Calculation, 15)

[1991Mas] Maslenkov, S.B., Filin, S.A., Abramov, V.O., “Effect of Structural State and Alloying of

Transition Metals on the Degree of Hardening of Ternary Solid Solutions Based on Nickel

Monoaluminide”, Russ. Metall. (Engl. Transl.), (1), 115-118 (1991) (Mechan. Prop.,

Experimental, 10)


Ternary Systems”, Mater. Sci. Eng. A, 146, 123-130 (1991) (Equi. Diagram,

Experimental, 5)

[1991Nas] Nash, A., Nash, P., “The Ni-Ta (Nickel-Tantalum) System”, in “Phase Diagrams of Binary

Nickel Alloys, Monograph Series on Alloy Phase Diagrams”,Vol. 6, ASM-Intl., Materials

Park, Ohio, 320-325 (1991) (Equi. Diagram, Crys. Structure, Review, 38)

430


MSIT®

Al–Ni–Ta


NiAl”, Intermetal. Comp. - Struct. Mechan. Prop., Proc. Conf., 877-881 (1991) (Equi.

Diagram, Mechan. Prop., Abstract, 10)

[1991Zha] Zhao, J.T., Celato, L., Parthe, E., “Structure Refinement of Monoclinic 12-Layer TaNi3with -NbPt3 Type. New Crystallographic Descriptions of this Type and of the Nb3Rh5

Type Based on Smaller Unit Cells”, Acta Crystallogr., Sect. C: Crys. Struct. Commun., C47,


[1993Kha] Khadkikar, P.S., Locci, I.E., Vedula, K., Michal, G.M., “Transformation to Ni5Al3 in a 63.0

at.% Ni-Al Alloy”, Metall. Trans. A, 24A, 83-94 (1993) (Equi. Diagram, Crys. Structure,

Experimental, 28)

[1993Zak] Zakharov, A., “Aluminium - Nickel - Tantalum”, in MSIT Ternary Evaluation Program, in



Assessment, 28)

[1994Jia] Jia, C.C., Ishida, K., Nishizawa, T., “Partition of Alloying Elements Between (A1),

`(L12) and (B2) Phases in Ni-Al Base Systems”, Metall. Mater. Trans. A, A25, 473-485


[1994Joh] Johnson, D.R., Oliver, B.F., “Ternary Peritectic Solidification in the NiAl-Ni2AlTa-NiAlTa

System”, Mater. Lett., 20, 129-133 (1994) (Equi. Diagram, Experimental, 11)

[1996Du] Du, Y., Schmid-Fetzer, R., “Thermodynamic Modelling of the Al-Ta System”, J. Phase

Equilib., 17, 311-324 (1996) (Equi. Diagram, Crys. Structure, Thermodyn., Assessment,

Calculation, 55)

[1996Mac] Machon, L., Sauthoff, G., “Deformation Behavior of Al-Containing C14 Laves Phase

Alloys”, Intermetallics, 4, 469-481 (1996) (Crys. Structure, Experimental, 41)

[1996Pal] Palm, M., Sanders, W., Sauthoff, G., “Phase Equilibria in the Ni-Al-Ta System”,

Z. Metallkd., 87, 390-398 (1996) (Equi. Diagram, Crys. Structure, Experimental, 27)

[1996Pau] Paufler, P., Faber, J., Zahn, G., “X-Ray Single Crystal Diffraction Investigation on

Ni1+xAl1-x”, Acta Crystallogr., Sect. A: Found. Crystallogr., A52, C319 (1996) (Crys.

Structure, Experimental, Abstract, 3)

[1999Roc] da Rocha, F.S., Fraga, G.L.F., Brandao, D.E., da Silva, C.M., Gomes, A.A., “Specific Heat


Physica B (Amsterdam), B269, 154-162 (1999) (Thermodyn., Phys. Prop., Experimental,

Calculation, 20)

[1999Sun] Sung, P.K., Poirier, D.R., “Liquid-Solid Partition Ratios in Nickel-Base Alloys”, Metall.

Mater. Trans. A, A30, 2173-2181 (1999) (Equi. Diagram, Experimental, 41)


Morphologies in Ni-Al-X Ternary Diffusion Couples”, Mater. Sci. Eng. A, A312, 168-175

(2001) (Kinetics, Thermodyn., Experimental, 21)


Ni-Solid Solution in Ni-Al-X (X: V, Nb And Ta) Ternary Systems”, J. Phase Equilib., 22,



Intermetallic Compounds by Thermal Conductivity Measurements”, J. Mater. Res., 16,

2314-2320 (2001) (Thermal Conduct., Crys. Structure, Experimental, Calculation, 63)

[2003Cor] Cornish, L., Dolotko, O., Rogl, P., “Al-Ta (Aluminium - Tantalum)”, MSIT Binary


International Services GmbH, Stuttgart; submitted for publication (2003) (Crys. Structure,

Equi. Diagram, Assessment, 3)

[2003Sal] Saltykov, P., Cornish, L., Cacciamani, G., “Al-Ni (Aluminium - Nickel)”, MSIT Binary


International Services GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi.


431


MSIT®

Al–Ni–Ta


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

( Al) hP2

P63/mmc

Mg

a = 269.3

c = 439.8

at 25°C, 20.5 GPa [Mas2]

( Al)

< 660.452

cF4

Fm3m

Cu

a = 404.96 at 25°C [Mas2]

(Ta)

< 3020

cI2

Im3m

W

a = 330.30 at 25°C [Mas2]

, (Ni)

< 455

TaxNi1-x

Ni1-xAlxTaxNi1-x-yAly

cF4

Fm3m

Cu

a = 352.40

a = 357.8

a = 355.0

a = 355.3 to 357.5

a = 359.3

at 25°C [Mas2]

at 8.% Ta, 1150°C, linear

da/dx [1984Och2, 1985Mis]

at 14.% Al, linear da/dx [1984Och2]

at 7.5 - 10 at.% Ta, 75 - 80 at.% Ni,

1000°C, quenched, sample contained ´

and 3 [1979Nas]

at 10 at.% Ta, 80 at.% Ni, 1250°C,

quenched, sample contained 3

[1979Nas]

´, Ni3Al

< 372

Ta1-xNi3Alx

cP4

Pm3m

AuCu3

a = 356.77

a = 358.9

a = 359.0 to 362.4

[1986Hua]

at 63 at.% Ni [1993Kha]

at 3 - 12 at.% Ta, 58.6 - 80 at.% Ni, 1000

- 1200°C, multiphase samples quenched,

linear da/dx [1972Min, 1979Nas,

1984Och2, 1985Mis]

Ni5Al3 723

oC16

Cmmm

Pt5Ga3

a = 753

b = 661

c = 376

32 to 36 at.% Al at 63 at.% Ni

[1993Kha]

, NiAl

< 1638

Tax(Ni1-yAly)1-x

cP2 a)

Pm3m

CsCl

a = 286.0

a = 287

a = 288.72 0.02

a = 287.98 0.02

a = 286.6 to 296.8

42 to 69.2 at.% Ni [Mas2]

[1987Kha]

at 63 at.% Ni [1993Kha]

at 50 at.% Ni [1996Pau]

at 54 at.% Ni [1996Pau]

3.0 - 20 at.% Ta, 50 - 70 at.% Ni.

Quenched from 1250 - 1000°C.

Samples were multiphase. [1979Nas]

TaNi8< 307

(Ta1-xAlx)Ni8

tI36

NbNi8

a = 760.5

c = 358.5

a = 767

c = 348

at 11.1 at.% Ta [1991Nas]

at 11.8 at.%, 83.3 at.% Ni, from EMPA,

1250°C, quenched, [1979Nas]

432


MSIT®

Al–Ni–Ta

, TaNi3

< 1547

mP48 b)

P21/m

TaPt3

a = 452.3

b = 512.6

c = 2544

= 90°

at 22.5 to 28.5 at.% Ta [1991Nas]

[1991Zha] single crystal

TaNi2< 1404

tI6

I4/mmm

MoSi2

a = 315.4

c = 790.5

32.5 to 35 at.% Ta [1991Nas]

at 33.3 at.% Ta [1991Nas]

, TaNi

< 1570

Ta(Ni,Al)

hR13

R3m

W6Fe7

a = 492.1

c = 2690.5

a = 491.9 to 497.8

c = 2714 to 2735

a = 496.1

c = 2504

a = 428.3

c = 2649

50 to 54 at.% Ta [1991Nas]

at 50 at.% Ta [1991Nas]

50 - 55 at.% Ta, 35 - 23 at.% Ni

[1968Hun]

9 at.% Ta, 58.8 at.% Ni 1250°C,

quenched, alloy with , , 2 [1979Nas]

20 at.% Ta, 50 at.% Ni 1250°C,

quenched, alloy with 11 [1979Nas]

, Ta2Al

< 2061

tP30

P42/mnm

CrFe

a = 986.4

c = 521.5

at ~20 to 40 at.% Al [Mas2, V-C]

TaAl

< 1446

mP* [1996Du]

Ta5Al7< 1345

hP* [1996Du]

Ta2Al3< 1226

cF* [1996Du]

Ta39Al69

1548 - 1183

cF432

F43m

[1996Du]

TaAl3< 1608

tI8

I4/mmm

TiAl3

a = 383.7

c = 855.0

[V-C]

* 1, TaNiAl hP12

P63/mmc

MgZn2

a = 496.9

c = 798.5

a = 501.5

c = 817.1

[V-C] alloy 20 at.% Ta, 50 at.% Ni

1000°C, quenched, alloy with and 2

[1979Nas]

* 2, TaNi2Al cF16

Fm3m

BiF3

a = 594.9

a = 580 to 594

[V-C] 9 - 20 at.% Ta, 50 - 58.8 at.% Ni,

1000 - 1250°C, quenched. Multiphase

samples [1979Nas]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

433


MSIT®

Al–Ni–Ta

a) When quenching from 1250°C NiAl transformed to a body centered tetragonal martensite with a = 261.0 and

c = 337.6 pm [1979Nas].

b) As a result of heavy cold work TaNi3 transforms to “TaNi3-cw” with the TiAl3 type, tI8, a = 362.7, c = 745.5 pm

[1991Nas]. This form was obtained as phase by [1979Nas] when quenching Al-Ni-Ta alloys from 1250°C; the

lattice parameters in these alloys varied from a = 357.1 to 364.8 pm and c = 741.9 to 748.7 pm. “TaNi3” with

the TiCu3 type, oP8, a = 512.2, b = 452.2 and c = 423.5 pm was listed by [1991Nas] as a metastable phase due

to surface contamination TaNi3Ox. This phase was observed as phase in Al-Ni-Ta alloys when quenched from

1000°C [1979Nas]; the lattice parameters in these alloys varied from a = 509.4 to 512 pm, b = 437 to 452.7 and

c = 423 to 424.7 pm.


