ISI J International. Vol. 29 (1989), No. 5, pp. 353-377Review

Development of ShapeMemoryAlloys

Shuichi MIYAZAKIand Kazuhiro OTSUKAInstitute of Materials Science, University of Tsukuba, Tsukuba, Ibaraki-ken, 305 Japan.

(Received on December19. 1988,• accepted in the final form on January 20. 1989)

Recent developmentof shapememoryalloys is reviewed, emphasisbeing placed on the Ti-Ni. Cu~)asedand ferrous

alloys which are considered as practical materials for applications amongmanyshapememoryalloys. Crystal structures

of the parent and martensitic phases are described, and the crystallography of the martensitic transformations is also

briefly explained Theorigins of the shapememoryeffect and the shapememorymechanismsare discussed on the basis

of the crystal structure and the crystallography of the martensitic transformations. Since an applied stress also inducesthe martensitic transformations, successive stages of the martensitic transformations are reviewed briefly in Cu-basedandTi-Ni alloys, which showmartensite-to-martensite transformations upon loading. Then, characterization of the shape

memoryalloys are reviewed in detail; i.e.,phasediagrams, transformation temperatures, transformation process, stress-

induced transformation, aging effects, cycling effects, fracture, fatigue, grain refinement, and so on

KEYWORDS:shape memoryeffect; pseudoelasticity; superelasticity; martensite; martensitic transformation; R-phase;rhombohedral phase; Ti-Ni, Ni!Ti; Cu-Al-Ni; Cu-Zn-AI; ferrous alloys; shape memoryalioy.

l .Introduction

The shape memoryeffect (SlviE) appears in somespecial alloys which sho_w crystallographically revers-ible martensitic transfbrmations. The martensitictransformation is accompaniedby a large shear-like

deformation associated with a diffusionless structural

change; the deformation generally amounts to about20 times more than the elastic deformation. Themartensite is deformable and it mayalso be inducedfrom the parent phasc by loading) both defbrmationmodesbeing associated with no permanent strain in

the shape memoryalloys (SMAS)' Thus, a large de-

formation induced in the SMAScan recover perfectly

by heating to temperatures above the reverse-trans-

formation flnish temperature (Af) after unloading(shape memoryeffect or SME)or simply by unload-ing at temperatures above Af (pseudoelasticity orsuperelasticity (PE)).

The SMEwas first found in a Au-Cd alloy in

1951,1) and then in an In-TI alloy in 1953.2,3) How-ever, the possibility for using the SMEin actual ap-plications wasrealized after the SMEwasfound in aTi-Ni alloy in 1963.4) Since the Ti-Ni alloy has

manycomplicated features and a difficulty in makingsingle crystals, the basic understanding of this alloy

wasnot possiblc until the carly 1980s. Onthe otherhand, a Cu-Al-Ni alloy wasalso found to reveal the

SMEin 19645) and in this alloy it wasdemonstratedthat the SMEis closely related to the thermo_elastic

martensitic transformation.6) Since then, the basic

understanding of the origin of the SME,6) the shape

memorymechanism,7~9) the crystallography of thestress-induced martensitic transformationro,11] hasbeen established for Clu-based alloys in the 1970s.

This is partly due to an easiness of making single

crystals of these alloys.

However, the Ti-Ni alloy has been the mostimportant material for applications, because the Cu-based alloys are brittle in a polycrystalline state.12)

In the 1980s, the basic understanding of the thermo-mechanical treatment,13,14) the deformation behav-ior,15-20) the shape memorymechanism,21,22, the na-ture of the R-phase,23~31] the crystallography32~36) ofboth the R-phaseand the martensitic transformations,

and the phase diagram37) including the crystal struc-

ture of metastable precipitates37~40) has been obtainedfor the Ti-Ni alloys.

In the same 1980s, manypatents of applications

using the SMEand PEhave been applied for; nowthe total numberof the applied patents amounts to

more than 4OO0.41) At this stage> the price of theshape memoryalloys is also one of the key factors for

applications. Recently, it was also found that someferrous alloys also exhibit perfect SMEunder certain

conditions.42~45) Therefore, many efforts are nowbeing devoted to developing Cu-basedalloys and fer-

rous alloys as well as Ti-Ni alloys.

This paper will first review the fundamental aspectsof SMAs,i.e., the crystal structures, the crystallogra-

phy of martensitic transformations, origins of SME,shape memorymechanism, successive stress-induced

transformation and so on. And then most part ofthis paper will be devoted to reviewing the recent

developmentand characterization of the Ti-Ni alloys,

Cu-basedalloys and ferrous alloys.

Th,cre are many good text books46~52) and re-views8,53~62) on SMEand PE, and they will be useful

for obtaining a deeper understanding of the funda-mentals of the SMAs.

C1989 ISI J 353

ISI J International, Vol. 29 (1989), No. 5

2. Fundamental Aspects of Shape MemoryA110ys

2. 1. Crystal Structure

Most of SMAShave superlattice structures, thesublattices of the parent phases being body-centeredcubic (BCC), as shown in Table 1. The parentphases of someother alloys (In-Tl, In-Cd, Mn-Cu)have a disordered lattice, i.e., face-centered cubic(FCC). If wetreat these alloys as exceptional cases,the parent phases of all alloys listed in the table havesuperlattices associated with the BCCstructure, andthese are classified as ,9-phase alloys, whenelectron

per atom ratio (e/a) is close to l.5.

Recently, it has been found that someferrous alloys

also show the SMEas shown in Table 2. If weexclude Fe3Pt, which has a superlattice associated

with the FCCstructure, all ferrous alloys have a dis-

ordered structure (FCC) in the parent phases.

Most of the P-phase alloys are divided into twotypes according to the superlattice or compositionratio. Onetype is denoted by fi2-phase, which has aCsCl-type B2 superlattice and about 50: 50 composi-tion ratio. The other type is denoted by P1-phase,

which has an Fe3Al-type D03superlattice and about75 : 25 cornposition ratio. Crystal structures for bothcases are shownin Figs. I and 2, where (a) illustrates

the three-dimensional structure, (b) the arrangementofthe atoms in the (1 lO) plane, and (c) the arrange-ment in the (110) plane above or below (b). Thecubic structure in (a) can be constructed by alternate

stacking ofthe planes in (b) and (c).46)

The martensitic phases of the fi-phase alloys also

have superlattices inherited from the crystal structure

A2

oo

- [110]

OCs (b) (110) plane (c) the (110) planc(a) unrt cell

OCl above and belowthe plane in (b)

A2' B2: (1 10) planes alternately stacked46'

Fig. l. Crystal structure of B2 type structure (~2 Parentphase).

A*(llO

,

L~l lr ~ f~~

l~I

(r- f7~l~

f~;

Fig

A*(llO)

i

A] Bl

- [110]

(a) unlt cell

e Fe OAl(b) (110) plane (c) the (110) plane

above and belowthe plane in (b)

A1' B1: Planes alternately stacked46)

2. Crystal structure of D03type structure (PI parentphase).

d

Table l. Non-ferrous alloys exhibiting perfect shape memoryeffect and pseudoelasticity.

Alloy Composition(ato/o)

Structure change Ternperaturehysteresis (K)

Ordering

Ag-CdAu-CdCu-ZnCu-Zn-X

(X=Si, Sn, Al, Ga)Cu-Al-Ni

Cu-SnCu-Au-Zu

Ni-AlTi-Ni

In-TlIn-CdMn-Cu

44-49Cd46

. 5- 50Cd38. 5-41

.5Zn

A few ato/o

28- 29A1

3- 4.5Ni

- 15Sn23- 28Au45- 47Zn36- 38Al49-5lNi

l8- 23Tl

4- 5Cd5- 35Cu

B2-2HB2-2HB2-9R, rhombohedral M9RB2 (D03)-9R, M9R(18R, M18R)

D03-2H

D03-2H, 18RHeusler-- 18R

B2-3RB2-monoclinicB2-rhombohedralFCC-FCTFCC-FCTFCC-FCT

-15-15

- lO

- lO

~35

~6

- 10

20- 100

1-2~4~3

OrderedOrderedOrderedOrdered

Ordered

OrderedOrdered

OrderedOrdered

DisorderedDisorderedDisordered

1

Table 2. Ferrous alloys exhibiting perfect or nearly perfect shape memoryeffect.

Alloy Composition Structure change Temeraturehystresis (K)

Ordering

Fe-PtFe-PdFe-Ni-Co-TiFe-Ni-CFe-Mn-SiFe-Cr-NiMn-Si-Co

-25 ato/ ptlo

-30 ato/o Pd33 o/o Ni, 10 o/o Co, 4o/o Ti (wto/o)

31 9/0 Ni, 0.4 o/o C(wto/o)

-30 O/o Mn, -5 o/o Si (wto/o)

-lO o/o Cr, o/o Ni, Mn,o/o Si, o/o Co(wto/o)

L12-0rdered BCTFCC-FCTFCC-BCTFCC-BCTFCC-HCPFCC-HCP

SmallSmallSmallLargeLargeLarge

OrderedDisorderedDisorderedDisorderedDisorderedDisordered

354

ISIT International, Vol. 29 (1989), No. 5

of the parent phase. Uponmartensitic transforma-

tion, each (110) plane of the parent phase deformsnearly to a hexagonal network and shifts in the [llO]

direction by a shear, resulting in the arrangement of

atoms in the (110) plane as shown in Fig. 3 or 4.

Accordingly, the resulting crystal structures of the

martensites can be constructed by stacking the 3types

atomic planes (A, B, C) in Fig. 3 or the 6 types (A,

B, C, A', B'. C') in Fig. 4, all types ofthc martensite

crystal structures being shownin Fig. 5. Martensitic

phases al' Pi, ~i' and ri produced from pl-phase haverespectively 6R, 18R(1), 18R(2) and 2H structures,

while martensitic phases a~, P~ and r~ produced fromfi2-phase have respectively 3R, 9Rand 2Hstructures.

Themartensitic phase of the Ti-Ni has a different

structure, which is three-dimensionally close packed

as shownin Fig. 6, although the parent phase has aCsCl-type B2 superlattice. The Ti-Ni alloy also

shows another phase transition prior to the marten-sitic transformation under a certain condition. This

phase transition consists of two processes; i.e., B2=~incommensuratephase-> commensuratephase (rhom-bohedral phase or R-phase) transition. The lattice

does not change upon the former transition, while it

changes to the rhombohedral lattice upon the latter

transition. The R-phase can be formed by elongat-

ing along any one of the l> directions of the B2structure as shownin Fig. 7.

A B

L,LiJc

' OCs OCI

Fig. 3, Three types of close packed stacking planes in

martensite with a periodic stacking structure pro-duced from B2 type p2 parent phase.46)

60 "

T"L

r' a--r

A B

2.2. Crystallography of Martensitic Transformation

Martensitic transformation occurs in such a waythat the interface between the martensite variant andparent phase becomesan undistorted and unrotatedplane (invariant plane or habit plane) in order to

minimize the strain energy. In order to form such amartensite variant (habit-plane variant), it is neces-

sary to introduce a lattice invariant shear such astwins, dislocations or stacking faults.64~66) The lat-

tice invariant shear is generally twinning, which is

reversible, in the shapememoryalloys.