* 3, Ta0.5Ni3Al0.5

< 1393

hP16

P63/mmc

TiNi3

a = 510.5 to 513.7

c = 831.9 to 836.6

Al-rich [1965Gie, 1972Min, 1974Var,

1979Nas, V-C, 1984Wil, 1986Wil2]

* 4 cF96

Fd3m

NiTi2

a = 1150 [1965Ram] most probably metastable


Ta Ni Al

L + 1 ~1550 e1 (min) L

1

15.5

1.0

30.0

42.25

48.5

36.0

42.25

50.5

34.0

L 3 + 1387 e2 (max) L

3

16

15

22

75

75

75

9

10

3

L ´ + 3 1372 e3 (max) L

´

3

11

10

13.5

75

75

75

14

15

11.5

L + ´ + 3 ~1365 U1 L

´

3

13

11.5

6

13.5

72.5

73.5

71

74

14.5

15

23

12.5

L + ´ + 3 ~1360 U2 L

´

3

11.5

10

11

13

78.5

78.5

84

76

10

11.5

5

11

L + 3 + ~1360 U3 L

3

14

14.5

7

22

71.5

71.5

70

74

14.5

14

23

4

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

434


MSIT®

Al–Ni–Ta

Table 3: Equilibrium Compositions of and ´ Phases and Ta Partition Coefficient [1994Jia]

Table 4: Equilibrium Compositions of ' and Phases and Ta Partition Coefficient [1994Jia]

L + 3 + ~1355 E L

3

14.5

12.5

15

20

78.5

83.5

77.5

77

7

4

7.5

3

+ TaNi8 + 3 1330-1250 U4 - - - -

Temperature [°C] (at.%) ´ (at.%) Partition

coefficient kTa/ ´

Ta Al Ta Al

1300 0.09 19.9 0.25 22.9 2.78

0.21 19.4 0.56 22.3 2.67

0.86 18.4 1.52 20.8 1.77

1200 0.57 18.9 1.20 22.4 2.11

1100 0.23 15.9 0.68 19.9 2.96

0.40 15.4 1.23 19.2 3.08

1000 0.45 14.7 1.05 20.2 2.33

0.61 13.0 2.01 20.7 3.30

800 0.64 10.4 1.79 14.4 2.80

Temperature [°C] ´ (at.%) (at.%) Partition

coefficient kTa´/

Ta Al Ta Al

1300 0.9 23.0 0.39 34.0 2.31

1100 0.78 23.3 0.34 34.0 2.29

1000 1.44 27.2 0.09 40.0 -

2.64 25.9 0.26 39.3 10.2

900 0.74 27.2 0.39 34.3 1.90


Ta Ni Al

435


MSIT®

Al–Ni–Ta

l + γ γ´

1372 p1 L τ

3 + δ

1387 e2max

L + γ´ β + τ3

ca.1365 U1

l γ + δ1360 e

5

l γ´ + β1369 e

4

γ + δ TaNi8

ca.1330 p2

L γ´ + τ3

1372 e3max

L + γ´ γ + τ3

ca.1360 U2

L + τ3

β + δca.1360 U3

L γ + τ3 + δca.1355 E

γ + δ TaNi8 + τ

3<1330 U

4

L+β+τ3

L+γ+τ3

γ´+β+τ3

γ+γ´+τ3

τ3+β+δ

γ+τ3+δ

γ+TaNi8+τ

3 δ+τ3+TaNi

8

L+β+δ

Fig. 2: Al-Ni-Ta. Reaction scheme

Al-Ni Al-Ni-Ta Ni-Ta

L β + τ1

ca.1550 e1min

Ni, at.%

La

ttic

epa

ram

ete

r,p

m

485

40 30 20 10

790

50

500

495

490

505

800

810

820

830

Fig. 1: Al-Ni-Ta.

Lattice constants a

(triangle) and c

(circle) of the 1

phase

436


MSIT®

Al–Ni–Ta

10

20

80 90

10

20

Ta 30.00Ni 70.00Al 0.00

Ni

Ta 0.00Ni 70.00Al 30.00 Data / Grid: at.%

Axes: at.%

1450°C

1425°C

1400

°CMaximum solidsolubility

γ

Fig. 4: Al-Ni-Ta.

Partial solidus surface

10

20

30

40

50

50 60 70 80 90

10

20

30

40

50

Ta 60.00Ni 40.00Al 0.00

Ni


Axes: at.%

γ

γ'

β

δ

e5

E

U2

e3maxU

1U3

τ3

e2 m

ax

e4p

1

1500

1450

14001375

1450

1425

1400

1500

1450

1400

1375

e1min

10

20

30

40

50

50 60 70 80 90

10

20

30

40

50

Ta 60.00Ni 40.00Al 0.00

Ni


Axes: at.%

γ

γ'

β

δ

e5

E

U2

e3maxU

1U3

τ3

e2 m

ax

e4p

1

1500

1450

14001375

1450

1425

1400

1500

1450

1400

1375

e1min

Fig. 3: Al-Ni-Ta.

Partial liquidus

surface

437


MSIT®

Al–Ni–Ta

Ta, at.%

Te

mp

era

ture

,°C

0

1350

5 10 15 20

1400

1450

1500

18

7

6

54

2

3

1 Ni-Ta system

2 3 at.% Al

3 7 at.% Al

4 9 at.% Al

5 11 at.% Al

6 13 at.% Al

7 15 at.% Al

8 17 at.% Al

−−−−−−−−

Fig. 6: Al-Ni-Ta.

Dependence of the

phase liquidus (solid

lines) and solidus

(dash lines)

temperatures vs Ta

content at constant Al

contents

Al, at.%

Te

mp

era

ture

,°C

0

1350

5 10 15 20

1400

1450

1500

1

2

6

5

Maximum solid solubility

3

4

1 Al-Ni system

2 2 at.% Ta

3 4 at.% Ta

4 6 at.% Ta

5 8 at.% Ta

6 10 at.% Ta

−−−−−−

Fig. 5: Al-Ni-Ta.

Dependence of the

phase liquidus (solid

lines) and solidus

(dash lines)

temperatures vs Al

content at constant Ta

contents

438


MSIT®

Al–Ni–Ta

20

40

60

80

20 40 60 80

20

40

60

80

Ta Ni


Axes: at.%

β

γ'

γ

TaNi8δTaNi

2

µ

τ2

τ3

L

τ1

TaAl3

Ta39

Al69

Ta5Al

7

TaAl

Fig. 8: Al-Ni-Ta.


1250°C

10

20

80 90

10

20

Ta 25.00Ni 75.00Al 0.00

Ni


Axes: at.%

1327

°C

1227°C

1127

°C10

27°C

927°

C82

7°C

γ+γ´

γ

10

20

80 90

10

20

Ta 25.00Ni 75.00Al 0.00

Ni


Axes: at.%

1327

°C

1227°C

1127

°C10

27°C

927°

C82

7°C

γ+γ´

γ

Fig. 7: Al-Ni-Ta.

The ´/( + ´) solvus

isotherms

439


MSIT®

Al–Ni–Ta

10 20

1000

1100

1200

1300

1400

1500

1600

Ta 0.00Ni 75.00Al 25.00

Ta 25.00Ni 75.00Al 0.00Ta, at.%

Te

mp

era

ture

, °C

L

γ'

γ'+τ3

τ3

τ3+δ

δ

1550°C

10 20

1000

1100

1200

1300

1400

1500

1600

Ta 0.00Ni 75.00Al 25.00

Ta 25.00Ni 75.00Al 0.00Ta, at.%

Te

mp

era

ture

, °C

L

γ'

γ'+τ3

τ3

τ3+δ

δ

1550°CFig. 10: Al-Ni-Ta.

Experimentally

determined partial

vertical section

Ni3Al - TaNi3

20

40

60

80

20 40 60 80

20

40

60

80

Ta Ni


Axes: at.%

β

γ'

γ

TaNi8δTaNi

2

µ

τ2

τ3

L

TaAl

Ta2Al

3

τ1

Ta5Al

7

TaAl3

Ni2Al

3

Fig. 9: Al-Ni-Ta.


1000°C

440


MSIT®

Al–Ni–W

Aluminium – Nickel – Tungsten

Konstyantyn Korniyenko, Vasyl Kublii, Olga Fabrichnaya, Natalia Bochvar

Literature Data

Experimental investigations of the phase equilibria in the Al-Ni-W system are limited to the Ni-rich range

of compositions [1958Bud, 1978Gul, 1983Nas, 1986Nov, 1986Udo, 1991Udo, 1994Jia]. Solid solubility of

tungsten in the Ni3Al ( ´) phase is presented in [1966Arb, 1983Och, 1984Och1, 1984Och2, 1985Mis,

1985Nas] and the , (Ni) solvus surface is described in [1989Hon1, 1989Hon2] and [1991Mis]. [1991Sas]

studied the alloying effect of tungsten, on the solidification of the NiAl ( ) phase. [1958Bud] investigated

about 60 ternary alloys, that were prepared by high-frequency melting using corundum crucibles under a

protective layer of basic slag. The starting components were A-000 aluminium, nickel (99.9 mass%) and

tungsten (99.98 mass%). After subsequent stepwise annealing in high vacuum at 1200°C for 24 h, at 1000°C

for an additional 100 h and at 800°C for another 100 h, respectively, each batch of alloys was partly

quenched and cooled to room temperature. All losses were less than 0.2 to 0.5 mass% due to the fact that

the Al was introduced by means of premelted Al-Ni master alloys. Three isothermal sections of the Al-Ni-W

system in the Ni rich range at 1200, 1000 and 800°C and part of the NiAl-W pseudobinary section are

plotted. Since then the Ni-NiAl-W partial system has been reinvestigated frequently and partial isothermal

sections have been established by various research groups: at 1370°C [1986Udo], 1250°C [1983Nas],

1200°C [1986Udo, 1987Pri], 1150°C [1991Udo], 900°C [1991Udo] and for the temperature range from the

end of alloy solidification up to the beginning of solid state reactions, i.e. from ~1350 to 1000°C [1986Nov].

[1978Gul] recorded the liquidus and solidus temperatures in the Ni-rich corner by means of calibrated

thermocouples. [1986Nov] presented also the liquidus surface projection and reaction scheme for the partial

Ni-NiAl-W system.

In most cases the alloys were prepared by arc melting of the elements with 99.9 mass% minimal purity on

a water-cooled copper hearth under an argon atmosphere using a nonconsumable tungsten electrode

[1966Arb, 1983Nas, 1983Och, 1984Och1, 1984Och2, 1985Mis, 1985Nas, 1986Nov, 1986Udo, 1987Pri,

1989Hon1, 1989Hon2, 1991Mis, 1991Sas, 1991Udo], by induction melting under an argon atmosphere

[1994Jia] as well as by high-frequency melting [1986Nov, 1991Udo]. Due to rather sluggish reaction

kinetics, it was difficult to obtain equilibrium and homogenization treatments up to 500 h at various

temperatures were carried out. After heat treatment the samples usually were quenched to room

temperature. Methods of experimental investigation of the alloys were: X-ray diffraction [1966Arb,

1983Nas, 1983Och, 1984Och1, 1984Och2, 1985Mis, 1986Udo, 1986Nov, 1987Pri], metallography

[1983Nas, 1986Nov, 1986Udo, 1987Pri, 1991Sas, 1991Udo, 1994Jia], electron microprobe analysis

(EMPA) [1983Nas, 1986Nov, 1987Pri, 1989Hon1, 1989Hon2, 1991Mis, 1991Udo, 1994Jia], differential

thermal analysis (DTA) [1986Nov, 1989Hon1, 1989Hon2, 1991Mis] and [1991Udo] as well as

microhardness measurements [1987Pri]. The critical review of literature data on the phase equilibria in the

system was carried out by [1993Ale] within the MSIT Evaluation Program and is continued and updated by

the present evaluation.

Binary Systems

The descriptions of the Al-Ni, Al-W and Ni-W systems are accepted from [2003Sal], [2003Sch] and

[Mas2], respectively.

Solid Phases

No ternary phases have been found. Crystallographic data on the known unary and binary phases are listed

in Table 1. Based on earlier investigations of [1966Arb], the extent of the ´ phase field on alloying with

tungsten as well as the mode of atom substitution has been studied by [1983Och, 1984Och1, 1984Och2]

and [1985Mis] who arrived at a maximum solubility of less than 5 at.% W at 1000°C with W replacing Al

441


MSIT®

Al–Ni–W

in WxNi3Al1-x [1983Och, 1984Och1, 1984Och2, 1985Mis], whereas in general W was found to substitute

for both lattice sites in Wx(WyNi1-y)3Al1-x [1966Arb]. Data on solubility of W in the ´ phase are

contradictory and need more accurate definition. So, [1973Mul] and [2001Sav] prepared single crystals

WxNi3Al1-x with 3 at.% W replacing Al. A slightly larger solubility of 6 at.% W in ´ at 1000°C was

obtained by [1983Nas] on the basis of EMPA data, and the solubility of W in the ´ phase was found to

decrease from ~6 to ~4.5 at.% as the temperature increases from 1000 to 1250°C [1983Nas, 1985Nas]. But

[1987Pri] obtained the solubility of W in the ´ phase ~2 at.% at 1200°C, while [1991Udo] reported ~4 at.%

W at this temperature (Table 1). Similarly the solubility of W in the phase was shown to be ~0.2 at.% at

1250°C [1983Nas, 1985Nas], whereas [1958Bud] reported a value of 10 at.% at 1600°C, and 6 at.% at

1500°C. The value of W solubility of ~2 at.% at 1250°C, obtained by [1993Ale] from extrapolation of

[1958Bud] data, is ten times higher than the result of [1983Nas]. As a whole, it ought to agree with

judgement of [1987Pri] that interpretation of the obtained experimental data is very difficult because the ´

and phases crystal structures are the superstructures of fcc and bcc lattices, respectively. In the ternary

system in the ranges of their coexistence with W or Ni it is very difficult to fix the superstructure reflexes

at small relative amounts of intermetallic phases. The and ´ phases in the ranges of their existence are

very similar both by chemical compositions, lattice parameters, and by microhardness.


The partial pseudobinary section NiAl-W given by [1958Bud] has been corrected by [1993Ale] to account

for the results of [1983Nas] revealing a much lower solid solubility of W in NiAl (see Section “Solid

Phases”); the composition of the eutectic point is now at 1.4 at.% W. The pseudobinary section NiAl-W

in the range of compositions 0 to 10 at.% W is presented in Fig. 1 according to [1991Ale], with small

changes in the melting temperature of the phase as 1651°C [2003Sal], whereas [1993Ale] accepted

1640°C according to [Mas2] data.


Based on a theoretical analysis of the phase reactions in the Ni-NiAl-W ternary as well as from the

experimental investigation of six selected alloys, [1986Udo] reported two invariant equilibria: L+ + ´

at 1380 10°C and L + + ´ at 1340 15°C. The same reaction scheme has been constructed by

[1986Nov] based on the experimental liquidus surface, thereby confirming the phase triangulation of

[1983Nas] and superseding earlier results by [1968Bud]. The calculated liquidus surface and the calculated

isothermal section at 1323°C [1974Kau, 1975Kau], however, suggest two invariant equilibria: L+ + ´

and L + + at 1367°C which are different from [1986Udo]. The reason for this inconsistency could be

the differences in binary systems used by [1975Kau] and accepted by [1986Udo]. The other reason could

be that [1975Kau] used data of [1958Bud], which contradict recent experimental results. [1993Ale]

proposed a new reaction scheme (Fig. 2) based on more recent experimental data by [1991Udo] with some

adjustments of data of [1986Nov, 1986Udo, 1987Pri] to the Al-Ni phase diagram of [1987Hil], which is

essentially the same as assessed by [2003Sal]. The main feature of this reaction scheme is the change of the

character of invariant reactions due to the changes in character of l+ ´ reaction in the Al-Ni system.

According to [1993Ale] L+ + ´ is a peritectic reaction and L+ + ´ is a transition reaction. The

tentative compositions of the invariant equilibria that were derived are based on the boundary binary phase

diagrams, on the experimental isothermal section at 1250°C and on extrapolations of the solubility of W in

the ´ phase from 1000°C (6 at.%) through 1250°C (4.5 at.%) to 1400°C (3.5 at.%), see Table 2.


The liquidus surface shown in Fig. 3 contains four fields of primary crystallization: , , and ' and is

primarily based on the findings of [1986Nov] and on the data of [1978Gul] for the Ni-rich corner. With

respect to the adopted boundary systems the paths of the tie-triangle liquidus vertex and univariant curves

shown in Fig. 3 are adjusted to the reaction scheme (Fig. 2) and to the compositions of the invariant

equilibria (Table 2). The content of W in the liquid for the reaction L + is taken from the data of

442


MSIT®

Al–Ni–W

[1993Ale]. One of the main features of the liquidus surface is the relatively small area of primary

crystallization of the ' phase.

The (Ni) solvus was determined by [1991Mis] using DTA as major experimental technique. Chemical

analyses using energy-dispersive X-ray spectroscopy proved high accuracy of DTA results and provided

information on the phase relations such as the three-phase triangle neighboring the + ´ two-phase field.

The results on the determination of the solvus in Al-Ni-W system are presented in the form of solvus

isotherms in Fig. 4. The solvus is shown by solid lines connected to broken lines at clearly defined

inflections, which indicates the appearance of the three-phase equilibrium + + ´.

Isothermal Sections

[1958Bud] presented the isothermal section of the partial Ni-NiAl-W system at 1200, 1000 and 800°C. But

calculations of isothermal sections carried out by [1975Kau] for 1423, 1323, 1123 and 923°C, showed

essentially different phase equilibria, compared with those by [1958Bud]. [1975Kau] showed the existence

of the ´+ two-phase region, which was confirmed later by experimental results of [1983Nas, 1986Nov,

1986Udo, 1987Pri, 1989Hon1, 1989Hon2, 1991Eno, 1991Udo]. Besides that, [1958Bud] and [1986Nov]

showed a high solubility of tungsten in the phase (up to 10 at.% at 1600°C), but [1983Nas] established,

that it is not more than 0.2 at.% at 1250°C. These different values can not be reconlied but the data

[1983Nas] can be preferred because a direct method (EMPA) was used to determine the phase composition,

while [1958Bud] used indirect methods of DTA and X-ray diffraction. Results of [1994Jia] obtained by

diffusion couples method, also indicate a small solubility of tungsten in the phase, not higher than 0.53

at.% in the temperature range 1300 to 900°C. The isothermal section at 1250°C on the basis of [1983Nas]

data is presented in Fig. 5. The experimental phase compositions given in [1986Udo] and [1991Udo] are in

good agreement with this isothermal section. But according to the [1987Pri] data for 1200°C, the

three-phase + + ´ field is much narrower along the Al-content. Therefore further investigations of the

isothermal sections are necessary. The position of the phase boundary in Fig. 5 is slightly corrected

according to the accepted Ni-W binary system. The composition of the vertex of the + + ´ three-phase

field has been taken from [1989Hon1, 1989Hon2, 1991Mis] for 1227°C. The position of the boundary of

the ´ phase at the Ni side calculated by [1991Eno] is slightly shifted in the direction of increasing Ni

contents, compared with [1983Nas]. [1991Eno] employed the cluster variation method (CVM) which

utilizes the tetrahedron approximation and the phenomenological Lennard-Jones pair interaction potential.

Thermodynamics

Information on thermodynamic properties of the Al-Ni-W alloys is not quite complete. The thermodynamic

activities of Al in the ternary system for the Al-Ni0.9162W0.0838 section with aluminium content from 0 to

9 at.% have been determined by [1968Mal] using the emf method. The measurements were conducted at

temperatures of 772 and 907°C. The obtained values for the excess integral Gibbs energies and for the

activity coefficients at 772 and 907°C are presented in Table 3.


Al-Ni alloys with additions of a refractory metal, in particular tungsten, are interesting materials for the

production of in situ composites of eutectic superalloys. In spite of very complicated alloy compositions,

commercial superalloys generally consist mainly of two phases, namely, and ´. The phase has been

used for surface coating of the superalloys because of its high resistance against oxidation [1987Woo].

[1991Sas], studying the alloying effect of tungsten on the solidification of the phase, classified it as the

eutectic-phase containing compounds. The grain sizes appeared to correlate with the melting temperatures

of the compounds, similar to NiAl based ternary phases, formed by other elements. [1995Juj] investigated

the tensile properties of the ´ phase reinforced with continuous tungsten fibers. Model composites were

fabricated by isothermal forging of sandwiched tungsten fibers between boron-doped ´ plates at

temperatures from 1100 to 1200°C. It was found that the use of cold rolled ´ plates for hot forging enables

better consolidation and a lower forging temperature than the use of recrystallized ´ plates. Tensile test of

443


MSIT®

Al–Ni–W

the / ´ composites at ambient temperature to 1050°C reveals that the composites are stronger than

monolithic ´ alloys at the test temperatures. [2001Kai] studied the effect of tungsten on the morphological

stability of the interface between the ´ and phases using the Al-Ni-W ternary diffusion couples annealed

at temperatures ranging from 900 to 1300°C. Nonplanar interfaces with the Widmanstaetten-like structure

were formed in the couples. Measurements of electrical resistivity were used by [2001Sav] to study the

kinetics of ordering in the ´ phase. Investigations were conducted contactless levitating the sample by a

rotary magnetic field. The ordering and disordering processes are described by C-shaped time-temperature

transition diagrams. The activation energy values of the ordering and disordering are estimated using the

Arrhenius equation.

References

[1958Bud] Budberg, P.B., “Study of Alloys of the Ternary System Nickel-Aluminum-Tungsten”, Russ.

J. Inorg. Chem., USSR, 3, translated from Zh. Neorg. Khim., 3(3), 694-698 (1958) (Equi.


[1966Arb] Arbuzov, M.P., Kachkovskaya, E.T., Khayenko, B.V., “Structural X-Ray Diffraction Study

of the Compound Ni3Al Alloyed with Ti, Cr and W”, Russ. Met. Phys. Met. Sci., 21(6),

46-49 (1966), translated from Fiz. Met. Metalloved., 21(6), 854-857 (1966) (Crys. Structure,

Experimental, 15)

[1968Mal] Malkin, V.I., Pokidyshev, V.V., “The Effect of Alloying Elements on the Thermodynamic

Properties of Ni-Al Alloys” (in Russian), Sb. Tekhn. Trud. Nauchno-Issled. Inst. Chern.

Met., 59, 94-99 (1968) (Equi. Diagram, Experimental, Thermodyn., 4)

[1973Mul] Mulford, R.M., Pope, D.P., “The Yield Stress of Ni3(Al,W)”, Acta Metall., 21, 1375-1380

(1973) (Experimental, 24) as quoted by [1993Ale]

[1974Kau] Kaufman, L., Nesor, H., “Computer Calculated Phase Diagrams for the Ni-W-Al, Ni-Al-Hf,

Ni-Cr-Hf and Co(Cr,Ni)-Ta-C Systems”, Report No. NASA CR-134608, 55 (1974) (Equi.

Diagram, Theory, #, 28) as quoted by [1993Ale]

[1975Kau] Kaufman, L., Nesor, H., “Calculation of the Ni-Al-W, Ni-Al-Hf and Ni-Cr-Hf Systems”,

Can. Metall. Q., 14, 221-232 (1975) (Equi. Diagram, Theory, #, 22)

[1978Gul] Gulyaev, B.B., Grigorash, E.F., Efimova, M.N., “Investigation of Solidification Ranges of

Nickel Alloys” (in Russian), Metalloved. Term. Obrab. Metallov., 11, 34-37 (1978) (Equi.


[1983Nas] Nash, P., Fielding, S., West, D.R.F., “Phase Equilibria in Nickel-Rich Ni-Al-Mo and

Ni-Al-W Alloys”, Met. Sci., 17(4), 192-194 (1983) (Equi. Diagram, Experimental, 20)

[1983Och] Ochiai, S., Oya, Y., Suzuki, T., “Solubility Data in Ni3Al With Ternary Additions”, Bull.