Crystallographic characteristics of martensitic

transformations are now well understood by the

phenomenological crystallographic theory.64~66) Thistheory describes that the transformation consists of

the ibllowing three operational processes: (1 )a lattice

deformation B creating the martensite structure fromthe parent phase, (2) a lattice invariant shear P2

(twinning; slip, or faulting) and (3) a lattice rotation

R. Thus, the total strain (or the shape strain) asso-ciated with the transformation is written in the fol-

lowing matrix form:

A A

ZhdanovRamsdel

~

B. B'

AB. B'

A(11)

2H

AC'

BA

CB

A(1)3

3RA

cA

A

B

c

B

BA'

BA'

A'

Ac'

Ac'

c'

B'

c

c

B

B

BA'

A'

A'

Ac'

c*

c'

B'

B C C CA' B' B'B

C C C AB' B'B'B

A AA A(1)6 (113D3(21)6(21)3

6R 18R(2)18R(1)9RFig. 5. Periodic stacking structurcs with various stacking

sequences.3s'

Ti

NiIst layer

Al C'

O Fe O Al

Fig. 4. Six types of close packed stacking planes in mar-tensite with a periodic stacking structure produced

from a D03-type P1 parent phase.46]

[100]

t tiO Ni

Othla:er

lO~\[] []

97t 780

O-~ [OOl]

Fig' 6. The crystal structure of Ti-Ni martensite' viewed

from [OIO] direction'30)

355

ISIJ International, Vol. 29 (1989), No. 5

Fig. 7.

~~~

~~~a[100JB2

brolO]B2

(a) B2

at tllO]e2

ct:~~ll{~2[IOOlf~

~ c tOOllBa

a[lOO, Ie2,L*

O(

b' [71~]B2 b[OIO]B2[OIOIR

(b) R-phase

Unit cell of (a) the parent B2 phase and (b) the R-phase. The principal axes in the lattice deforma-tion associated with the R-phase transition is also

shownin (b).31)

P1= RP2B ............(1)

This theory requires that the shape strain producedby the martensitic transformation is described by aninvariant plane strain, i.e., a plane of no distortion

and no rotation, which is macroscopically homogene-ous and consists of a shear strain parallel to the habitplane and a volumechange (an expansion or contrac-tion normal to the habit plane). Thus, the shapestrain can be represented in the following way:

P1= I +mldlpl, .•••...........

......(2)

where) dl: a unit column vector in the direction ofthe shape strain

I: the (3 x 3) identity matrix

ml : the magnitude of the shape strain

pl : a unit row vector in the direction normalto the invariant plane.

Ifwe knowthe lattice parameters of the parent andmartensitic phases, a lattice correspondence and alattice invariant shear, the matrix P1 can be deter-minedby solving Eq. (1 )under invariant plane straincondition. Then, all crystallographic parameters suchas pl' ml' dl and orientation relationship are deter-mined.

Although the phenomenological theory has beenapplied to manyalloys, overall agreement betweentheory and experiment wasobtained only fbr Au-Cdand {3, 10, 15} transformation in Fe-Pt, Fe-Ni-Cand Fe~1-C alloys. However, recently excellent

agreement has been obtained in every respect for

Cu-Al-Ni and Ti-Ni alloys by introducing Type IItwins.36,67) The former alloy includes the MType 11 twinning as the lattice invariant shear,67,68)

while the latter one the 1>MType 11 twinning 32 36)

2.3. Origins of ShapeMemoryEffect

The driving force for the reverse-transformation is

the difference betweenthe chemical free energy of theparent and martensitic phases above A*, and thecomplete shape recovery lies in that the original ori-

entation ofthe parent phasecan be restored (crystallo-graphic reversibility). Thc crystal structures of non-ferrous shape memoryalloys exhibiting thermoelasticmartensitic transformations are ordered, except forthe alloys exhibiting FCC- FCTtransformations.

The origin of complete reversibility upon reversetransformation is explained for an ordered alloy inFig. 8.69) If the ordered arrangement of the atoms is

ignored, there are 3 equivalent lattice correspon-dences; these arc represented by the rectangles markedA to C in Fig. 8(a). If the re.verse transformationtakes the path A in the flgure, the original B2 super-lattice structure can be restored as shown in (b).

However if the reverse transf,

ormation takes the pathB or C, the product phase will be as shown in (c),

which is not the B2 structure anymore. In this case,the chemical free energy of the product phase is

higher than that of the B2 structure, and thus the

reverse transformation taking the path B or C be-

comesunfavorable. Therefore, the original orienta-tion of the parent phase can be restored.

The FCC->FCTtransformations in In-TI andother alloys also occur in a crystal]og_ raphically re-versible manner in spite of disordered structures.This is because the lattice deformation upon trans-formation is extremely small, and hence taking thepath to the original orientation of the parent phaseis easier than taking the other paths to the differentorientations.

Mo_st offerrous shapememoryalloys also have spe-cial situations for restoring the original orientation ofthe parent phase upon reverse transformation, al-

though they have disordered structures. In the Fe-Ni-Co-Ti alloy, fine coherent T2 Precipitates with L12order are formed in the austenite matrix by ausag-ing.43) Thus, the hardness of the austenite and thetetragonality of the martensite becomehigh, resultingin an elastic strain accumulation around the marten-sites and a high mobility of the interface between theaustenite and martensite phases.

It is nowknown that Fe-Mn-Si alloys show thecomplete SMEassociated with the stress-induced FCCto hexagonal close-packed (HCP)transformation.44,45)

The origin of the shape recovery in this case is ex-plained to be the preferential multiplication of a sin-

gle type of Shockley partial dislocations upon thetransformation, and they accumulate a stress fleld

which assist the partial dislocations movebackwardto

restore the original o_rientation oi' the parent phase

upon reverse transformation.

2.4. ShapeMemoryMechanism

Although the habit plane is an invariant plane tomakethe strain energy of the interface between the

parent and martensite phases nearly equal to zero,

(

~

~

356

ISIJ Internationall Vol. 29 (1989), No. 5

o• ee eOo(p\ oy A~\~\~;~

.

(10~1)B2

d,

'x:;r;b\~o_'ol

e e / e' e* ~O'_/) / O o ~!

e

~(\"-'~

o e c'+d e oe5 -

)(--\'/'i~'

[1 OO]B19

(a) (C )(b)

Fig. 8(a). Three possible lattice correspondences in re-

verse transformation ol' the B2 to B19 trans-

formation.(b). Parent phase (B2 structure) resulting from re-

verse transformation with lattice corrcspon-dence A.

(c). Parent phase (different from B2 structure) re-

sulting from reverse transformation with laLtice

correspondence B. 69)

the martensitic transformation itself is accompaniedby a large shear strain. Therefore, the self-accom-

modation of martensite variants becomesalso impor-tant in order to minimize the overall strain energy,resulting in nearly no macroscopic shape change in aspecimen.

By applying a force, the most favorable variant

outgrows in a self-accommodating morphology, andflnally the maximumrecoverable strain is attained;

the strain remains even after unloading. By heatingthe deformed specimen to a temperature above Af '

the original shape recovers completely by reversible

reverse transformation.In order to understand the shape memorymech-

anism, it is necessary to clarify the above processes;i.e., (a) self-accommodation, (b) variant coalescence

to produce the most favorable variant upon loading,

and (c) shape recovery upon heating. These proc-esses will be explaincd in the following. The de-scriptions are largely based upon those by Saburi andcoworkers8,9) for Cu-based alloys and those by Miya-zaki et al.21,22) for Ti-Ni alloys.

2.4. I .Self-accommodation

Three types of self-accommodating morphologieshavebeenobserved; i.e., a diamond-shapedmorpholo-

gy in !9-Phase alloys which produce long period stack-

ing order structure martensites, a triangular mor-phology associated with the martensitic phase in the

Ti-Ni alloy, and a cross-marked morphology asso-ciated with the R-phase in the Ti-Ni alloy. Recent-ly, all these self-accommodating rnorphologies havebeenwell undcrstood; each of themwill be explainedin order in the following.

There are, for example, 12 crystallographically

equivalent lattice correspondences in the D03-~ 18Rtransformation. Because there are 6 {OI l}Do* Planesand in each of them there are two shear directions( 1>and ) available for producing 18Rmar-tensite from the D03structure, the combination ofeach one of the planes with each one of the shear

directions results in 12 crystallographically equivalentlattice correspondences. There are two solutions inphenomenological theoretical calculations for eachlattice correspondence, resulting in 24 martensitevariants as shownin Fig. 9, where numberswith orwithout prime indicate the lattice correspondencesand label (+) or (-) distinguishes two types of solu-tions for each lattice correspondence.9)

Each diamond-shaped self-accommodating mor-phology consists of 4 martensite variants clustering

about each {lIOJ pole, and hence there are 6 self-

accommodatingmorphologies in all. Anexampleofthc diamondconsisting of 4variants is shownin Fig.lO. The net strain of the diamond is quite smalland the interface between cach two of the four mar-tensite variants is a twinning plane (T 2 8)18R or(1 O10)18R.

There are also 12 Iattice correspondences betweenthe parent B2 and martensite B19' in the Ti--Ni alloyalso. A martensite variant is formed by combiningtwo of the 12 Iattice correspondences, resulting in 24martensite variants as shownin the standard (OOl)B2stereographi_c projection in Fig, I l, where the closedcircles indicate the habit planes and the arrows theshear directions.22]

With reference to Fig. I l, various possible self-ac-

commodatinggroups are conceivable. They involvehabit planes symmetrically clustered about {OOl},{Oll} and {111} poles relative to the parent phase.However, the trace analysis revealed that the self-

accommodatingmo_rphology is triangle and consists of

3of 4 martensite variants clustered around one of 3{OOl}1~2 poles. An example of the triangular mor-phology is shownschematically in Fig. 12, where theinterface betweeneach two of 3variants correspondsto a twinning plane; i.e., M Type 11 twinningplane or (OOl)M compoundType I twinning plane.22)

The calculated shape strain matrix of the triangle is

quite small, although non-zero. There are 16 group-ings which fbrm such a triangular self-accommodatingmorphology about each of the {OO1}B2poles, result-

ing in 48 possible combinations in all to form suchtriangles.

TheR-phase transition in the Ti-Ni alloy is essen-tially a martensitic transformation, and hence the R-phase va,riants form a typical cross-marked self-ac-

commodatingmorphology consisting of all 4 types ofvariants.21) Three types of the possible minimumunit of self-accommodating morphology are showninFig. 13. There arc 4 variants A to D cxpressed bythe elongated axes I11>B2 mthe figure, where thetwinning planes {O11} and {OOl}, which combinetwo variants with a twinning relationship, are also in-

dicated. The shape strain matrix of each morpholo-

gy wasnearly the unit matrix.

2.4.2. Variant CoalescenceuponLoadingAs explained in the preceding section, each mar-

tensite forming a habit plane consists o_f two lattice-

correspondence martens-ites which have a twinningrelationship with each other in fi-phase and Ti-Nialloys, while a single lattice-correspondence R-phase

357

ISI J Internatlonal Vol. 29 (1989), No. 5

Too

oTo

TTO

.2'f+ )

I,4(-)

5'(-,.

EE

6(+) -

H5(t). Tol

TT1 2[f)

21(_]

ITo

,3'(-)

.4'( -]

OTI

.3(*,

6/(_l

I'f -)

OOI

4'(

+)OII

5(-)'

TII

n4(+)

6(-)

m

TIO

lll 2C-,101,

I (+)

I'Ct]

.

"IL'I 1J l+1

lll

,61( t l I

5(-~

3'(t'

IC-]

iOO

IIO

OIO

Fig. 9.