P.M.E., 52, 1-17 (1983) (Equi. Diagram, Experimental, 7) as quoted by [1993Ale]

[1984Och1] Ochiai, S., Oya, Y., Suzuki, T., “Alloying Behaviour of Ni3Al, Ni3Ga, Ni3Si and Ni3Ge”,

Acta Metall., 32(2), 289-298 (1984) (Experimental, 90)

[1984Och2] Ochiai, S., Mishima, Y., Suzuki, T., “Lattice Parameter Data on Ni( ), Ni3Al( ') and

Ni3Ga( ') Solid Solutions”, Bull. P.M.E. (T.I.T.), 53, 15-28 (1984) (Crys. Structure,

Experimental, 66)

[1985Mis] Mishima, Y., Ochiai, S., Suzuki, “Lattice Parameters of Ni( ), Ni3Al( ') and Ni3Ga( ')

Solid Solutions with Additions of Transition and B-Subgroup Elements”, Acta Metall.,

33(6), 1161-1169 (1985) (Crys. Structure, Experimental, 64)

[1985Nas] Nash, P., “Nickel-Base Intermetallics for High Temperature Alloy Design”, High-Temp.

Ordered Intermet. Alloys, Boston, Mat. Res. Soc. Conf., 423-427 (1985) (Equi. Diagram,

Review, 15)

[1986Nov] Novikova, M.B., Budberg, P.B., “Phase State of Cast Alloys of Ni-NiAl-W System”, Russ.

Metall., (4), 407-111 (1986), translated from Izv. Akad. Nauk SSSR, Met., (4), 104-108


[1986Udo] Udovskii, A.L., Alekseeva, Z.M., Lukovkin, A.I., “Phase Equilibrium Diagram of the

Nickel-Aluminum-Tungsten System in the Range 1200-2000°C for the Concentration

444


MSIT®

Al–Ni–W

Region Ni-Ni0.5Al0.5W”, Sov. Phys., Dokl., 288(4), 496-499 (1986), translated from Dokl.

Akad. Nauk SSSR, 288(4), 935-939 (1986) (Equi. Diagram, Experimental, #, 13)

[1987Hil] Hilpert, K., Kobertz, D., Venugopol, V., Miller, M., Gerads, H., Bremer, F.J., Nickel, H.,

“Phase Diagram Studies of the Al-Ni System”, Z. Naturforsch., 42A, 1327-1332 (1987)

(Equi. Diagram, Experimental, #, 17)

[1987Pri] Prima, S.B., “The Isothermal Section of the W-Ni-Al Phase Diagram in the Range of

W-Ni-NiAl at 1200°C” (in Russian) in “Stabilnye i Metastabil'nye Fasy v Materialakh”,

Stable and Metastable Phases in Materials, Kiev, IPM, 97-105 (1987) (Equi. Diagram,

Experimental, #, 9)

[1987Woo] Wood, J.E., Goldman, E., in “Superalloys II”, Sims, C.T., Stoloff, N.S., Hagel, W.C. (Eds.),

New York, John Willey & Sons, 359-384 (1987) (Experimental) as quoted by [1994Jia]

[1989Hon1] Hong, Y.M., Nakajima, H., Mishima, Y., Suzuki, T., “The Solvus Surface in Ni-Al-X

(X: Cr, Mo and W) Ternary Systems”, I.S.I.J. International, 29(1), 78-84 (1989) (Equi.


[1989Hon2] Hong, Y.M., Mishima, Y., Suzuki, T., “Accurate Determination of ´ Solvus in Ni-Al-X

Ternary Systems”, Mat. Res. Soc. Symp. Proc., 133, 429-440 (1989) (Equi. Diagram,

Experimental, 35) as quoted by [1993Ale]


Compositions in Nickel-Base Superalloys by Cluster Variation Method”, Calphad, 15(2),

143-158 (1991) (Assessment, Calculation, Equi. Diagram, 34)

[1991Mis] Mishima, Y., Hong, Y., Suzuki, T., “Determination of the Solvus Surface in Ni-Al-X

Ternary Systems”, Mater. Sci. Eng. A, 146, 123-130 (1991) (Equi. Diagram,

Experimental, 5)


NiAl”, Intermetal. Comp. - Struct. Mechan. Prop., Proc. Conf., 877-881 (1991) (Abstract,


[1991Udo] Udovskii, A.L., Oldakovskii, I.V., Moldavskii, V.G., “Theoretical and Experimental

Investigations of Phase Equilibria in the Al-Ni-W System in the Range 900 to 1500°C” (in

Russian), Izv. Akad. Nauk. SSSR Met., 4, 112-123 (1991) (Equi. Diagram, Experimental)

[1993Ale] Alekseeva, Z.M., “Al-Ni-W (Aluminium - Nickel - Tungsten)”, MSIT Ternary Evaluation




[1994Jia] Jia, C.C., Ishida, K., Nishizawa, T., “Partition of Alloying Elements Between (A1), ´

(L12) and (B2) Phases In Ni-Al Base Systems”, Metall. Mater. Trans. A, 25, 473-485


[1995Juj] Jujur, I.N., Hanada, S., “Tensile Properties of W/Ni3Al Composites at Elevated

Temperatures”, Mater. Sci. Eng. A, 192/193, 848-855 (1995) (Equi. Diagram, Review, 20)


Morphologies in Ni-Al-X Ternary Diffusion Couples”, Mater. Sci. Eng. A, 312, 168-175

(2001) (Equi. Diagram, Experimental, Thermodyn., 21)

[2001Sav] Savin, O.V., Stepanova, N.N., Akshentsev, Yu.N., Rodionov, D.P., “Ordering Kinetics in

Ternary Ni3Al-X Alloys”, Scr. Mater., 45(8), 883-888 (2001) (Crys. Structure, Electr.

Prop., Experimental, Kinetics, 18)





[2003Sch] Schuster, J., “Al-W (Aluminium - Tungsten)”, MSIT Binary Evaluation Program, in MSIT


Stuttgart; submitted for publication, (2003) (Crys. Structure, Equi. Diagram,

Assessment, 22)

445


MSIT®

Al–Ni–W


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Al)

< 660.452

cF4

Fm3m

Cu a = 404.96

(Al) dissolves up to 0.1 at.% of Ni at 639.9°C

and up to 0.024 at.% W at 650°C

pure Al, T = 25°C [Mas2]

, WxNi1-x-yAly< 1455

WxNi1-x

cF4

Fm3m

Cu

a = 352.40

a = 356.35

a = 358.8

a = 357.4

a = 357.3

a = 357.0

0 x 0.11 0 y 0.2 at 1250°C [1983Nas]

0 y 0.203 at 1200°C [1987Pri]

0 x 0.16 y = 0 at 1002°C [Mas2]

0 x 0.175 at 1495°C [Mas2]

0 y 0.202 x = 0 at 1385°C [Mas2]

pure Ni, at 25°C [1984Och2, Mas2]

x = 0.09 [1985Mis], linear da/dx

x = 0.155 [1984Och2], linear da/dx scaled

from diagram

x = 0.05 y = 0.1783, annealed at 1200°C,

together with phase [1987Pri]

x = 0.053 y = 0.2018, annealed at 1200°C,


x = 0.0431 y = 0.2027, annealed at 1200°C,


, (W)

< 3422

cI2

Im3m

W

a = 316.52

dissolves up to 2.6 at.% Al [2003Sch];

up to 0.05 Ni at T = 1187 °C and 0.6 at.% Ni

at 1927°C [Mas2]

pure W, at 25°C [Mas2]

´, Ni3Al

< 1372

cP4

Pm3m

AuCu3 a = 356.6

a = 356.77

a = 356.32

a = 357.92

a = 357.30

a = 357.0

a = 357.8

a = 357.3

a = 357.6

a = 357.6

a = 357.0

a = 358.87

´ contains 73 to 76 at.% Ni [Mas2] and up to

4-6 at.% W [1983Nas, 1984Och1]

[2003Sal]

[2003Sal]

Disordered [2003Sal]

Ordered [2003Sal]

at 75 at.%Ni [1966Arb]

at 75 at.% Ni [1984Och2]

3W-75Ni (at.%) [1984Och2], linear da/dx

[1985Mis]

1.9W-73.38Ni (at.%), annealed at 1200C,


0.94W-74.12Ni (at.%), annealed at 1200°C,

together with and phases [1987Pri]

2W-74.12Ni (at.%), annealed at 1200°C,

together with and phases [1987Pri]

1.4 at.% W [1966Arb]

3 at.% W, 75 at.% Ni, annealed at 1227°C

(6 h) [2001Sav]

Ni5Al3< 723

oC16

Cmmm

Pt5Ga3

a = 744

b = 668

c = 372

63 to 68 at.% Ni [2003Sal, Mas2]

63 at.% Ni [2003Sal]

446


MSIT®

Al–Ni–W

, NiAl

< 1638

cP2

Pm3m

CsCl a = 287.04

a = 287.26

a = 286.0

a = 287.0

a = 288.72 0.02

a = 287.98 0.02

a = 286.6

a = 286.4

, NiAl contains 42 to 69.2 at.% Ni [Mas2]

0.2-1 at.% W [1983Nas, 1985Nas, 1993Ale]

57.7 at.% Ni [L-B]

46.6 at.% Ni [L-B]

[2003Sal]

63 at.% Ni [2003Sal]

50 at.% Ni [2003Sal]

54 at.% Ni [2003Sal];


together with ´ phase [1987Pri]


together with and ´ phase [1987Pri]

´, Ni3Al4< 702

cI112

Ia3d

Ni3Ga4

a = 1140.8 0.1 [2003Sal, V-C]

, Ni2Al3< 1138

hP5

P3m1

Ni2Al3

a = 402.8

c = 489.1

36.8 to 40.5 at.% Ni [Mas2]

[2003Sal, V-C]

, NiAl3< 856

oP16

Pnma

NiAl3

a = 661.3 0.1

b = 736.7 0.1

c = 481.1 0.1

[2003Sal]

Ni2Al9 mP22

P21/c

Ni2Al9

a = 868.5 0.6

b = 623.2 0.4

c = 618.5 0.4

= 96.50 0.5°

Metastable

[2003Sal]

NixAl1-x tP4

P4/mmm

AuCu

m**

a = 383.0

c = 320.5

a = 379.5

c = 325.6

a = 379.5

c = 325.6

a = 379.5

c = 325.6

a = 379.9 to 380.4

c = 322.6 to 323.3

a = 371.7 to 376.8

c = 335.3 to 339.9

a = 378.00

c = 328.00

a = 418

b = 271

c = 1448

= 90°

= 93.4°

= 90°

Martensite, metastable, 0.60 x 0.68

[2003Sal]

62.5 at.% Ni [2003Sal]

63.5 at.% Ni [2003Sal]

66.0 at.% Ni [2003Sal]

64 at.% Ni [2003Sal]

65 at.% Ni [2003Sal]

[2003Sal]

[2003Sal]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

447


MSIT®

Al–Ni–W

Ni2Al hP3

P3m1

CdI2

a = 407

b = 499

Metastable

[2003Sal]

D1 P105mc or

P105/mmc

a = 373.3

c = 407.3

[2003Sal]

D4 - - Decagonal, contained some Si [2003Sal]

WAl12

< 697

cI26

Im3

WAl12

a = 758.03 [2003Sch]

WAl6 mC56

C2/c

MoAl6

a = 514.4 0.3

b = 1298.9 0.5

c = 1348.4 0.4

= 94.03 0.39°

In the Al-B-W alloys rich in aluminium

[2003Sch]

WAl5< 870

hP12

P63

WAl5

a = 490.20

c = 885.70

[2003Sch]