Stereographic projection showing calcu-

lated habit plane poles for the 24 mar-tensite (18R) variants of the Cu-Zn-Gaalloy. Four habit plane poles are clus-

tered about cach {110} pole. Each ofthe six {110} groups is designated bythe Romannumerals I -VI as shown.9)

I

(

1'

Fig. 10.

2

5

l~T5)003

(T o 1)Oo3, f~ 28)18R

e' (o To)003' (1 o

(~15)003

10)18R

Geometrical relations amongfour habitvariants in 18Rmartensite (group 11).9)

plane

forms a R-phasevariant in the Ti-Ni alloy. Besides,

each variant is connected with another variant by atwinning plane. These twinning planes moveeasily

at a low stress upon loading, resulting flnally in form-ing the most favorable martensite or R-phasevariant.If twinning planes remain in the martensite variant,

they also movefbr producing the most favorable lat-

tice-correspondence martensite. At this stage, the

maximumrecoverable strain is attained. Generally,the maximumrecoverable strain is larger than theshape strain calculated by the phenomenological the-

ory.Themaximumrecoverable strain can generally be

OI~O

~lOO

Tl~O

Jl \ TIO

\Ill Tol Tll

/ ~~\~~~

/ l~7 \OTl OOl Ol l

\ /;/ \~\ /~> cl

ITl IOllll

\\ /ll~O f llO

OIO

IOO

11. Standard (OO1)B2 stereographic projection show-ing calc.ulated habit planes (1) and correspondingshape-strain directions (-~) of the Ti-Ni alloy.22)

d

1

Fig.

calculated by the lattice distortion (B) which general-ly consists of distortions along the principal axes anda shear componentparallel to the basal plane of themartensite.8) Examplesof the calculated results for

the pl -> fii transformation in Cu-Zn-Ga, the B2-~

358

ISIT International, Vol. 29 (1989), No. 5

\11'~\\to/~!~~'

:~ei~/ Il~)O)

O'~o~~'

~o/,\:Q'.

//~"(e

l

l

C\,

CO

I~5

Z

!

l

:~

oo~I l:O.;A

~)I/2 O~~)

'Ol/ .V~~9:/

2 cPll

2l/ (O.72

, I, I)M '

2ll

/ 2Il

I

Fig. 12.

Schematic diagram of triangular B19-

martensite morphologydepicting crys-

tallographic relationships betweenvariants in the Ti-Ni alloy.22)

/

2/ 1 2/ l 21

(0.78 ,0.39 ,0.48)B2

Il

2

~:\~~

or~.

:~~~\r'~))~(~ro;1;r~ ~~

2/ l2/~rl ~)~\1(Q~~

/~~(O;>cP

I

'O)6!~)

'l)~~~)'

"/)4t~.

Type r

(OOl,

Type ll

(lOO,

Type ID:

(OOi)

Fig. 13. The minimumunit of self-accomTnodating R-phase

variants morphology in the Ti-Ni alloy. Threepossible types ofvanant combmationsare shown21)

B19' transformation in Ti-Ni, and the B2-> Rtransi-

tion in Ti-Ni are shownin Figs. 14 to 16, respective-

ly.

2.4.3. ShapeRecoveryShaperecovery occurs by the reverse transforma-

tion accompaniedby the movementof the interface

between the parent and martensite phases uponheat-

ing. The shape recovery starts at A, and completes

at Af' However, the shape recovery occurs in a dif-

ferent mannerupon the reverse R-phase transition as

shownin Fig. 17, which shows the shape change vs.

temperature diagram upon cooling, Ioading and heat-

ing.21) Although the reverse transition takes place at

TB, most of the strain recovers by the changeof unit

cell upon heating to TB. In fact, during heating, nomovementof interface between the R-phaseand par-ent phasewasobserved.

2.5. Successive Stress-induced Transformation

It has been found that stress affccts the martensitic

-8

-6 IIO

-4) lOl

111-2

I11 o4 2

6g8 OIO

OOl P Oll

lil 111

O1

llO

1oo

Fig. 14. Stereographic rcpresentation of the calculated

maximumrecoverable strain occurring in the Pi

to ~{ (18R) transformation in the Cu-Zn-Ga9)alloy.

transformation not only by changing the transforma-

tion temperatures but also by inducing newkinds of'

martensites. Fig. 18 shows stressstrain curves of' aCu-Al-Ni single crystal deformed at various tempera-tures.11) In each curve each stage corresponds to astress-induced martensitic transformation. By plot-

t sfor-ting the stress for each stage representing a ran

359

ISIJ International, Voi. 29 (1989), No. 5

8.0- -7. O- -

e(7, 5)

6.0- -

5.0 - -e4O- - (s.s)

3.0- ::- 4s) e

e(10.6 )

(1 1I] [cal ('1.)]

(10.5]_ _ _10.5

e( Io o)

- -10.0(s4,

e- -9,0

(OO1) (O 11)

Orientation dependenceof the maximumrecover-able strain associated with the martensite in theTi-Ni alloy. Numbersin the parentheses repre-sent the experimcntal results, and the contour lines

the calculated ones.34)

I,O

~s~oo

'~ 0.5~~

o

(a)

l

(3) (b)

Heati~g $

(2) Loading

(I)Cooling(c)

$

Fig. 15

/ ///tI ' /t l / /I l I /l llI ' llll (C) I Il

TTl l

/ ///// /'~ '~\/ / \

\\ 1 Il Il ILlL\ 11'

/ / \/ -~f f/ (A)+~ \_ \ 1lII +0.8 1 l

-0.4 t t\

x\\ IL\ \

\ \\\ \ ~\ \

l \\\ \ /

I n 'l +v.6 /l+0.4

\ \\ \\ \ 'L

\ \ \ / / / /\ / / /\ /

\ \ \ \\ ~~ ._- / /// /i0.2 \ \\~,\ \ \

\ \ \\ \\\ 0.0 \\ \ \\

+0.2 \ III \\ \

\\( D!) \\\

\:~

\ \ \\\ \ \0.4 \ \ \\\ \\ \ \\ \ \

\ \\ ~~

OI~+-_~l L_-+0.2\ \ --0.0'-'~\ \ ,,,,

\ \ \ - -0.2'-\

\ III_(B) ~ ~0.4-

\ \ \ \\ \ \\ \

\ \

Fig. 16. Standard (OOI)B2 Projection showing the calculatedorientation dependenceof the recoverable strainassociated with the R-phasein the Ti-Ni alloy.31)

mation, wecan construct a phase diagram represent-ing several martensite phases in stress-temperaturecoordinates as shownin Fig. 19, which is a schematicrepresentation.

There are three types of transformation sequencesaccording to the test temperature. Figs. 18(a) and18(b) reveal successive transformations upon loadingthrough the path I in Fig. 19, i.e., r{ -~ Pi' -~ al trans-formation sequence. In this case, the transformationsequence is not crYstallographically reversible; af -~pl -~ rl transformatlon sequenceappears upon unloading. Figs. 18(c) and 18(d) involve a fil -~ ri -~ Pi' -~ai transfbrmation sequence. Finally, Figs. 18(e) to18(h) involve a fil -~ Pi a{ crystallographically re-versible transfbrmation sequence.

Ti-Ni alloys also show successive stages of trans-formation in the stress-strain curve. The deforma-tion is associated with both the R-phaseand the mar-tensite in Ti-Ni alloys which include a high density

Fig. 17.

240 260 Z80 300 320Temperature(K) t

TR

(I) Formation of self-accommodatedmorphologycomposedof 4variants uponcooling

(2) Formation of the most favorable single vari-

ant with respect to thc applied stress bytwinning

(3) Recovery of strain due to the change ofrhombohedral angle uponheating

Shape change vs. temperature diagram showingthe shape memoryeffect associated with the R-phase transition in thc Ti-Ni alloy.21)

of dislocations and/or fine Ti3Ni4 precipitates.30)Therelbre, the deformation behavior is sensitive totest temperature; it is divided into 6 categories ac-cording to the relative relationship between test tem-perature and transformation temperatures as schemat-ically shownin Fig. 20.

In range I (T only one stage associated withthe rearrangement of martensite variants appears asshownin Fig. 20(a). In range 2 (Mf T boththe R-phaseand the martensite coexist, revealing twostages associated with the rearrangement of the R-phase and martensite variants as shownin Fig. 20(b).In the figure, the stress-strain curve associated withthe R-phase is drawn by a solid line, while that asso-ciated with the martensitc by a broken line; thedashed line shows the shape recovery associated withthe two reverse transformations upon heating. Inrange 3 (M. the specimen is in a fully R-phase state prior to loading, and hence deformationfirst proceeds by the rearrangement of the R-phasevariants to a favorable one as shown in Fig. 20(c).Uponfurther loading the martensite is stress-inducedin the second stage. In range 4 (Af thePEassociated with the martensitic transformation ap-pears, although a part of the deformation is still asso-ciated with the rearrangement oi'the R-phasevariants.In range 5 (Tl~ T T,,), the R-phase is also stress-induced, exhibiting two-stage PE (Fig. 20(e)). Thecritical stresses for inducing both the R-phaseand themartensite satisfy the Clausius-Clapeyron relation-ship. Sincetheslopeofthe stress--temperature rela.tionfor the R-phaseis steeper than that for the martensite,both lines cross each other at a temperature T..'rhus, the deformation associated with the R-phasedoes not appear in range 6 (T* T) as shownin Fig.

2O(f ).

,

d

360

ISIJ International, Vol. 29 (1989), No. 5

(a) 223K pH, ai (a' ) (b)273K(T pl~)aj

reorientat ion f l

f-I,f$C~

fA~ I p'c C'I

v'l

~,(-A

(~)

rl'~~ P1

I~rP\

(d)

7~*) pl

1,','-~

~-• al

pll'

~ ('1

(c) 303K(Mf T PfH)al

reorientation 'l' SIM '~_-_

,,f~i+,

yp,(ra;tl~r

,

rll~ pT

(d)317 5K(Ms A'-'ai

yl:.pi f

If~(- a;

/A-P1~71 1(r F1

l

a,

ce,

h

CV1'(6R)

(pl" (18R2))

~yl'(2H)

pl' (18Rl)

@C

ll(pl' )

pl'(D03)

Fig. 19.

(d')~H) arl

(e) 354K(Af Tc) ~' --) (lrl (f)400 7K ~~"j

Temperature

Schematic phase diagram of Cu-Al-Nialloy in temperature-stress coordinates.u)

1,') p;

A:r al pf '(-OE1 A-) P1

pl~ p; P1(~'iA'~ C,1

y,~ P; ~- pl

~500:E~; 400

~SOOul*200"c 100

~ O

(g)444 5K( Tc

pl'~ pl

pl(- p{(r OLl

pl- a ,(g')

pl~ pl

(CI) T Mf

pa(~ p;

(b)Mf (c) Msa;M '/'/'1~r ! 'l7~'// / / olM~r~~i ~r~~~7! i / o:T:.LLr:/:::,/

-~~~.J______,'l // f

~LJ-- -J

(h)477 5K ~L)a;

p,- p;

Ft~G'\

r rlpl

(rf)

cr)

(d)Af

~r, -~ ~~

!~~--J1

(e )TR T TX

o'iM

~fl~,t, l

o':

"'~L' Jt

tf)Tx T-r'~-1! lf,~.___. J

iL_.

o

Fig. 18.

5 10 15

E[ongat ion ('1.)

Stress-strain curves assuccessive stress-induced

(mass?•'~) alloy.11)

a function of temperature representing

transformations in Cu-14.0Al-4.2Ni

Fig. 20.

Strain

Classification of stress-strain curves ac-cording to test temperatures for Ti-Nialloys cxhibiting both the R-phase andmartensitic transfonnations.30)

3. Developmentof ShapeMemoryAlloys

3. 1. T, i-Ni Alloys

3. I . I .PhaseDiagram

Phasediagram is important to understand howto

heat-treat and develop alloys. H.owever, there havebeen a lot of difrilculties interferring with constructing

the phase diagranl of the Ti-Ni system due to the

following reasons :(a) Ti is so active that it easily combineswith O, C,

N, etc, at high temperatures.(b) There are several metastable phases at an inter-

mediate temperature region and one of themaffects strongly the transformation process andthe shapememorycharacteristics.

Therefore, there is no complete phase diagram of the

Ti-Ni system yet at present.However, according to a recent investigation,37) it

has beenfound that these metastable phasesare Ti3Ni4

and TiENi3 as shown in Fig. 21, the former beingthe important phase. Thesemetastable phasesfinally

transform to an equilibrium phase TiNi3 after a longperiod of' aging. The crystal structure of the Ti3Ni4precipitates has a unit cell with lattice parametersa=0.670 nm, a=113.8', and contains 14atoms.38~40)

If these precipitates are formed by a constrained

aging, a speciflc stress field causes a two way SME(the so-called all-round SME)to appear.70)

3. I .2. Trarisformation Process

Oneof the causes of the confusion in understandingthe martensitic transformation of the Ti-Ni alloy is

that so-callcd " premartensitic transition " appearsprior to the martensitic transfbrmation. This transi-

tion has' recently been investigated in the Ti-Ni47-Fe3(ato/o) alloy in detail by X-ray diffraction, electron

diffraction, electron microscopy and neutron diffrac-

tion.24,25) The Ti-Ni binary alloy also shows the

361

ISIJ International, Vol. 29 (1989), No. 5

o

"aE~be

b'

C

90

800

700

600

500

o o

o

o TiNiA TiNi fTillNil4 fTi2Ni3D TiNi .Ti 2Ni3 fTiNi9

o

OIOl0=0lO OI

o o o

' TINi cllll

ATiNi tTi 2Ni 3• TiNi •TtN i3

~1' ',

'

.'r-TI'-1-

II----I~-- IIA\ l\ lA \\P l lIJL\\ \\ l lJL- - -

e e

e e

l.

e

,

~e e

~,l ~o\:\

O.1 1 10 100Aging Time(hrs )

1000 10000Fig. 21.

TTTdiagram describing aging behaviorfor Ti-52Ni.37'

J

i

d

" premartensitic transition ", if it is subjected to aspecial thermomechanical treatment which introduces

a high density of dislocations and/or fine Ti3Ni4 pre-cipitates.30) Thetransformation proceeds in the abovecases as shownin Fig. 22. The electrical resistance

starts to increase on cooling at a critical temperatureTI' and on further cooling the resistance increasing

rate becomesrather gradual at a critical temperatureTB which is about 10 K below TI' The phase pro-duced flrst is incommensurate, since the superlatticereflections are deviated from the " one-third posi-tion " and the unit cell remains cubic. The incom-mensurate phase is locked into the commensurateoneat T~. The commensuratephase is called the R-phase, since the phase is rhombohedrally distort-ed.28,71) Thus the transformation in the Ti-Ni-Fealloy and Ti-Ni alloy containing dislocations and/orprecipitates proceeds as B2- incommensuratephase

- R-phase=~ martensite.However, solution-treated Ti-Ni binary alloys do

not reveal the R-phase transition. A solution-treatedequiatomic TiNi alloy showsno electrical resistanceincrement and exhibits the following transformation

sequence upon cooling, i.e., B2- martensite, while asolution-treated Ni-rich Ti-Ni alloy shows an elec-trical resistance increment at TI and reveals the fol-

lowing transformation sequence, i.e., B2- incommen-surate phase- martensite.30,72)

TheR-phasetransition usually appears prior to themartensitic transformation whenM* is more loweredby somemeansthan T1~' There are manyfactorseffective to depress M, as follows.

l) Increasing Ni-content73,74)

2) , Aging aftcr solution-treatmentl3,14,75)

3) Annealing at temperatures below the recrystalli-

zation temperature aftcr cold working 13,75)

4) Thermal cycling76~81)

5) Substitution of a third element23,75)

Amongthese factors, factors 2) to 5) are effective to

:h

aJ::

O::

7;

LLJ

-M~-M+R--~--R-~l~B2-BZ : Parent Phase

I :Incommensurate Phase

R :CommensuratePhase(Rhombohedral Phase)

M:Martensitic Phase

Mf

~

M

Ms~

TR+t ~ 'l

t Af ~As

M+RR~I ~-B2-

~

!

150 200 300 350250Temperature(K)

Fig. 22. Electrical resistanc.e vs. temperature curve for aTi-50.5ato/oNi alloy aged at 673 K fbr 3.6 ks aftcr

solution-treatment at 1273 K.21)

reveal the R-phase transition.

3. I .3. Deformation BehaviorThe martensitic transformation occurs by loading

as well as by cooling. Therefore, martensites arestress-induced even at temperatures above M*. Thecritical stress ((TM) for inducing martensites increaseswith increasing test temperature as shownin Fig. 23,

and it satisfies the Clausius-Clapeyron relationship,i.e., a linear relationship between all and tempera-ture.15)

However, if the critical stress ((Ts) for slip is low asindicated by asL in the flgure, there is no tempera-ture region in which the perfect PEappears. Thisis becauseslip deformation always occurs prior to thestress-induced transformation. If (rs is high as indi-

cated by aSHin the figure, the perfect PEappears.Fig. 24 showsthe schematic stress-strain curves for

(I

'~

362

ISIJ Interndtional, Vol. 29 (1989), No. 5

C..r.itical Stress for Sli asHc'~'-

tl,

u,

: S.ha.pe MemoryEffect ,q)

' : !~:/lJTransfcrmation 5~~''

' '

g;'!

' lpseudoelastla- \e~~

- ty c~~~:'

:~ '

~~~

~~:~oe'

'\c~

~~ \oc5\$:'b~,~-

•i'(Jc;~ Critical St

,(;~:~\~ ~~- -ress for S,ipasl

Mf MsAs AfTemperature --)

(Jc;~ Critical St(;~:~\~ ~~- -ress for S,ip

asL

aM: The critical stress for inducing martensites

asL: The low critical stress for slip

(rsH : Thehigh critical stress for slip

Fig. 23. Schematic diagram representing region of shape

memoryeffcct (SME) and pscudoelasticity (PE)in stress-temperature coordinates.

(/)

o

(a ) (b) (c) (d)

CiM-

ai,1-El

EI (11R'

,,,

u,u,u,

(;1R u)u,

~~Et

'l)L ~~

e~ &~ ~~

E2E2

l IE2I E2

Strain o Strain o Strain ~l~h~~atin9 Strain

Fig. 24. Schernatic stress-strain curves for

pseudoelasticity characteristics.

Ep

assessing the

assessing the PEcharacteristics (b) to (d), (a) beingfor the elasticity of a normal material.19) The areaE1 is the energy density which is dissipated during

one cycle, while E2 the energy density per unit volumewhich is stored and available uponunloading. Sincethe total work done on the specimen is (El+E2), theefhciency for energy storage r/ is defined as E2/(E1+E2)' This PE characteristics may be applied for

storing mechanical energy, becauseE2associated withPE in curve (c) is muchlarger than that with elas-

ticity in curve (a). For example, E2 for the Ti-Nialloy subjected to a special thermomechanical treat-

mentamountsto morethan 40 times larger than thatfor an ordinary steel spring.13) This meansthat only2.5 1of Ti-Ni alloy is enough to store the kinetic

energy of an automobile which weighs about I t andis running with a speed of 50 kmjh.

By comparing curves (b) and (c), it is easy tounderstand that both E2 and ~ increases with increas-

ing aM. However, in case 'T~l exceeds crs, PEbecomesincomplete, since the permanent residual strain (sp)

is introduced as shownin curve (d). Thus, in orderto obtain stable SMEand PE characteristics, it is

necessary to raise ('s'

For that purpose, the following three factors are

)important: (1) anncaling tempcrature, (2 aging tem-perature, and (3) Ni-concentration.13,14) Annealingat an intermediate temperature lower than the re-crystallization temperature thermally rearranges ahigh density of dislocations which were introduced bythe preceding cold work.13) Aging Ni-rich alloy pro-duces fine Ti3Ni4, causing the precipitation harden-ing. The density of such fine precipitates increaseswith increasing Ni-concentration, and hence agingeffect appears more strongly in Ni-rich specimens.13)

These internal structures prevent the movementofdislocations, resulting in raising (Ts'

An example of such effects is shownin Fig. 25,13)

where stress-strain curves obtained at various tem-peratures are shown for (A) a specimen solution-

treated and (B) a spccimenannealed at a temperaturebelow the recrystallization temperature after coldwork. It is clear that specimen (B) shows excellent

SMEand PE characteristics whencomparedto thesolution-treated specimen, which shows a large per-manentstrain.

3. I .4. Cycling Effect

Thegreat advantage of SMEand PEis that thesefunctions can be used manytimes. However, if thealloy is sub.jected to thermal cycling76~85) or stresscycling,16,19) the transformation temperatures and de-formation behavior will change.(1 ) Thermal Cycling

Fig. 26 shows a typical example of the effect ofthermal cycling on the transformation temperaturesin the solution-treated Ti-50.6atc~)Ni alloy.80) Theseelectrical resistance vs. temperature curves show that

M, decreases with increasing number oi' thermalcycling. Transmission electron microscopy observa-tion revealed that dislocations were introduced bythermal cycling and the density of the dislocations

increased with increasing numberof cycling, 78,80)

Such kind of' thermal cycling effects can be sup-pressed by introducing flne Ti3Ni4 precipitates and/or

a high density of dislocations. Fig. 27 shows theeffect of thermal cycling in a Ti-50.6ato/oNi alloy

which was solution-treated followed by aging at 673

K for 3.6ks to introduce fine precipitates.80) Nochange was observed in the transformation tempera-tures.

If a specimen is age-treated at a higher tempera-ture (773 K) to introduce rather larger precipitates,

thermal cycling causes M, to decrease and TR to in-

crease.81) It is suggested that these changes in trans-formation temperatures are caused by a specific stressfield around the Ti3Ni4 precipitates.(2) Stress Cycling

Fig. 28 shows the effect of stress cycling on thestress-strain curve of the Ti50.5ato/oNi alloy; the

stress is controlled by choosing a suitable test tem-perature. General features of the effect of cyclic de-formation are as follows:

(1) Residual strain increases as shownby the de-viation of the starting point of the stress--strain curvefrom the original point.

(2) Critical stress OMdecreases.

363

ISIJ Internationall Vol. 29 (1989), No. 5

(B)

eoo

o200

o200

(~

2 o

c~ 200

C:Q'

oh 400

200

o

to)210K

- 6

(A)

(b,228X (c)247K

- 6 - o

(d,269K (e,28eK

o

,f)295K (9,304X

o

,I )325t(,h ,3f 4K

'a'208K (b'20eK (c'219K 'd'230K (e'239K

o ____

200

o o~~~ o ~~ o ~400

(f,248F( (9'258K (h'268t( (i)2771( (]'289K

-

200 l__ ____ _o oo (n)328K'k '29eK (I)300X eFn)318K

- 400~:~

,1' 200

(1)

~2 O'~~O

,

o

Ti - 49Sof ,,Ni1273X10 - 673XIQMs' 294KAf •318K

Sfrdn (")

eoo

400

200

(o)337XO (p,348K ,q)356K o [r )365K

Ti -49.8at ,,J\li

673, IOM$•251KAf •313K

,:)

)c,

e)OC

G,

oc!:,

v,

u,Q)O:

::,

o

oQ)

uJ

Fig. 25.

(c)

O 2 4 6 o 2 4 6 o 2 4 60 2 4 6 eStrorl (%)

(A) Solution-treated at 1273 K (B) Annealed at 673 Kafter cold workl3)

Effect of annealing temperature on the stress-strain curves at various test temperatures in the Ti-49.8ato/oNi alloys.

~lsIl

f

lI

lr

T~

lifT

N= 100

N= 50

a,

c~:

o

o:

c,

a,

uJ

v

~

*

~f

*

Ti-50. eat "/*Ni

1273h , hr IQ

-• 673KIhr IQ

,Ifl

llIlIll

N= lO

Numberof

therrnal cycies(N)

N= 100

N=50

N=10

N•5

N= I

1

1

N=5

~

N=1Ti - 51

. 6at'/. Ni

1273K Ihr IQ

O IOO 15O 200 2: 50 300 350Temperature (K)

Fig. 26. Effect of thermal cycling on the electrical resist-

ance vs. temperature curve for the solution-treated

Ti-50.6at,~Ni alloy.80,

O 100

Fig. 27.

I50 20O 25O 300 35OTcmperature (K)

Effect of thermal cycling on the electrical resist-

ance vs. temperaturc curvc for the 'ri-50.6atoA]Ni

which was age-treated at 673 K aher solution-

treatment .8o)

(3) Strain or stress hysteresis becomessmall.

The cause for the residual strain is the occurrenceofslip deformation during the preceding deformation.Internal stress formed by these slip deformation mayassist the formation of the stress-induced martensites;thus (T~l decreases. As the stress-induced transforma-tion occurs like the Ltiders dcformation in stage 1,16)

ahl is kept constant for the first cycle. However, sincethe internal stress fleld generally has a gradient in its

364

ISIJ International, Vol. 29 (1989), No. 5

strength, (;M increases with increasing strain even in

stage I after cyclic deformation. As a consequence,the hysteresis becomessmaller by cyclic deformation.The dcgree of thc change in the residual strain, a_hl

and the strain hysteresis increase with increasing ap-plied stress. However, all these values becomein-

sensitive to cycling with increasing numberof cycles,i.e., a steady statc of the PEcharacteristics exists cor-responding to the cyclic stress level. This fact can beexplained by the work hardening due to the introduc-tion of dislocations during every cycling. Therefore,it is also important to raise the critical stress for slip

in order to stabilize the PE characteristics againstcyclic deformation.

g,

~:~:

u,u,Q,

V,

a,

u,

Fq,

(a)

O(b) 298.5 K

2 2

40

2ao

N=1 10 20 50 IOO

2 2 o 2 o 2(c) eo8.5K

2

400

2oo

N• 1 5 10 50 lOO

2 o 2 o 2Stra in ('/.)

2 2Stra in ('/.)

Fig. 28. Effect of cyclic deformation on stress-strain curvesat various temperatures in the age-treated Ti-50 5ato/ Ni 19)

' /o '

3. I .5. Fatigue Properties(1 ) Fatigue Life

Fatigue test modeis commonlya tension-compres-sion cycling for normal metals and alloys, becausethey are subjected to such a cycling modewhenusedas structural materials, Iarge or small size. The first

fatigue data of Ti-Ni alloys were obtained in such away.86) Although the deformation modeof the Ti-Ni alloy seemsto bc reversible macroscopically, theTi-Ni alloy also showedfatigue failure.86,87)

Since thc Ti-Ni alloy is a functional materialcharacterized by PEand SME,the fatigue modesub-jected to such a matcrial in use is commonlya load-

ing-unloading cycling without or with heating after

each unloading. Fig. 29 shows such data obtainedin a tension-unloading fatigue modein the Ti-50.8ato/~Ni alloy.88) The data are represented by twostraight lines. Deformation modefbr each region is

different to each other; the short life region is for

cyclic stress-induced transformation, while the longlife region is for cyclic elastic deformation. However,as an exception, the solution-treated specimen (1 273

K I hr IQ) does not reveal PE, because the stress-

induced transformation is accornpanied by a large

amount of slip deformation in this case. If' we areconcerned with PE, we concentrate our attention onthe short life region for the other two types specimens,

which are subjected to age-treatment after solution-

treatment (1 273 K I hr IQ-~673 K I hr IQ) andannealing at an intermediate temperature after cold

work (673 K I hr IQ), respectively. The fatigue life

for the latter is about ten times longer than that for

the ibrmer; the former incl.udes fine precipitates only

as the internal structure, while the latter includes notonly the precipitates but also a high density of dis-

locations which were introduced during cold working.Therefore, it is clear that the internal structure withboth precipitates and dislocati,ons is effective to im-

prove the fatigue life to a certain degree. However,it is necessary to investigate the effect of other factors

systematically in order to further improve the fatiguelife.

(2) Fatigue Crack Nucleation89)

The fatigue data mentioned above were obtained

Fig. 29.

Effect of heat-treatments on the fatiguelife of the 'I~i-50.8at~~Ni alloy which

was tested at 293 K.e8)

800

oO_~:

~600u)v)q)

cr)

400(1)

u)

c(1)

H200

OlOI