WAl4< 1326

mC30

Cm

WAl4

a = 527.2

b = 1777.1

c = 521.8

= 100.2°

[2003Sch]

(Al-W)

h**

t**

a = 787.0 0.4

c = 2380 3

a = 714.5

c = 787.4

More rich in Al than WAl12, T < 580°C

[2003Sch]

T = 650°C [2003Sch]

[2003Sch]

´(Al-W) c** a = 692 8 Metastable (?), T = 200°C; transforms into

WAl12 [2003Sch]

cI* a = 766.4 [2003Sch]

(Al-W)

1300 < T < 1344

at 24 at.% W [2003Sch]

´(Al-W)

1317 < T 1420

at 30 at.% W [2003Sch]

´´(Al-W)

1335 < T 1650

at 33 at.% W [2003Sch]

W50Al50 t** a = 613

c = 418

Body-centered (?), metastable (?), from TEM

data [2003Sch]

WNi

1060

o**

MoNi

a = 776

b = 1248

c = 710

at 50 at.% W [V-C]

W2Ni tI96 a = 1040

c = 1090

at 66 at.% W [V-C]

WNi4 tI10

I4/m

MoNi4

a = 573 1

c = 355.3 0.1

at 20 at.% W [V-C]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

448


MSIT®

Al–Ni–W


Table 3: Integral Excess Gibbs Energies of Al-Ni-W Alloys and Activity Coefficients of Aluminium at 772

and 907°C [1968Mal]


Al Ni W

L + ~1400 U L

´

25.0

1.5

30.8

23.0

73.0

1.5

69.0

73.0

2.0

97.0

0.2

4.0

L + + ´ ~1405 P L

´

23.5

1.5

10.0

23.5

73.5

1.5

80.0

73.0

3.0

97.0

10.0

3.5

xAl fGes [J mol-1] lg Al

772°C 907°C 772°C 907°C

0,01

0.03

0.05

0.07

0.09

- 1710

- 3680

- 5580

- 7240

- 8860

-1550

-3430

-5180

-6810

-8440

-5.17

-4.79

-4.47

-4.22

-3.99

-4.34

-4.00

-3.76

-3.62

-3.52

1500

1600

1700

1800

W 10.00Ni 45.00Al 45.00

W 0.00Ni 50.00Al 50.00W, at.%

Te

mp

era

ture

, °C

1600°C

~1651°C

~1.4 ~1

β

Lα+L

α+β

2468

Fig. 1: Al-Ni-W.

Pseudobinary section

NiAl-W in the range

of compositions 0 to

10 at.% W

449


MSIT®

Al–Ni–W

Fig. 2: Al-Ni-W. Partial reaction scheme

l γ + α1495 e

2

L + α + γ γ´1405 P

l + γ γ´

1372 p

L + α β + γ´1400 U

l β + γ´

1369 e3

L β + α1600 e

1(max)

L + α γ´

α + β + γ´

α + γ + γ´

Ni-W Al-Ni-W Al-Ni

10

20

30

40

60 70 80 90

10

20

30

40

W 50.00Ni 50.00Al 0.00

Ni

W 0.00Ni 50.00Al 50.00 Data / Grid: at.%

Axes: at.%

γ

α

β

γ'p,1372

e3,1369

P

U

e2,1495

e1,1600

1486

1440

1637

1600

Fig. 3: Al-Ni-W.

Liquidus surface of

the Ni-rich region

450


MSIT®

Al–Ni–W

10

20

80 90

10

20

W 30.00Ni 70.00Al 0.00

Ni

W 0.00Ni 70.00Al 30.00 Data / Grid: at.%

Axes: at.%

1227°C

1127°C

1027°C

927°C

827°C

γ+γ´

γ

Fig. 4: Al-Ni-W.

The (Ni) solvus

surface

20

40

60

80

20 40 60 80

20

40

60

80

W Ni


Axes: at.%

γ

γ'

β

α+β+γ'

α+γ+γ'

α

α+β

α+γ

Fig. 5: Al-Ni-W.

Partial isothermal

section at 1250°C

451


MSIT®

Al–Ni–Zr

Aluminium – Nickel – Zirconium

Gautam Ghosh

Literature Data

[1966Mar1] was the first to report the isothermal section of the Al corner at 800°C. They prepared 99

ternary alloys, containing up to 25 at.% Zr and 75 at.% Ni, in an arc furnace under Ar atmosphere using

elemental metals of the following purity: AV-000 grade Al (99.997 mass%), very pure grade Ni (99.99

mass%) and iodide Zr (99.96 mass%). The alloys were annealed at 800°C for 700 h in evacuated quartz

tubes together with Zr foil followed by quenching in cold water. Later, [1969Bur] investigated the

constitutional equilibria of the Zr corner. The authors prepared 150 ternary alloys using metals of the above

mentioned purity and the alloys were annealed at 800°C for 2100 h followed by quenching in cold water.

In both cases phase analysis was performed by microstructural and X-ray diffraction techniques. The

constitutional equilibria of the Ni corner at 1100 and 1000°C were determined by [1983Jay1, 1983Jay2],

using metallography, electron microprobe analysis, and X-ray diffraction techniques. These results were

reviewed by [1991Nas] and [1993Gho]. Only a brief review of phase equilibria was presented by

[1990Kum].

Recently, Miura et al. [1999Miu, 2001Miu] investigated the solid-liquid phase equilibria of Ni-rich ternary

alloys using DTA, XRD and SEM-WDS analysis. [2001Miu] prepared three ternary alloys using 99.99

mass% Al, 99.95 mass% Ni, and 99.6 mass% Zr. [1991Mis] determined the solvus boundary of (Ni) using

DTA and SEM-EDX analysis.

Other recent investigations of the ternary system involve measurement of heat of formation of liquid phase

by calorimetry [1999Wit, 1999Zho], prediction of glass-forming ranges [2002Shi, 2003Shi], and synthesis

of amorphous alloys by liquid quenching/mechanical alloying, and study their crystallization behavior

[1990Bha, 1990Ino, 1995Gaf, 1997Kuh, 1998Sri, 1998Tur, 1999Hel, 2000Ill, 2001Cho, 2003Ele,

2003Yan].

Binary Systems

The Al-Ni binary phase diagram is accepted from [2003Sal] and the Al-Zr binary phase diagram is accepted

from [2003Sch]. Miura et al. [1999Miu, 2001Miu] determined the liquidus of Ni-rich alloys containing up

to 13 at.% Al. Unlike Hilpert et al. [1987Hil], Miura et al. [2001Miu] observed a maximum (1466°C) in the

liquidus at about 2 at.% Al. Except for [2001Miu], this feature has not been considered in the

thermodynamic modeling of the Al-Ni system [2003Sal]. The Ni-Zr binary phase diagram is accepted from

[1984Nas] in which the assessed temperature for invariant reaction L (Ni) + ZrNi5 is 1170°C. However,

recent experiment [2001Miu] shows that it occurs at 1196°C.

Solid Phases

The solid solubility of Ni in (Al) decreases from 0.11 at.% at 639.9°C to 0.01 at.% at 500°C. By rapid

solidification processing enhancement of the solid solubility is observed to as much as 7.7 at.% Ni [Mas2].

(Ni) dissolves both Al and Zr. The limit of solubility is given up to 21.3 at.% at the peritectic temperature

1372°C by [2003Sal]. Solution of Al in (Ni) causes a linear increase in the lattice parameter from 352.32

pm for pure Ni to 353.88 pm at 8 at.% Al [1985Mis]. Also, the limit of solubility of Zr 1.6 at.%, and this is

associated with a linear increase in the lattice parameter. The rate of increase in the lattice parameter, da/dc,

is reported to be 1.0 pm/at.% Zr [1984Och, 1985Mis]. Figure 1 shows the solubility isotherms of (Ni)

[1991Mis].

( Zr) can dissolve up to 25.5 at.% Al [Mas] and 2.92 at.% Ni [1984Nas], and the corresponding values for

( Zr) are 9.5 at.% Al [Mas] and 0.2 at.% Ni [1984Nas].

None of the binary intermediate phases is reported to dissolve more than 1 at.% of the third element

[1966Mar1, 1969Bur] at 800°C. Contrary to [1969Bur], [1971Bla] found that ZrNi5 can dissolve up to

452


MSIT®

Al–Ni–Zr

about 16.7 at.% Al at 800°C by replacing Ni and causing a linear increase in the lattice parameter to 674.8

pm at 16.7 at.% Zr; however, it is not known whether the stability of ZrNi4(Ni,Al) is enhanced by the

presence of oxygen. [1959Gua] reported that Ni3Al dissolves very small amount of Zr at 1150°C though no

specific value was given. On the other hand [1969Tho] found that, in directionally solidified Ni3Al-Zr2Ni7eutectic, Ni3Al dissolves about 2.7 at.% Zr as measured by an electron probe microanalyzer. In a review

paper [1985Nas] reported that Ni3Al can dissolve about 5 at.% Zr and Zr2Ni7 can dissolve about 11 at.%

Al, but the source of such information was not mentioned. Careful X-ray diffraction measurements

[1983Och, 1984Och, 1985Mis] showed, however, that Ni3Al can dissolve about 1 at.% Zr at 1000°C.

Solution of Zr in Ni3Al causes a linear increase in the lattice parameter, and the rate of increase in the lattice

parameter, da/dc, is 0.79 pm/at.% Zr [1984Och].

So far eight ternary phases have been reported. The ternary phase ZrNiAl ( 1) was first reported by

[1964Mar] and subsequently confirmed by [1966Mar2, 1967Kri, 1968Dwi, 1974Fer]. According to

[1968Dwi], the structure of the ZrNiAl phase can be better described by introducing a slight variation in

stacking sequence and by doubling the c-parameter.

Originally, [1964Sch, 1965Ram] reported that ZrNiAl has the Cu2Mg type structure with lattice parameter

a = 735 and 734 pm. Since subsequent investigations confirmed that ZrNiAl has the Fe2P type structure, the

ZrNiAl phase as designated by [1964Sch] and [1965Ram] is certainly the 2 phase (ZrNixAl2-x,

0.2 x 0.5) which has the Cu2Mg type structure and similar lattice parameter as confirmed by

[1966Mar2] and [1969Bur]. The lattice parameter of the 2 phase increases with increasing Al content

[1966Mar2]. The existence of ZrNi2Al ( 2) has been confirmed several times [1962Hei, 1964Mar, 1964Sch,

1965Ram]. [1962Hei] reported its structure to be of the CsCl type with a = 302.0 pm, but subsequent

investigations confirmed the structure to be of the MnCu2Al type with a = 609 to 612.3 pm [1964Mar,

1964Sch, 1965Ram, 1967Hof]. The ternary phase Zr2NiAl5 ( 5) was reported to be present in the as-cast

alloy [1965Ram], but was not reported by [1969Bur] in the 800°C isothermal section. Also, [1966Mar1] did

not observe the 5 phase in the annealed alloys (900°C for 700 h). Nevertheless, minor impurity levels can

significantly influence the stability of the AuCu3 type (or L12) phase [1990Kum]. Originally, ZrNi2Al5 ( 6)

was designated as Zr3Ni6Al16 [1969Bur]. The crystal structure of Zr5Ni4Al ( 7) is not known [1969Bur].

The details of the crystal structures and lattice parameters of stable solid phases are listed in Table 1.


Based on the DTA results and microstructure observations, [2001Miu] proposed that (Ni,Al) and ZrNi5form a pseudobinary eutectic (e2(max)). Even though the details are not known, the eutectic temperature

must be greater than 1196°C [2001Miu]. Ni3Al and Zr2Ni7 form a pseudobinary eutectic (e3(max)) at

1193°C [1969Tho]. This was also confirmed by [1978Hao] and [1983Jay1]. The composition of the eutectic

point was claimed by [1969Tho] to be at 10.9Zr-73.4Ni-15.7Al (at.%), however this can not be correct

because this point would appear outside of the Ni3Al-Zr2Ni7 section.