~~~-

D

Ti- 50.8 at'/. Ni

o 1273KlhrlQA 673K Ihr IQ

D 1273KlhrlQ-->673Klhr IQCycled at 293 K

A~\o

ACl

Cl

[]-~

o~

I02 I03Nf

ICf l05 06

365

ISIJ International, Vol. 29 (1989), No. 5r

Fig.

~=~~j~~;~~

~~. ~:

(b) : Thc magnification of (a)89)

30. Scanning electron micrographs of a fatigue fractured surface of

a Ti-50.8ato/oNi alloy.

,,

using specimens produced in a carbon crucible by ahigh frequency induction melting method. There-fore, these specimens contained TiC inclusions. Fa-tigue cracks were frequently observed to nucleatearound these inclusions. One of the examples is

shownin Fig, 30, in which (a) showsa TiC inclusionat the crack nucleation site, and (b) exhibit the mag-nified photograph of (a). Therefore, it is expectedthat purification of alloys lengthen the fatigue life,

However, specimens produced by an electron beammelting methodshowedalmost the samefatigue life

as that ofthe specimencontaining TiC inclusions.

Wenowask whypurification of the alloy hardlyaffects the fatigue life. To answer this question, weneed to clarify the fatigue crack nucleation site inspecimens produced by an electron beammeltingmethod. Optical micrographic observation revealedthat a crack was nucleated at a grain boundary asshownin Fig. 31. It is important to notice that thefatigue crack nucleates at about 10 o/o of the totallife, irrespective of whether the crack nucleation site

is at a TiC inclusion or along a grain boundary. This

meansthat purification of alloys hardly affects thetotal life.

Since the TiC inclusion does not transform, trans-formation of the surrounding matrix produces a largestrain incompatibility along the interface between theinclusion and the matrix, causing a large stress con-centration. In case of a grain boundary, transforma-tion in one grain causes the second variant to beformed in thc neighboring grain in order to makethe strain incompatibility along the grain boundarydecrease. If the contribution of the second variantis not large enough to reduce the strain incompati-bility, a large stress concentration will be created atthe grain boundary. Thus both the TiC inclusion

and such grain boundary maybecorne crack nuclea-tion sites. However, a hundred percent of the trans-formation strain contributes to the strain incompati-bility around the interface of the TiC inclusion,

while part of the transformation strain will be elimi-

Fig. 31. Fatigue crack nucleated along a grain

boundary. 89)

nated from contributing to the strain incompatibilityalong the grain boundary. This is the reason whyafatigue crack preferentially nucleates at the TiC in-

clusion rather than along the grain boundary if' speci-

menincludes a high density oflarge TiC inclusions.

Onthe basis of the above discussion, it is possible

to improve the fatigue life ofthe Ti--Ni alloy by intro

ducing a combination of purification of alloys anddevelopmentof a favorable texture.(3) Fatigue Crack Propagation Rate

As mentioned earlier, the martensitic transforma-tion is the main deformatio_ n modein SMAs,and thisis the essential feature which distinguishes SMASfromthe other normal metals and alloys. Thus, the mar-tensitic transformation is expected to serve to relaxthe stress concentration at crack tips in SMAs.86)Since olu: is a function of test temperature (T), it is

expected that the crack propagation rate varies withT.90)

Fig. 32 shows the relationship between the crackpropagation rate (Aa/AJV) and the stress intensityfactor range (AK) at various temperatures.90) Thesedata can be divided into three temperature regions

as follows. In range I (T the crack propaga-tion rate shows the minimumvalue. The marten-sitic phase will be introduced at a low stress level

during the first loading or it mayexist prior to load-ing. In both cases, the martensitic phase remains at

a crack tip. Deformation proceeds by the movementof twin boundaries or interfaces between variants in

the martensitic phase, and hence the critical stress for

the deformation is very low and almo_st constant ir-

respective of temperature in this range. Therefore,the stress concentration at the crack tip will be relaxed

most effectively in this situation. This is the reasonwhy the crack propagation rate becomesthe mini-

mum.

1

1

,

366

ISIJ International, Vol. 29 (1989), No. 5

10-s

E'-,.>

~1~

10-7

10~~ e/AA

e~~i'~("/ /

l!

/L!D~!D~

Dl

D

Df

Test temperature

OexAAQ

423K373K333K318K299K253K

Ti - 50.80t 'l.Ni

673KIhrlO

Ms=266KAs=279KAf =311K

101 10 100

AK(MNm~3!2)Fig. 32. Crack propagation rate as a functron of stress m

tensity factor range for various temperatures in the

Ti-50.Sat~i, Ni alloy.90)

In range 11 (M, Tp) the crack propagation

rate increases in proportion to test temperature. Themartensite will be induced at the propagating crack

tip in every cycle to relax the stress concentration at

the crack tip. The critical stress for inducing the

martensitic transformation increases with increasing

test temperature. This meansthat the stress relaxa-

tion at the crack tip becomesdiffrcult with increasing

test temperature. Thus, the crack propagation rate

increases with increasing test temperature up to T1"

above which the stress-induced martensitic transfor-

mation does not occur or (T~d: is too high to relax the

stress concentration at the crack tip effectively.

In range 111 (Tp therefore, the crack prop-agation rate becomesinsensitive to test temperature.It corresponds to the crack propagation rate in the

parent phase. Thus, the experimental data in this

range almost satisfies the empirical equation (AajAJV=

5.1 x 106 (AK!Y3.5) m!cycle)86} relating the crack prop-agation rate and the stress intensity factor range for

normal metals and alloys, the empirical equation

being shownby a dotted line in Fig. 32.

3.1.6. Effect ofthe T'hird Element on PE91)

As mentioned previously, the third element addi-

tion is effective to changethe transformation tempera-tures. However, it is also effective to change the PEcharacteristics. Amongmanyelements, Cuwasibund

to be the most effective element to improve the PEcharacteristics. Fig. 33 showsthe stress-strain curvesof the Ti-Ni40-Culo alloy which exhibits the mostsuperior PE characteristics amongTi-Ni-Cu alloys

investigated. Superior PE can be obtained eventhough the specimen is deformed at a stress as high

as 600 MPaas shownin the curve at 365 K. Table

3 shows the stress-hysteresis of the Ti-Ni-Cu alloys.

It is clear that stress-hysteresis strongly depends onthe alloy composition. Theminimumstress-hystere-

sis is about 100 MPa,which was attained in 3 typesof compositions, i.e., Ti-Ni40~Culo, Ti-Ni44.5-Curo andTi-Ni45.5-Curo. Therefore, these 3alloys can be used

Q*:~

(1)

o

333K

3421

50Ti-40Ni-lOCu(at '/.]

IQ

Fig. 33.

24 024 024Strain (%,

Stress-strain curves at various temperatures in the

Ti-Ni40~Culo alloy. 91 l

Table 3. Stress-hysteresis of Ti-Ni-Cu alloys.91)

Specimen (at0/~)Stress-hysteresis

(MPa)

50.OTi-40

.ONi- IO.

OCu45

.5Ti-44

.5Ni- IO.

OCu44

.5Ti-45

.5Ni- IO.

OCu41

.5Ti-48

.5Ni- 10

.OCu

45.5Ti-49

.5Ni-5

.OCu

44.5Ti-50

.5Ni-5

.OCu

lOO

1OO

1OO400300200

as an energy saving material with high efficiency of

energy storage.

3.2. Cu-basedAlloys

3.2, I .PhaseDiagrams

The range of compositions where Cu-Al~~:i shows

SMEis the region where the single pphase exists at

high temperatures, and is limited to around Cu-14Al-4Ni(wto/o)' The phase diagrams and martensitic

transformations of Cu~~1-Ni alloys are basically the

sameas in Cu-AI binary alloys. Fig. 34 shows the

phasediagram for binary Cu-Al. There is a pphaseregion with a BCCstructure around the composition

ns.92) Inof 12 wt?/o Al in the high temperature regio

an equilibrium state, the ~-phase decomposesinto ana-phase (FCC) and r2-phase (r brass type structure)

at 838 K (565'C) by the eutectoid decomposition.

However, if a specimen is rapidly quenchedfrom the

psingle phase region, the eutectic decomposition is

suppressed and the martensitic transformation occursat temperatures below the M, temperature. Themartensite phase formed in this process differs ac-

cording to the Al concentration : starting with the

low Al concentrations, P', Pf, pi+ri and r{ phases

appear, the subscript indicating a superlattice. TheM, temperatures at which these martensite phases areformed can be represented by a continuous single line

in the diagram. With A1 concentrations more than

l I wtcyo' the disordered plattice transforms into the

ordered fil structure (D03 or Fe3Al type structure) at

367

ISIJ International, Vol 29 (1989), No. 5

~(u~

~,

4Jc~'

l~(L'

~~(U

H

Al concentration (wt'/. )9 lO ll 12 13 14 15 16

lOOO

900 a+p p p+ y2

800

700 a t

Ms _-~~~~r~\ -- 1I \ p+pl /I \ ll \ ." p] \ ~

I \ a+T2I

y2

I bI

'L

600 l \J l

I\

I \, I ~

p'l p+pl :I tI I. ,

l~

500 I\lI

LII

)\

400 a+a2 a2 II J

,lp] 1 l 7] a2+ r2

IlI

II

,1Pi+ri L300 I l I h

18 20 22 24 2

Fig. 34.

26 28 30 32Al-concentration (at'/.)

l.4 1.5 1.6

ela

Phasediagraln for Cu-AI binary system.92}

1200

1000

U~~::

800

&~ha'

600

400

a+e

a

L+a

L :LIQUID

L L+p

a+p

L+a+p

*+NiAl a+P+NiAl

p

p+r2

a+a+NiAl

a+p+y2

a+NiAl+ r=

+r2

the order-disorder transition temperature T., whichis shownby a single line between the eutectic trans-formation temperature and M, in the phase diagram.The order-disorder transition can not be preventedeven by rapid quenching, resulting that the marten-site phase formed in this composition range takes onan ordered structure by inheriting the order structureof the parent phase.

In the Cu-AI binary system with high Al concen-trations, even rapid quenching can not suppress theprecipitation of the T2-Phase, resulting that thermo-elastic martensitictransformations do not occur. SinceNi addition is effective to suppress the diffusion of Cuand Al, the p-phase becomesstabler in the ternaryCu-Al-Ni system than in the binary Cu-AI system.Fig. 35 shows the cross section of the phase diagramfor Cu-A1-Ni with fixed 3wto/o Ni 93) By comparingFig. 34 with Fig. 35, it is found that the boundarybetweenthe p- and (fi+r2)-phases shifts toward highconcentrations of Al by adding Ni, indicating thestabilization of the fi-phase; hence thermoelastic mar-tensitic transformations occur in the Cu~L1Niter-

nary systern. Since M, is around room temperaturein the vicinity of 14 wto/o Al as shownin Fig. 34, thecomposition range is used for practical applications,the martensitic phase formed being the ri• In theAl-rich composition range, Ni addition is necessary to

suppress the precipitation of the r2 phase.In the Cu-Znbinary system as shownin Fig. 36,94)

M* is too low in the composition range where thermo-elastic transformations appear. Therefore, it is neces-sary to adjust the transformation temperature by add-ing a third element such as Al, Ge, Si, Sn, or Be.Fig. 37 shows the cross section of the phase diagrarnof Cu-Zn-AI ternary system with flxed 6wto/) Al.95)

By comparing Fig. 37 with Fig. 36, it is found thatthe P-phase region largely shifts toward the low Znconcentration region by adding Al, resulting in mod-erate transformation temperatures.

Fig. 35.

O 5 10 15

AI Concentration (wto/.)

Phase diagram for Cu~l-Ni ternary system.Vertical cross-section with fixed 3wt~/o Ni.93)

900

800

~~ 700

(L)

*~*q)l:~

E 600t,

H

500

400

Zn Concentration (wt"/•)

5030 40 60

31.9(32,5)

902'

a

+p

p

38.3 454'

(39,0 ) 44.8(45,5)

p+

468'

48.2(48.9)

834'

p

r

Fig.

30 40 50 60

Zn Concentration (wt"/o)

36. Phasediagram for Cu-Znbinary systcm.94)

In Cu-Zn-AI alloys, there is a P-phase with a dis-

ordered structure in a high temperaturc region as for

Cu-A1Nialloys. However, since there is an order-disorder transition at an intermediate temperature,the p2 ordered structure (B2 or CsCl type) is formedduring quenching; this is the parent phase ol' theCu-Zn-AI alloys. In the range of Al-rich composi-tions, though, a B2~~ D03ordering transition occursat a relatively high temperature region. Therefore,the D03 structure appears in Cu-Zn-AI with theAl-rich composition.

368

ISIJ International, Vol. 29 (1989). No. 5

~,

e)~::$

":h

q,l:~

~,L)

H

1000

900

800

7oo

600

500

400

L+a L L+pL+*+fi

a+p p

a+pa+r *+r

P+7

II

7

90

4

85

9

~1 350e,

a, 300CLEQ,

Hco 250

(1~

E~20Q

=,~

H

Fig.

Cu- XAI - 4.0Ni (wt.%)

AfAsM*Mf

80 75 70 65 60 55CuConcentration Cwt"/o)

14 19 24 29 34 39

Zn Concentration (wt"/*)

37. Phasediagram for CuZnAlternary system fixed

6wto//o Al.95)

Martensitic transformation start temperature [K('O]

90:~8(,~e~~ 7()

- 60/\;~••()

4r)

30

Zn1020

301:~/

40 rc500J

6070

ao10

8090

Cu '~(),~o 7vr;(':Tv fo';020 /r) Al

Cu(o/o)