Solid solubilities of Zr in Ni3Al were also measured by electron microprobe and X-ray analyses after

annealing the samples at 1100 and 1000°C [1983Jay1, 1991Nas], and they were found to be 3.8 and 3.1

at.%, respectively. It is believed that Zr resides primarily on the Al-sublattice of Ni3Al [2001Ter]. The solid

solubility of Zr in Ni3Al, as measured by electron microprobe on dendrites adjacent to the eutectic, was 2.7

at.% [1969Tho].


Figure 2 shows the tentative reaction scheme for the solidification of Ni-rich alloys. Two saddle points,

e2(max) and e3(max), are due to [2001Miu] and [1969Tho], respectively. Since e2(max) feeds the binary

eutectic L (Ni)+ZrNi5 at 1196°C [2001Miu], it must occur above 1196°C. [2001Miu] observed both

(Ni)+ZrNi5 and Ni3Al+ZrNi5 microstructures in as-cast alloys. Based on these observations, they proposed

the existence of (Ni)+Ni3Al+ZrNi5 phase field which is the product of the ternary eutectic reaction E1

estimated to be occurring around 1186°C. It is important to note that the three-phase field,

(Ni)+Ni3Al+ZrNi5, has also been observed at 800°C [1966Mar1, 1969Bur]. The invariant reaction U2 gives

rise to a three-phase field Ni5Al+ZrNi5+Zr2Ni7 as proposed by [1969Tho] but it was not considered by

453


MSIT®

Al–Ni–Zr

[1983Jay1] and [1991Nas]. Other three ternary invariant reactions account for the observed microstructures,

but their temperatures of occurrence are not known.

Liquidus Surface

[1983Jay1] and [1991Nas] presented a speculative liquidus surface in the Ni-corner, which involves two

ternary eutectic reactions, two transition reactions and a pseudobinary eutectic reaction. They proposed a

ternary eutectic reaction L (Ni)+ZrNi5+Ni7Zr2 resulting in a three-phase field (Ni)+ZrNi5+Zr2Ni7 on

isothermal sections [1991Nas]. However, this conflicts with the observation of Ni3Al+ZrNi5 microstructure

by [2001Miu]. Therefore, the existence of ternary eutectic reaction L (Ni)+ZrNi5+Zr2Ni7 can be ruled out.

Based on the aforementioned results, the probable liquidus surface of the Ni-corner is shown in Fig. 3.

[1999Miu] determined the effect of Zr addition on the liquidus temperature of Ni-rich alloys. They found

that in alloys containing up to 13 at.% Al addition of Zr increases the liquidus temperature, while in alloys

containing more than 13 at.% addition of Zr decreases the liquidus temperature.

Isothermal Sections

Figures 4 and 5 show the partial isothermal sections in the Ni corner at 1100 and 1000°C [1983Jay1,

1991Nas], respectively. Some of the phase boundaries involving (Ni), Ni3Al, Zr2Ni7 and ZrNi5 are not

known exactly because of limited experimental work and also due to inconsistency between the X-ray data

and electron microprobe analysis. A number of amendments in Figs. 4 and 5 have been made to comply

with the accepted binary phase diagrams. [1969Tho] observed the Ni3Al-Zr2Ni7 eutectic microstructure in

the directionally solidified Ni-15Al-10.9Zr (at.%) alloy. Based on this observation, [1969Tho] proposed a

tentative phase diagram of the Ni corner characterized by the presence of a Ni3Al+ZrNi5+Zr2Ni7 phase field

and a Ni3Al+Zr2Ni7 phase field, which are consistent with the reaction scheme shown in Fig. 2. [1978Hao]

also observed Ni3Al-Zr2Ni7 as-cast eutectic microstructure in the above ternary alloy. Even though some

phase boundaries are shown dashed in Figs. 4 and 5, it is important to note that the above ternary alloy falls

in the Ni3Al+Zr2Ni7 phase field.

There are at least two reasons to doubt the results of [1971Bla] where it is claimed that ZrNi5 and ZrNi4Al

form a continuous solid solution. First, [1971Bla] noted a poor agreement between calculated and observed

X-ray intensities of AlNi4Zr which was attributed to preferred orientation due to cleavage parallel to (311)

plane, and the presence of oxygen in some positions in the structure. Secondly, [1983Jay1] reported that

ZrNi4Al is not a single-phase alloy.

Figure 6 shows the isothermal section at 800°C [1966Mar1, 1969Bur]. It should be noted that the solid

solubility ranges of Ni3Al, Zr2Ni7 and 2 are shown to be drastically reduced between 1000 and 800°C but

to be virtually unchanged between 1100 and 1000°C, which may be considered unlikely. The phase

relations of the Ni corner at 800°C as reported by [1969Bur] differ from that at 1000°C given by [1983Jay1].

Also, the observation of equilibrium between ZrAl3 and 3 phases at 780°C is inconsistent with the 800°C

isothermal section, but it would be consistent if one assumes a transition reaction 2+ 6 ZrAl3+ 3

occurring just below 800°C [1991Nas]. Similarly, the presence of transition reaction at higher temperatures

given in Fig. 2 is not necessarily in conflict with Fig. 6. The absence of phase fields 2+NiAl+Zr2Ni7,

Ni3Al+ZrNi5+Zr2Ni7 and Ni3Al+NiAl+Zr2Ni7 in the 800°C isothermal section can be due to various

solid-state reactions that may take place between 800 and 1000°C.

Thermodynamics

[1999Wit] determined the heat of formation of Al-Ni-Zr liquid alloys at 1292 5°C by high-temperature

calorimetry. Also, [1999Wit] used an empirical relationship for excess entropy, and derived the Gibbs

energy of mixing of liquid alloys. The heat capacity of undercooled Zr60Ni25Al15 liquid alloy was measured

by differential scanning calorimeter [1994Zap] and by an adiabatic calorimeter [1999Zho]. Both heat of

formation of liquid alloys and the heat capacity of undercooled liquid have been analyzed in terms of an

association model [1999Zho, 2000Kru]. A maximum in the capacity near the liquidus temperature was

454


MSIT®

Al–Ni–Zr

attributed to temperature dependent chemical short-range ordering [1999Zho]. Hoch [1995Hoc] proposed

an alternate model to describe the specific heat of undercooled liquid alloys.

The heat of formation of ternary amorphous alloys has been determined by solution calorimetry [1998Tur]

and direct reaction calorimetry [1999Hel]. These results also indicate the evidence of strong chemical

short-range ordering in amorphous alloys.

[1999Dar] measured the low-temperature (3.2 to 10.3 K) specific heat of ZrNi2Al ( 2) using an adiabatic

calorimeter, and analyzed the specific heat data in terms of electronic, Debye lattice and Einstein models.

The analysis of experimental data yields the Debye temperature D = 5°C. They also calculated the

electronic structure by tight-binding linearized muffin-tin orbital (TB-LMTO) method. Their results

underscore the importance of electron-phonon coupling on the phase stability.


[1998Sri] studied the microstructure and hardness of rapidly solidified Zr1Ni10Al89 alloy which was

subsequently aged at 150, 250, 350 and 450°C. The as-solidified ribbons, containing nanoscale precipitates

of NiAl3, exhibit hardness up to 4.5GPa. Aging treatment results in the precipitation of metastable cubic

ZrAl3, and the hardness decreases. They also observed that during aging the metastable ZrAl3, precipitates

in (Al) matrix and inside NiAl3 phase.

Miscellaneous

Existence of metastable phases have been reported in mechanically alloyed specimens, and also during

crystallization of amorphous alloys. Mechanical alloying of elemental powders produced an amorphous

phase having composition Al-12.5Ni-25Zr (at.%) (i.e. Zr2NiAl5 or 2), which crystallizes into ZrAl3 (L12)

and Zr6Ni8Al15 ( 3) upon heating to 780°C [1991Des] which is not the equilibrium state according to Fig. 6.

[1995Gaf] reported the formation of an fcc phase with lattice parameter of 460 pm in mechanically alloyed

Zrx(NiAl)1-x, 0.05 x 0.5. [2003Ele] observed the formation of two cubic phases during mechanical

alloying of Zr60Ni25Al15. The phase that forms first has lattice parameter of 1228.2 pm, and upon further

milling it transforms to another cubic phase with lattice parameter of 454.49 pm. They also found that

during crystallization both these phases give same end products, viz., Zr5Ni4Al ( 7) and Zr6NiAl2 ( 4);

however, the crystallization temperatures are different.

The structure, low-temperature specific heat [1987Yam] and crystallization behavior [1990Bha, 1997Kuh,

2001Cho] of some Al-Ni-Zr metallic glasses have also been reported. Mechanical alloying of elemental

powders produced an amorphous phase having composition Al-12.5Ni-25Zr (at.%) (i.e. Zr2NiAl5 or 2),

which crystallizes into ZrAl3 and Zr6Ni8Al15 ( 3) upon heating to 780°C [1991Des], which is not the

equilibrium state according to Fig. 6.

[1990Ino] reported the formation of amorphous alloys in the composition range of 3 to 67 at.% Ni and 0 to

37 at.% Al by melt spinning. In these alloys, the difference between glass transition temperature (Tg) and

crystallization (Tg) can be as large as 77°C. Also, the reduced glass transition temperature (Tg/Tm) can be

as high as 0.64. Shindo et al. [2002Shi, 2003Shi] have employed a quasi-chemical approach to predict the

critical composition range for bulk metallic glasses.

[1995Gaf] noted that mechanical alloying of (NiAl)1-xZrx, 0.05 x 0.5, does not yield fully amorphous

phase. They obtained up to 50% amorphous phase when x = 0.5. On the other hand, [1997Kuh] claimed to

obtain fully amorphous phase by mechanical alloying of Zr55Ni25Al20 and by rapid solidification of

Zr52Ni26Al22. Also, [1997Kuh] found that the observed phases in fully crystallized specimens do not

correspond to the expected equilibrium phases.

[2001Cho] studied crystallization of Zr55+xNi25Al20-x, 0 x 10, amorphous alloys prepared by rapid

solidification. Based on the change in crystallization temperature, they concluded that local environment of

atomic pairs is important for the stability of amorphous and supercooled liquid, and the retardation of

crystallization process. These amorphous alloys crystallize to Zr3Al2, (Zr) and ZrNi phases. Once again,

these do not correspond to the expected equilibrium phases in Fig. 6. A study of liquid-quenched

Zr60Ni25Al15 glassy alloy suggests the presence of high density of quenched-in nuclei [2003Yan].

455


MSIT®

Al–Ni–Zr

References

[1959Gua] Guard, R.W., Westbrook, J.H., “Alloying Behavior of Ni3Al ( ´ Phase)”, Trans. AIME,


[1962Hei] Heine, W., Zwicker, U., “Phases of the B2 Type (CsCl) in Ternary Systems Containing Cu

and Ni” (in German), Naturwissenschaften, 49, 391 (1962) (Crys. Structure,

Experimental, 1)


of MnCu2Al and MgZn2 Types Containing Al and Ga”, Sov. Phys.- Crystallogr., 9, 619-620

(1965), translated from Kristallografiya, 9, 737-738 (1964) (Crys. Structure,

Experimental, 4)

[1964Sch] Schubert, K., Raman, A., Rossteutscher, W., “Some Structure Data on Metallic Phases” (in

German), Naturwissenschaften, 51, 506 (1964) (Crys. Structure, Experimental, 0)

[1965Ram] Raman, A., Schubert, K., “On the Crystal Structure of Some Alloy Phases Related to TiAl3.