/h,b:~

~CqF\~

20/~;~

=_~~~~~J~~~~t~~:~~()~~F~F~'o~~~~~*

305~f'~~ ~f~...~~.~~:;~'~~~~:,~~s~~;s)~.1~~-~~1

~~~~L

e 168(-105)

/187( 86)

2Jr 209( 64)

Fig.

14.213.6 13. 8 14.414.0

Al Concentration (x: wt.O/o )39. Relation between A1 concentration and transfor-

mation temperatures in Cu-Al-Ni.96)

Table 4. Relation betweenNi concentration and trans-

formation temperatures in Cu-Al-Ni alloys.

All specimens were water-quenched after

solid-solution treatment at I 193 K.99)

Cu(wto/o)

M* MfNiAl(wt"/~) (tw"/.) (K) (K)

A*

(K)Ar(K)

~\v_ro~o

:l

7271

70

6968

2828282828

O1234

285265270248

l80

~37(- 36)

240250269233172

i\267(- 6) \

1~3r- 90)

332, :i~6(~,9,r,3)l 31 7. 3f)•(44. 32)

70

230277283260200

320290290280230

106Jr

Table 5.

426rl53)443rl 70)339(166)

15346(73)

. 3.',7(~4)

/311(38)3~1(48)

2~~(1"') /"iL)1r48) 75'~40(67)

l:\~~:!9~

(L;~

Relation between cooling rate and Ms in

Cu Al-Ni alloy.

All spe{:imens were quenchedfrom 1273 in-

to 288 K or 373 K.The cooling rate was changed by varying

the tcmperature of the quenching media100)(Tgm)'

Composition (wt'/.) Tq" (K) M, (K) Tg~ (K) M, (K)

Cu-14.OAl-3

.9Ni

Cu-14. IAl-4

.ONi

Cu-14.2Al-4

.ONi

288282262228

373

363333

308

421(148)419rl46)

10 80572(299) 'll

Fig. 38. Relation betweencomposition and Msin Cu-Zn-A1 alloy.95]

3.2.2. Transformation TemperaturesThe transro_ rmation temperatures of Cu-based al-

loys are sensitive to their compositions as well as thequenching rate. Fig. 38 shows the relation betweenthe composition and transfbrmation temperature for

Cu-Zn-Al.95) The triangle on the left shows thecomposition range for the ternary system; the blackcolored area is magnified to produce the right flgure,

where the measuredM*temperature are shown.Fig. 39 shows the transformation temperatures vs.

Al concentration relationship in Cu-Al-Ni alloys with

a flxed Ni concentration of 4.0 ato/0'96) The straight

lines represent the averages for more than 50 single

crystals and polycrystals, the scatter being about 20to 30 K. The transformation temperature decreases

with increasing Al concentration. Even though the

single crystal and polycrystal were produced from the

sameingot, the M, of the single crystal was about20 Khigher than that of the polycrystal. The lower

M, of the polycrystal is due to the constraining forces

from the surrounding grains.97,q8)

TheNi concentration also changes the transforma-

tion temperatures as shownin Table 4, although the

effect is not so large as the Al concentration.99) Thetransformation temperatures decrease with increasing

Ni concentration, the Al concentration being fixed.

Finally, the cooling rate also affects the transforma-

tion temperatures, as shown in Table 5.100) Theeffect ofthe cooling rate maybe caused by variations

in the Al concentration in the matrix due to the

precipitation of the r2'phase, exc~ss vacancies frozen

in by quenching, variations in the degree of order,

369

ISII International Vol. 29 (1989), No. 5

or thermal stresses induced during quenching. How-ever, the mechanismfor the shift of the transforma-tion temperature is not clearly understood at present.

Thermal cycling also changes the transformationtemperatures due to introducing dislocations or chang-ing the degree of order.101-ro4)

3.2.3. Aging Effects

There are several causes conceivable for the effectsof aging. The aging effects appear differently ac-cording_ to thc aging temperature and dependinguponwhether the aging occurs in the parent phase or inthe martensitic phase. In any case, the ag_ing effects

are associated with atomic diffusion. Thusthe effectsof the addition of elements which changethe diffusioncoefncients are also important. These problems willbe described in the following.(1 ) Aging in the Parent Phase

Generally, several causes are conceivable f 'o_r agmg

effects in Cu-based alloys : (1) change of the degreeo_f order after quenching and (2) the formation ofprecipitat,es.

In A3B type alloys, the ordering reaction duringquenching occurs in two stages; i.e., a disorderedlattice changes into a B2 at higher tem-peratures, and then the latter into a DO3superlat.tice.

Since the fi-B2 ordering transition occurs at a tem-perature around 773 K (500'C) in the Cu-Zn-Alsystem, this process will be completed after quench-ing. The B2-D03 transition shows a strong com-position dependenceand can be lowered even to roomtemperature. In such a case, an incompletely or-dered structure maybe obtained by quenching.

If specimens thus obtained are aged at 312 K, thetemperatures M* and A* which are the tempera-tures measuredby a differential scanning calorimeterand associated with M, and A,, respectively, increasewith increasing aging time as shown in Fig. 40.]05)

After 20-30 min, both M* and A* increase about20 Kand then becomeconstant. If these specimenswere subjected to flash annealing at temperaturesabove the B2-D03transition temperature fbllowed byquenching, M*and A* reverted to the values beforeaging.

Since the B2-D03ordering transition temperaturedependsstrongly uponalloy composition,ro6) the com-position affects the structure obtained by quenching.Fig. 41 shows the calculated results of the B2-D03ordering transition temperature in the Cu-Zn-AI ter-

nary system.105) Theshadedarea in the figure showsthe range of composition for which M, is about 323 K(50'C), indicating that the transition temperature fororde~.ring increases with increasing Al concentrationin the shaded area. Whenthe ordering transition

occurs at higher temperatures, the B2-~ D03atomicrearrangement occurs easily to form a complete D03structure. Therefore, Cu-Zn-AI alloys with highconcentrations ofAl will be ofthe D03structure, andthose with low Al concentrations will have the B2structure.

Another cause of the aging effects which appear inthe parent phase is the precipitation oi' r2.107) Fig.42 showshowthe Vickers hardness and the percentageshape recovery vary with the aging time for Cu-15Al-3Ni(wto/o) aged at various temperatures from473 to 673 K. Since the increase in hardness anddecrease in shape recovery begin after the samcagingtime, the two changes must be based on the samecause, i.e., the precipitation of the r2 phase. Theact.ivation energies for both the changes of the hard-

ness and shape recovery are about 80 and 60 k.1lmol(about 0.83 and 0.62 eV/at), which are roughly equalto the values for the diffusion associated with the

excess vacancies.(2) Aging in the Martensitic Phase

Generally, the martensitic transformations occur bythe cooperative movementof atoms, and thus themartensite inherits the ordered atomic arrangemcntof the parent phase. Accordingly, a possibility re-mains ol' further reducing the free energy by therearrangement of the atoms in the martensiticphase.93,105,108) This phenomenonbrings about theincrease of the reverse-transfbrmation temperature,because the martensitic phase becomesmore stable.

Cu- 50Al

270

~(

S~260~E~!

~=~250

V M*e A*

ee

i~ 40o\~,

•s\:'~

,,~4(fZ,"~^~~30

~O ot

O,O,~O(;e'~

~,r///;t7~ Or~;

~ ~~i~~,~

,!;O 20 r,t\~r;i~~,~

z//zZ///z

10 e:

~l~:oo O!r:~~O0'O!r:O~f !r

lIIINlutlll MS273t50KB2-D03contours

~0V)o0•~fooof~

~,,,ZZ

~:1,2

A7//////z

BzO 1 10 100

Aging time (min)

Fig. 40. Relation betweentransformation telnperatures andaging time. A•fter solution-treatment followed byquenching, the specimenwasaged at 312 K.105)

Cu 20 Cu-50Zn3010 40Zn concenttation (at.o/o)

Fig. 41.

Calculated results of the relation between the com-position and the B2~' D03order-disorder transi-

tion temperature in Cu-ZnlAl.105]

370

ISI J International. Vol. 29 (1989), No. 5

) 300IQ,C:

11,V

=,L,

~e.U~ 200

(a )

e_'L~~l

673K 523K, 623K 573KJlr drh,'l ' e~lP '~i7!ee

e -/:473K'~

e e'~~llee _L:c~:~:'~~~r

Aging time (s in log)5 6

600:1:

1)

~* 400~>o

~

)~.'

>q,>O 100UQ,

q,

~ O(D~:V,

(b)

Cu- 28Al - xNi (wt~f~,)

x=10 2,0

e

3.0

e-e• •?~44,l

~ ~- ~h673K 573K 473K

1 2 53 4Aging time (S in log)

Fig. 42. (a) Changein thc Vickers hardness and (b) the

shape recovery ratio as a function of aging in Cu=Al-Ni

.Io7)

If SMAis used as a temperature detector and anactuator utili7.ing the force produced by the reverse-transibrmation, the aging in the martensitic phase will

change the operation temperature after the SMAis

installed. Since the rearrangement of atoms due tothc aging in the martensitic phase is also associated

with diffusion, it is important to suppress diffusion

in order to stabilize SMEcharacteristics in practical

applications.(3) Effect of the Third Element on Aging

Thehardness of Cu-Al-Ni alloys with different Niconcentrations was measuredas a function of aging

temperature as shownin Fig. 4399) for a fixed agingtime of 10 min. The point at which the hardnessbegins to rise corresponds to the temperature whereT2 Precipitation begins. The point shifts to higher

temperatures when the Ni concentration was in-

creased. This inclicates that the Ni operates to sup-press the diffusion of the Cu and Al. In fact, in

binary Cu-AI alloys, diffusion is rapid and t.he pre-cipitation oi' the T2 can not be suppressed even byquenching.

3.2.4. Fracture and Fatigue

Oneof the most serious problems in applying Cu-based alloys are the brittleness due to intergranularfractures.100,ro9,1ro) The following are considered pos-sible causes of the intergranular fractures :l) the large elastic anisotropy,2) Iarge grain size,

3) the large orientation dependenceof the trans-formation strain, and

4) grain boundary segregation.

Cases I)-3) may produce intergranular fracture

through large stress concentration along grain bound-aries. Cause4) Iowers the strength of the boundariesthemselves and results in intergranular fractures.

In order to clarify the cause for the intergranular

200 300 500 6004OOAging temperature (K)

Theaging time : 10 min

X : TheNi concentration99]

Fig. 43. Relation between aging temperature and micro-Vickers hardness in Cu-Al-Ni alloys.

fracture, well designed bicrystals are ideal speci-

mens.111,n2) Six bicrystals were used for the purpose.They can be classified into three types according tothe relative orientations of the two componentcrystals

as ibllows :l) Asymmetric bicrystals l, 2and 3with stress con-

centrations at the grain boundaries due to an-isotropic elasticity,

2) Asymmetric bicrystal 4 which have a special

orientation relationship which does not producestress concentrations upon elastic deformation,but produce upon transformation, and

3) A symmetric bicrystal which does not produceany stress concentration in the boundary eiiherelastically or after the transformation.

Asymmetric bicrystal I wasa combination of com-ponent crystals in which the difference betweenelastic

strains at the grain boundarybecameextremely large.

The intergranular fracture occurred in this specimendue to the_rmal stresses whenit was quenched. Thedifference in the elastic strains at the grain boundaryin asymmctric bicrystals 2and 3wassmaller than in

asymmetric bicrystal l, and the specimens did notcrack during quench.ing. However, bicrystal 2frac-

tured in thc elastic deformation stage as '-hown in Fig.

44. If we assumethe cause of the cracking in this

case was stres*s concentration due to the difference in

transformation strain in the grain boundary, the frac-

ture stress should showa temperature dependenceasthe stress for inducing martensite does. However,grain boundary cracking occurred almost at the samestress irrespect.ive ofdeformation temperature. There-fore, it can be concluded that the cause of the inter-

granular fracture in asymmetric bicrystal 2was thestress concentrations due to the large elastic anisotropyofthis alloy.

In the special asymmetric bicrystal 4, stress con-centrations do not arise elastically but are insteadgenerated on the grain boundary by the difference in

transformation strains in component crystals. Thisspecimen did not fracture elastically even though thedefbrmation temperature or the composition werealtered, however, intergranular fractures always oc-

371

ISIJ International, Vol. 29 (1989), No. 5

curred after the stress induced martensitic transfor-

mation as shownin Fig. 45. Therefore, the cause forthe intergranular fracture lies in the stress concen-trations due to the large difference in the transforma-tion strains in componentcrystals and hence the frac-

ture stress dependson the deformation temperature orthe composition in this case.

In the symmetric bicrystal, both componentcrys-tals rotated around [100] and the tensile axis inclinedmutually ancl symmetrically lO' awayfrom the [OOl]direction in each crystal. Fracture occurred trans-granularly after deformation more than 20 o/o asshown in Fig. 46, where the stress-strain curves forsingle crystal, the symmetric bicrystal and asyn]metricbicrystal 2are compared.

Recently, an Auger electron analysis revealed that

oO*21"1"

2~5

~2'~ ic:oF

lOO