III. Investigations in Several T-Ni-Al and T-Cu-Al Alloy Systems” (in German),

Z. Metallkd., 56, 99-104 (1965) (Crys. Structure, Experimental, 14)

[1966Gan] Ganglberger, E., Nowotny, H., Benesovsky, F., “On Some New G-Phases” (in German),

Monatsh. Chem., 97, 219-220 (1966) (Crys. Structure, Experimental, 3)

[1966Mar1] Markiv, V.Ya., Matushevskaya, N.F., Rozum, N.S., Kuzma, Yu.B., “Investigation of

Al-Rich Zr-Ni-Al Alloys” (in Ukrainian), Izv. Akad. Nauk SSSR, Neorg. Mater., 2,


[1966Mar2] Markiv, V.Ya., Kripyakevich, P.I., “Compounds of the Type R(X'X")2 in Systems with

R= Ti, Zr, Hf; X'= Fe, Co, Ni, Cu; and X" = Al or Ga and Their Crystal Structures”, Sov.

Phys.- Crystallogr., 11, 733-738 (1967), translated from Kristallografiya, 11, 859-865


[1967Hof] Hofer, G., Stadelmaier, H.H., “Co, Ni and Cu Phases of the Ternary MnCu2Al Type” (in


[1967Kri] Kripyakevich, P.I., Markiv, V.Ya., Melnik, Ya.V., “Crystal Structure of Zr-Ni-Al,

Zr-Cu-Ga and Analogous Compounds” (in Ukrainian), Dopov. Akad. Nauk Ukr. RSR, (A8),


[1968Dwi] Dwight, A.E., Mueller, M.H., Conner, R.A., Downey, J.W., Knott, H., “Ternary

Compounds with the Fe2P-Type Structure”, Trans. TMS-AIME, 242, 2075-2080 (1968)


[1969Bur] Burnashova, V.V., Markiv, V.Ya., “A Study of the Zr-Ni-Al System” (in Ukrainian), Vest.

Lvov. Univ. (Khim.), (11), 34-37 (1969) (Equi. Diagram, Experimental, #, *, 8)

[1969Tho] Thompson, E.R., Lemkey, F.D., “Structure and Properties of Ni3Al( ') Eutectic Alloys

Produced by Unidirectional Solidification”, Trans. ASM, 62, 140-154 (1969) (Equi.


[1970Kri] Kripyakevich, P.I., Burnashova, V.V., Markiv, V.Ya, “Crystal Structures of the Compounds

Zr6FeAl2, Zr6CoAl2, Zr6NiAl2” (in Ukrainian), Dopov. Akad. Nauk Ukr. RSR, (A9),


[1971Bla] Blazina, Z., Ban, Z., “X-Ray Studies in the System ZrNi5-xAlx”, Croat. Chem. Acta., 43,


[1974Fer] Ferro, R., Marazza, R., Rambaldi, G., “Equiatomic Ternary Phases in the Alloys of the Rare

Earths with Indium and Nickel or Palladium”, Z. Metallkd., 65, 37-39 (1974) (Crys.


[1978Hao] Haour, G., Mollard, F., Lux, B., Wright, I.G., “New Eutectics Based on Fe, Co and Ni”,

Z. Metallkd., 69, 149-154 (1978) (Equi. Diagram, Experimental, 14)

[1982Mar] Markiv, V.Ya., Kripyakevich, P.I., Belyavina, N.M., “Crystal Structure of the Compound

ZrNi2Al5” (in Ukrainian), Dopov. Akad. Nauk Ukr. RSR, A3, 76-78 (1982) (Crys. Structure,

Experimental, 6)

456


MSIT®

Al–Ni–Zr

[1983Jay1] Jayanth, C.S., “Phase Equilibria in the Ni-Al-Zr and Ni-Al-V Systems”, M.S. Thesis, Illinois

Institute of Technology, (1983) (Equi. Diagram, Experimental, #, *, quoted in [1991Nas]

[1983Jay2] Jayanth, C.S., Nash, P., “Phase Equilibria in the Ni-Rich Region of the Al-Ni-Zr System”,

Bennett, L.H., Massalski, T.B., Giessen, B.C, (Eds.), Materials Research Society, Pittsburg,

PA, 395-398 (1983) (Experimental, 10)

[1983Och] Ochiai, S., Oya, Y., Suzuki, T., “Solubility Data of Ni3Al with Ternary Additions”, Bull.

P. M. E., (52), 1-17 (1983) (Equi. Diagram, Experimental, Review, 39)

[1984Nas] Nash, P., Jayanth, C.S., “The Ni-Zr (Nickel-Zirconium) System”, Bull. Alloy Phase

Diagrams, 5, 144-148 (1984) (Equi. Diagram, Review, #, *, 38)

[1984Och] Ochiai, S., Mishima, Y., Suzuki, T., “Lattice Parameter Data of Ni( ), Ni3Al( ') and

Ni3Ga( ') Solid Solutions”, Bull. P. M. E., (53), 15-28 (1984) (Crys. Structure,

Experimental, 66)

[1985Mis] Mishima, Y., Ochiai, S., Suzuki, T., “Lattice Parameters of Ni( ), Ni3Al( ') and Ni3Ga( ')

Solid Solutions with Additions of Taransition and B-Subgroup Elements”, Acta Metall., 33


[1985Nas] Nash, P., “Ni-Base Intermetallics for High-Temperature Alloy Design” in

“High-Temperature Ordered Intermetallic Alloys”, Koch, C.C., Liu, C.T., Stoloff, N.S.,

(Eds.), Materials Res. Soc., Pittsburg, PA, 423-427 (1985) (Equi. Diagram, Review, 15)

[1987Hil] Hilpert, K., Kobertz, D., Venugopal, V., Miller, M., Gerads, H., Bremmer, F.J., Nickel, H.,

“Phase Diagram Studies on the Al-Ni System”, Z. Naturforsch. A, 42A, 1327-1332 (1987)

(Equi. Diagram, Experimental, #, 17)

[1987Yam] Yamada, Y., Iton, Y., Mizutani, U., Shibagaki, N., Tanaka, K., “Low-Temperature Specific

Heat and Soft X-Ray Spectroscopic Studies of Ni33Zr67-Based Metallic Glasses Containing

H, B, Al and Si”, J. Phys. F: Met. Phys., 17, 2303-2311 (1987) (Experimental, 12)

[1990Bha] Bhatnagar, A.K., Rhie, K.W., Naugle, D.G., Wolfender, A., Zhang, B.H., Callaway, T.O.,

Bruton, W.D., Hu, C.-R., “The Effect of Simple Metal (Al, Ga) Addition on the

Crystallization and Density of Amorphous Zr-Ni Alloys”, J. Phys.: Condens. Matter., 2,

2625-2636 (1990) (Experimental, 24)

[1990Ino] Inoue, A., Zhang, T., Masumoto, T., “Zr-Al-Ni Amorphous Alloys with High Glass

Transition Temperature and Significant Supercooled Liquid Region”, Mater. Trans., JIM,

31, 177-183 (1990) (Experimental, 12)


Cr, Mn, Fe, Co, Ni, Cu, Zn)”, Int. Mater. Rev., 35, 293-327 (1990) (Equi. Diagram, Review,

#, 158)

[1991Des] Desch, P., Schwarz, R.B., Nash, P., “Formation of Metastable L12 Phases in Al3Zr and

Al-12% X-25% Zr (X = Li, Cr, Fe, Ni, Cu)”, J. Less-Common Met., 168, 69-80 (1991)

(Experimental, 25)

[1991Mis] Mishima, Y., Hong, Y.M., Suzuki, T., “Determination g Solvus Surface in Ni-Al-X Ternary

Systems”, Mater. Sci. Eng., A146 123-130 (1991) (Equi. Diagram, Experimental, #, *, 5)

[1991Nas] Nash, P., Pan, Y.Y., “The Al-Ni-Zr System (Aluminum-Nickel- Zirconium)”, J. Alloy

Phase Equilibria, 12, 105-113 (1991) (Equi. Diagram, Review, #, *, 49)

[1993Gho] Ghosh, G., “Aluminium-Nickel-Zirconium”, in MSIT Ternary Evaluation Program, in



Assessment, 35)

[1994Zap] Zappel, B., Sommer, F., “Structural Enthalpy Relaxation in the Glass Transition Range”,

Mater. Sci. Eng. A, A179-180, 283-287 (1994) (Thermodyn., 10)

[1995Gaf] Gaffet, E., “Structural Investigation of Mechanicall Alloyed (NiAl)1-x(M)x (M=Fe,Zr)

Nanocrystalline and Amorphous Phases”, Nano-Structured Mater., 5(4), 393-409 (1995)


[1995Hoc] Hoch, M., “The Heat Capacity Cp of Undercooled Liquid Metals and Alloys”, Z. Metallkd.,

86(8), 557-560 (1995) (Theory, Thermodyn., 14)

457


MSIT®

Al–Ni–Zr

[1997Jou] Joubner, J.-M., Cerny, R., Yvon, K., Latroche, M., Persheron-Guegan, A.,

“Zirconium-Nickel, Zr7Ni10: Spase Group Revision for the Stoichiometric Phase”, Acta

Crystallogr., Sect. C; Cryst. Struct. Commun., C53(11), 1536-1538 (1997) (Crys. Structure,

Experimental, 12)

[1997Kuh] Kuhnast, F.A., Held, O., Ragnier, F., Illekova, E., “Calorimetric and Structural Analyses of

Mechanically Alloyed and Rapidly Quenched Zn-Ni-Al Alloys”, Mater. Sci. Eng. A,

A226-228, 463-467 (1997) (Crys. Structure, Equi. Diagram, Experimental,

Thermodyn., 10)

[1998Sri] Srinivasan, D., Chattopadhyay, K., “Formation and Coarsening of a Nanodispersed

Microstructure in Melt Spun Al-Ni-Zr Alloy”, Mater. Sci. Eng. A, A255, 107-116 (1998)


[1998Tur] Turchanin, A.A., Tomilin, I.A., “Experimental Investigations of the Enthalpies of

Formation of Zr-Based Metallic Amorphous Binary and Ternary Alloys”, Ber. Bunsen-Ges.

Phys. Chem., 102(9), 1252-1258 (1998) (Experimental, Thermodyn., 29)

[1999Dar] Da Rocha, F.S., Fraga, G.L.F., Brandao, D.E., Da Silva, C.M., Gomes, A.A., “Specific Heat


Physica B (Amsterdam), 269, 154-162 (1999) (Crys. Structure, Experimental, Theory,

Thermodyn., 20)

[1999Hel] Held, O., Braganti, J.P., Kuhnast, F.A., “Calorimetric and Structural Analysis of the New

Phase Al33Ni16Zr51 Produced by Direct Synhesis and Mechanical Alloying”, J. Alloys

Compd., 290, 197-202 (1999) (Experimental, Thermodyn., 15)


Ni-Solid Solution in Ni-Al-X (X: Ti, Zr, and Hf) Ternary Systems”, J. Phase Equilib.,

20(3), 193-198 (1999) (Equi. Diagram, Experimental, #, *,11)

[1999Wit] Witusiewicz, V.T., Sommer, F., “Thermodynamics of liquid Al-Ni-Zr and Al-Cu-Ni-Zr

Alloys”, J. Alloys Compd., 289, 152-167 (1999) (Experimental, Thermodyn., 11)

[1999Zav] Zavaliy, I. Yu., Pecharsky, V.K., Miler, G.J., Akselrud, L.G., “Hydrogenation of Zr6MeX2

Intermetallic Compounds (Me=Fe, Co, Ni; X=Al, Ga, Sn): Crystallographic and Theoretical

Analysis”, J. Alloys Compd., 283, 106-116 (1999) (Crys. Structure, Experimental, 31)

[1999Zho] Zhou, S.H., Sommer, F., “Calometric Study of Liquid and Undercooled Liquid Al-Ni-Zr

Alloys”, J. Non-Cryst. Solids, 250-252, 572-576 (1999) (Calculation, Experimental,

Thermodyn., 14)

[2000Ill] Illekova, E., Jergel, M., Kuhnast, F.-A., “On Structural and Thermal Relaxation in

Non-Crystalline Zr-Ni-Al Alloys”, Mater. Sci. Eng. A, A278, 27-35 (2000) (Crys. Structure,

Experimental, 32)

[2000Kru] Krull, H.G., Singh, R.N., Sommer, F., “Generalized Association Model”, Z. Metallkd.,

91(5), 356-365 (2000) (Review, Thermodyn., 46)

[2001Cho] Choi, H.W., Cho, J.H., Kim, J.E., Kim, Y.H., Yang, Y.S., “Calorimetric and Structural

Properties of Amorphous Zr-Al-Ni Alloys”, Scr. Mater., 44(8-9), 2027-2030 (2001) (Crys.

Structure, Equi. Diagram, Experimental, Thermodyn., 13)

[2001Miu] Miura, S., Unno, H., Yamazaki, T., Takizawa, S., Mohri. T., “Reinvestigation of Ni-Solid

Solution/Liquid Equilibria in Ni-Al Binary and Ni-Al-Zr Ternary Systems”, J. Phase

Equilib., 22, 457-462 (2001) (Equi. Diagram, Experimental, #, *, 9)




[2002Shi] Shindo, T., Waseda, Y., Inoue, A., “Prediction of Glass-Forming Ranges in Zr-Ni-Al

Alloys”, Mater. Trans., 43, 2502-2508 (2002) (Thermodyn., Theory, 25)

[2003Ele] El-Eskandarany, M.S., Saida, J., Inoue, A., “Structural and Calorimetric Evolutions of

Mechanically-Induced Solid-State Devitrificated Zr60Ni25Al15 Glassy Alloy Powder”, Acta


458


MSIT®

Al–Ni–Zr







GmbH, Stuttgart; submitted for publication (2003) (Crys. Structure, Equi. Diagram,

Assessment, 151)

[2003Shi] Shindo, T., Waseda, Y., Inoue, A., “Prediction of Critical Compositions for Bulk Glass

Formation in La-Based, Cu-Based and Zr-Based Ternary Alloys”, Mater. Trans., 44,

351-352 (2003) (Thermodyn., Theory, 20)

[2003Yan] Yan, Z., Li, J., He, S., Wang, H., Zhou, Y., “Study of the Crystallization Kinetics of

Zr60Ni25Al15 Glassy Alloy by Differential Scanning Calorimetry”, Mater. Trans., 44,

709-712 (2003) (Experiemntal, 17)


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Al)

660.452

cF4

Fm3m

Cu

a = 404.88 pure Al at 24°C [V-C]

(Ni)

1455

cF4

Fm3m

Cu

a = 352.32 pure Ni at 20°C [V-C]

( Zr)(h)

1855 - 863

cI2

Im3m

W

a = 356.8 [V-C]

( Zr)(r)

< 863

hP2

P63/mmc

Mg

a = 323.2

c = 514.7

[V-C]

NiAl3 854

oP16

Pnma

NiAl3

a = 661.3

b = 736.7

c = 481.1

[2003Sal]

Ni2Al3 1133

hP5

P3m1

Ni2Al3

a = 402.8

c = 489.1

[2003Sal]

58.7 to 63.9 at.% Al

NiAl

1638

cP2

Pm3m

CsCl

a = 286.00

to 288.72



Ni5Al3 700

oC16

Cmmm

Pt5Ga3

a = 753.0

b = 661.0

c = 376.0



Ni3Al

1372

cP4

Pm3m

AuCu3

a = 356.77

to 358.90



ZrAl3 1580

cP4

Pm3m

Cu3Au

a = 399.93

c = 1728.3

[2003Sch]

459


MSIT®

Al–Ni–Zr

ZrAl21660

hP12

P63/mmc

MgZn2

a = 528.24

c = 874.82

[2003Sch]

Zr2Al31590

oF40

Fdd2

Zr2Al3

a = 960.1

b = 1390.6

c = 557.4

[2003Sch]

ZrAl

1275 25

oC8

Cmcm

CrB

a = 335.9

b = 1088.7

c = 427.4

[2003Sch]

Zr5Al4 (h)

1550 - 1000

hP18

P6/mcm

Ti5Ga4

a = 844.8

c = 580.5

[2003Sch]

Zr4Al3< 1030

hP7

P6/mmm

Zr4Al3

a = 543.3

c = 539.0

[2003Sch]

Zr3Al2< 1480

tP20

P42/mnm

Zr3Al2

a = 763.0

c = 699.8

[2003Sch]

Zr5Al3 (h)

< 1400

tI32

I4/mcm

W5Si3

a = 1104.4

c = 539.1

[2003Sch]

Zr5Al3 (r) hP16

P63/mcm

Mn5Si3

a = 817.4

c = 569.8

[2003Sch]

Zr2Al

< 1350

hP6

P63/mmc

Ni2In

a = 489.39

c = 592.83

[2003Sch]

Zr3Al

< 1019

cP4

Pm3m

AuCu3

a = 437.2 [2003Sch]

ZrNi5 1300

cF24

F43m

AuBe5

a = 670.64

to 670.72

[1984Nas], 15.0 to 18.0 at.% Zr

Zr2Ni7 1440

mC36

C2/m

Zr2Ni7

a = 469.8 0.9

b = 823.5 1.3

c = 1219.3 1.6

= 95.83°

[V-C]

ZrNi3 920

hP8

P63/mmc

Ni3Sn

a = 530.9

c = 430.3

[1984Nas], 24.5 to 26.0 at.% Zr

Zr8Ni21

1180

aP29

P1

Hf8Ni21

a = 647.21

b = 806.45

c = 858.75

= 75.18°

= 68.00°

= 75.20°

[1984Nas]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

460


MSIT®

Al–Ni–Zr

Zr7Ni10

1160

oC68

Cmca

Zr7Ni10

a = 1238.1 1.0

b = 918.5 0.5

c = 922.1 1.1

[1997Jou],

measured on single crystal with 799

refections

Zr9Ni11

978 - 1170

tI40

I4/m

Zr9Pt11

a = 988.0

c = 661.0

[1984Nas]

ZrNi

1260

oC8

Cmcm

CrB

a = 326.8

b = 990.3

c = 410.7

[V-C]

Zr2Ni

1120

tI12

I4/mcm

CuAl2

a = 647.7

to 648.3

c = 524.1

to 526.7

[V-C]

* 1, ZrNiAl hP9

P62m

Fe2P

a = 691.57

c = 694.12

a = 692.1

c = 346.7

[1968Dwi]

[1974Fer]

* 2, ZrNi2Al cF16

Fm3m

MnCu2Al

cP2

Pm3m

CsCl

a = 611.47

a = 302.0

[1999Dar]

[1962Hei]

* 3, Zr6Ni8Al15 cF116

Fm3m

Th6Mn23

a = 1208.0 [1966Gan, 1966Mar1]

* 4, Zr6NiAl2 hP9

P62m

Zr6CoAl2

a = 792.0

c = 334.0

a = 792.8

c = 334.7

[1969Bur, 1970Kri]

[1999Zav]

* 5, Zr2NiAl5 cP4

Pm3m

AuCu3

a = 406.0 [1964Sch, 1965Ram], observed in

as-cast alloy

* 6, ZrNi2Al5 tI16

I4/mmm

ZrNi2Al5

a = 402.3

c = 1444.0

a = 401.0

c = 1441

[1982Mar]

[1969Bur]

* 7, Zr5Ni4Al - - [1969Bur]

* 2, ZrNixAl2-x cF24

Fd3m

Cu2Mg

a = 734.3

to 746.4

a = 746.4

a = 734.3

0.2 x 0.5 [1966Mar1,

1966Mar2]

at x = 0.2 [1966Mar1]

at x = 0.5 [1966Mar1]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

461


MSIT®

Al–Ni–Zr

10

90

10

Zr 20.00Ni 80.00Al 0.00

Ni

Zr 0.00Ni 80.00Al 20.00 Data / Grid: at.%

Axes: at.%

1127°C

1027

927

827

γ

Fig. 1: Al-Ni-Zr.

Solubility isotherms

of (Ni)

Fig. 2: Al-Ni-Zr. A tentative reaction scheme for the solidification of Ni-rich alloys

Al-Ni A-B-C

l+(Ni) Ni3Al

1372 p1

L+NiAl+τ2

?

Al-Ni-Zr

L+τ2

NiAl+Zr2Ni

7? U

1

Ni-Zr

l+Zr2Ni

7ZrNi

5

1300 p3

l Ni3Al+NiAl

1369 e1

l (Ni)+ZrNi5

1170 e4

L+τ2+Zr

2Ni

7

?

L (Ni)+ZrNi5

>1196 e2

L Ni3Al+Zr

2Ni

7

1193 e3

L+Zr2Ni

7Ni

3Al+ZrNi

5? U

2

L (Ni)+Ni3Al+ZrNi

51150 E

1

L Ni3Al+NiAl+Zr

2Ni

7? E

2

τ2+NiAl+Zr

2Ni

7

L+Ni3Al+ZrNi

5 Ni3Al+ZrNi

5+Zr

2Ni

7

(Ni)+Ni3Al+ZrNi

5

Ni3Al+NiAl+Zr

2Ni

7

L+NiAl+Zr2Ni

7

462


MSIT®

Al–Ni–Zr

20

40

60

80

20 40 60 80

20

40

60

80

Zr Ni


Axes: at.%

τ2

U1

NiAl

Ni3Al

e3(max)

Zr2Ni

7

(Ni)

e2(max)

ZrNi5

Zr8Ni

21

E2

E1

e1p

1

U2

p3 e

4

10

20

30

40

50

50 60 70 80 90

10

20

30

40

50

Zr 60.00Ni 40.00Al 0.00

Ni


Axes: at.%

Zr2Ni

7

ZrNi5

(Ni)

Ni3Al

NiAl

τ2

Fig. 3: Al-Ni-Zr.

A tentative liquidus

surface of the

Ni-corner

Fig. 4: Al-Ni-Zr.

Partial isothermal

section at 1100°C.

The dashed lines


phase boundaries

463


MSIT®

Al–Ni–Zr

20

40

60

80

20 40 60 80

20

40

60

80

Zr Ni


Axes: at.%L

NiAl3

Ni2Al

3

NiAl

Ni3Al

(Ni)

ZrNi5

Zr2Ni

7ZrNi3

Zr8Ni

21Zr

7Ni

10ZrNiZr2Ni(Zr)

Zr3Al

Zr2Al

Zr3Al

2

Zr4Al

3

ZrAl

Zr2Al

3

ZrAl2

ZrAl3

τ6

λ2

τ3

τ1

τ4

τ7

τ2

10

20

30

40

50

50 60 70 80 90

10

20

30

40

50

Zr 60.00Ni 40.00Al 0.00

Ni


Axes: at.%

Zr2Ni

7

ZrNi5

(Ni)

Ni3Al

NiAl

τ2

Fig. 6: Al-Ni-Zr.


800°C

Fig. 5: Al-Ni-Zr.

Partial isothermal

section at 1000°C.

The dashed lines


phase boundaries

light metal systems. part 3: selected systems from al-fe-v to al-ni-zr

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