50

(a)

Deformedat 286K

lOO

50

(b)

Ms=203KAf *228K

Defermed at 213K

O (1.0)(0.5) OStrain ('/.)

X : Apoint of final fractureill)

Fig. 44. Stress-strain curves ibr asymmetric bicrystal~, the

deformation tenrperature being (a) 286Kand (b)

213 K.

oxygen atoms were at grain boundaries in a Cu-Al-Ni alloy, while they are not in a Ti-dope.d Cu-Al-Nialloy; the Cu-Al-Ni alloy showedintergranular frac-

ture, while the Cu-Al-Ni-'1'i alloy transgranular frac-ture.113,114) This fact suggests that the segregation of

oxygen at grain boundaries is also one of the causesof intergranular fracture.

Table 6 shows the observed fracture pattern ofCu-Al-Ni, CuZn-AI and Ti-Ni alloys along withthe elastic anisotropy factor, orientation dependenceof translbrmation strain, the crystal structure of theparent phase, stress for slip. Although large stressconcentrations due to the strong orientation depen-dence. of transformation strain can be expected in everyalloy listecl in the tab]e, Cu-Zn-AI with B2 structureand Ti-Ni shows transgranular fracture. Since thecritical stress for slip is low (100-200 MPa)fbr both

fO

~:~

e,

C~

800

6oo

400

200

Symmetricbicrystal (2) X

/// ~ ~( X

Si,tgle crystal li

/\ symmetric bicrystal (1)///

XAsy'nmetnc blcrystal (2) Cu- 14.6Al - 4.0Ni (wt.%)

O 10 3020Strain (o/o)

X : Thepoint offracturell2]

Fig. 46. Stress-strain curves for a single crystal and sym-metric bicrystal of Cu-Al-Ni

.

Fig. 45.

Effect of conrposition and test temperature

on the stressstrain curves in asyrnmetricbicrystal 4.112)

~~

~

U'

Fo

600

400

20o

(ej

~1X

(a Jx

Cb]

x (a)(b)(c(d]

Chemical content(wt%]Cu-14. 2Al-4 . ONiCu-14. 6Al-4 . ONiCu-14 . 6Al-4 . ONiCu 14 2Al-4.0Ni

Ms(K)

258198198258

T (K)

R.T.R.T.573223

(d,

X3,4

o o o oStrain e~J

Table 6. Relation between fracture patterns and various factors.12)

AlloyAnisotropic Dependenceof

factor transformationstrain on orientation

Parent phaseSlip stress

(MPa) Fracture pattern

CuAlNi

CuZ*Al

TiNi

13

15

2

I.arge

LargeLargeLarge

D03B2

D03B2

-600-200High

- 100

Grain boundary fracture

Transgranular fracture

Grain boundary fracture

Transgranular fracture

372

ISIJ International. Vol. 29 (1989), No. 5

cases, it seems that the stress concentration can beeffectively relaxed by slip deformation. Onthe other

hand, since Cu-Al-Ni and Cu-Zn-AI with the D03structure have a high stress for slip, it is not effective

for these alloys to relax the large stress concentration;

and hence grain boundary fracture occurs. There-fore, it is concluded that the critical stress fbr slip is

also one of the i_mportant factors controlling the frac-

ture pattern.The fatigue life depends strongly upon the cyclic

derormation mode.109,110) Fig. 47 shows the fatigue

life of polycrystalline Cu-Al-Ni specimens (A to C)with different concentrations. The M, temperaturefor specimenA is 180 K, and specimenBwas 273 K,and specimen Cwas 420 K. Consequently, the de-

formation modeat room tempcrature is elastic defor-

mation in the parent phase ibr specimen A, thestrcss-induced transformation for specimenB, and the

rearrangement of therrnally formed Ti martensite for

specimen C. The fatigue strength was highest in

specimen A, and lowest in specimen B. This indi-

cates the deformation modeinvolving the movementof parent/martensite interface significantly deteriorate

the fatigue properties.

3.2.5. Grain-sizc RefinementThere are several techniques for grain-size refine-

ment, although the basic principlc is the same, i.e.,

to increase thc rate of' nucleation and to suppress thegrain growth. The fo]_10wing techniques arc effective

for such purpose :(1 ) Adding other elementsll3-120)

(2) Splat quen,chingl21)(3) Powdersintering.122)

In technique (1), added elements should be se-lected to meet either of' the following criteria: theyshould have low solubility (e.g., B, Cr, Se, Pb, V,and Ni),113,115-n7) or they should combine with the

other elements in the alloy to form fine compounds(e.g., Ti).li8-120) Techniclue (2) also restrain graingrowth: by splat quenching molten alloys are instan-

taneously solidifled. This prohibits the crystal nuclei

to grow and thus keeps the grains extremely small.121)

In technique (3), flne powders are sintered at high

temperatures and pressures.122) Since thc grains will

retain their original size, this technique is equivalent

to suppressing the grain growth.Grain reflnement by adding newelements or pow-

der sintering has improved the fracture strength andfatigue life.113,122) Oneexample of the improvementoffatigue life by grain refinement is shownin Fig. 48,

where the fatigue lives of fine and course grainedCu-Zn-AI alloys are compared.

Recently, the workability of the Cu-Al-Ni alloy

has been improved by adding Ti and Mn119,123) Theconcept of the alloy designing is explained in the fol-

lowing. Grain refinement was achieved by adding_

Ti, resulting in a good hot workability. In order to

improve the cold-workability, Al concentration wasreduced from 14 at(yo (hypereutectoid composition) to12 ato/o (hypoeutectoid composition). In a hyper-eutecto_ id composition region, the r2 precipitates areproduced, deteriorating the cold-workability. How-ever, in a hypoeutectoicl compositions region, there

are two phases, i.e., a' and p, above the eutectoidtransformation tempcrature. Since the (~ phase is

FCCstructure, it is easily deformed, resulting in agood cold-workabili_ty (- 10 o/o at roomtemperature).However, it is further neccssary to decrease the M.temperature which has increased by reducing the Alconcentration (e.g., I ato/o incrcment of A1 causes the

M, to increase about 120 K). For that purpose, theaddition of Mnis effective, the final product beingCu-12o/o Al-5 o/o Ni-2 (/o Mn-Io/o Ti (at ).

3. 3. Fe Alloys

Amongferrous SMASin Table 2, Fe-Ptl24-127) andFe-Pdl28) have been investigated only for the aca-demic interest, because Pt and Pd are very expensiveelements. The other ferrous alloys are now beingdeveloped for practical applications. One of the

700

600

5001~CL~~f 400~aE 300~""~~ 200

iE 100

o

Al[oy A A

HR.

A

v

A.CA

v

AI toy

Alloy

Alloy

~Jl'\v'\A~;,~~\1r

\A[[oy c

e""'

VA~•\\~~1~;\~\\v\\\

.

A\Z;\-lL

~~\\o e~\e e

o ~~~~oI~;

\~~Jl:'Fb'

r"

Atloy B

ABC

105 1061O1 103102 Iof

Numberof cycles to failure (Nf)

Fig. 47. Fatigue life ofsingle crystals of Cu-Al-Ni.110)

107

600

_ 500:~l

A*~:

~ 400~,"O~~V) 300

200

100

c

c fine-grained

, , .f speclmens(d =30pm)

t~~~tc'

coarse-grained ~\cspecimens \\~t, ~~~~

~(d=200- 2000pmlsA\

. 2~AclJ~!1~ ~A\~'~,~'. \~ ~~~l~'~\ \1'~~;*"\ \- i~);~~,.~L ~'~

~~ -'~;:~~'7~~.

I *1_\ ~\, '- T~\ \

" '\AA I o-

103 10' 105 106

Numberof Cycles before Fracture

Fig. 48. Fatigue lifc in fine- and coarse-grained

specimens.122)

373

ISI J International, Vol. 29 (1989), No. 5

commonfeatures of the ferrous SMASis that the mar-tensite morphology is thin plate type.43)

The thin plate martensite is chtlracterized by thefollowing features; i,e., the lattice invariant shear is

perfectly twinning and the stress due to the trans-formation strain is accommodatedby the elastic de-formation instead of the plastic deformation. Hence,dislocations are rare in the austenite in the vicituty ofthe thin plate martensite, resulting in the formationof the mobile parent/martensite interface. In orderto achieve tlle above situation, the ibllowing 4 factors

are considered to be effective.129)

(1) The yield stress oi' the parent phase is highand/or the elastic modulusis lo_ w.

(2) Thevolumechangeand transformation shearstrain upon transformation are small.

(3) The tetragonality (c/a) of the martensite is

large. The twinning shear and thc twinning bound-ary energy decrease with increasing c/a; hence, thetwinning becomeseasy in the martensite.

(4) Msis low. The twinning becomeseasier to

occur comparedto the slip deformation with decreas-ing Ms; The yield stress of the parent phase also in-

creases.All the above 4 factors makeslip deformation dif-

ficult to occur. In cases of Fe-Ni-C130) and Fe-A1-C13i) alloys, the transition temperature for martensitemorphology from lenticular to thin plate increaseswith increasing C-content, indicating that the thinplate martensite becomeseasier to appear with in-

crea-sing (~-content. Besides, the tetragonality of themartensite and the yield stress of the parent phaseincrease with increasing C-content. High C-contentalso raises the Curie temperature (Tr.'), resulting in alarge difference between Tc and M, and hence asmall volume changeupon transformation due to theinvar effect. The above characteristics of Fe-Ni-Cand Fe-Al--C alloys are effective to form the thinplate martensite and henceSME.

However, these alloys include carbon atomswhichcombinewith Ni or Al to form carbides in the mar-tensite upon reverse-transformation, causing the SMEto occur inc.ompletely by suppressing the mobility ofthe parent/martensite interface. Several newferrousalloys have been developed along this line, i.e., Fe-Ni-Ti-Co, Fe=Mn-Si and Fe-Cr-Ni-Mn-Si-Co al-

10ys.

The characteristics of each of these alloys are ex-plained briefly in the following.

(I) Fe-Ni-CoTi43) :The martensitic translbrmation of the Fe-33(yoNi-

40/0Ti,-lO(/oCo(wto/o) alloy becomes thermoelasticwith a small thermal hysteresis if it is ausaged at793 K ibr 1.8 ks. Themartensitic transformation oc-curs from a FCCaustenite (r) to BCT(a'). Theeffect of each element is as follows : High Ni-contentmakesM* decrease. Addition of Ti causes orderedfine r' (Ni3Ti) precipitates to be formed, resulting in

a high yield stress and a large back stress for assistingthe reverse-transformation. Addition of ComakestheCurie temperature increase, resulting in a small vol-

umechangeupon transformation and a low modulus.

Onedifficulty in practical application of this alloy is

at present the low value ofM, (below 200 K).(2 ) Fe-Mn-Si44, 45) :

The SMEis associated with the stress-inducedtransformation frorn the FCCaustenite to e(HCP)martensite. If the specimen is cooled below M.,thermally induced martensites are formed, suppress-ing part of the shape recovery. Ivin is effective toreduce t-he stacking fault energy of the austenite andhence the e phase is easily stress-induced. Since Mnreduces M., the yield stress of thc austenite increases.However, Mnraises the N~el temperature (TN). IfTNbecomeshigher than Ms' the stress-induced mar-tensitic transf ormation is suppressed. However, Si is

effective to decrease T,v' Si is also effective to reducethe stacking fault energy and raise the yield stress ofthe austenite; assisting a good SME. Therefore, thereis an optimumrange for Si- and Mn-contents in orderto achieve good SME. Fig. 49 shows the effect ofMn- and Si-contents on the magnitude of shape re-covery. The composition range of Fe-28-330/0Mn-5_60/0Si shows nearly perfect shape recovery. Oneof' the advantages of this alloy is that M, is aroundroomtemperature.(3 ) Fe-Cr-Ni-Mn-Si-C0132) :

TheSMEis associated with the r -~ e' stress-inducedtransformation, similar to the Fe-Mn-Si alloy. It

was recently reported that the composition range of7-150/0Cr, and O-150/0Co shows complete shape recovery if the defor-mation does not exceed 4 ()o Strain. Another advan-tage oi' this alloy is a high corrosion-resistance as astainless steel. This alloy also has an advantage inM.*, because the M* is between 173 and 323 K.

4. Prospects in Applications of ShapeMemoryAlloys

As mentioned in Introduction, the total numberofpatents for applications of SMASamounts to morethan 4OO0.41) They include almost all kinds of in-dustrial fields, i,e., electrical engineering, machinery,transportation, chemical enginecring, energy, medi-

~~~e

~

~cOo

CD

6

4

2

o

, •

o o70,,

50,,

ro,5

o o Jt e

A , ,

24 26 3428 30 32 36 38

Mncontent (wt%)

Fig. 49. Effect of Mnand Si contents on the magnitude ofshapememoryfor Fe-Mn-Si alloys.44)

J

374

ISIT International, Vol. 29 (1989), No. 5

cine, and so on. This fact indicates a great demandfor applications of SMAs. However, the types of ap-plication products in the market nowamountto onlyabout 40,133) which is only less than I o/o of the total

numberof the applied patents. Oneof possible rea-

sons for this is probably the current pricc of Ti-Nialloys.

However, most of the current price of Ti-Ni alloysis not the cost of raw materials but the manufacturingcost. Thus, the price of Ti-Ni alloys is expected tofurther decrease in the near future if the amountofalloy production increases. Besides, if the charac-teristics of Cu-based and ferrous alloys are furtherimproved to the level suitable for applications, therewill be no problem of cost.

SMAand PEcharacteristics of Ti-Ni alloys havebeenwell developed by using special thermo-mechani-cal treatments, and they are nowin usc in actual ap-plications. Examples of' such applications includepipe joints, rimless glasses, brassieres, PEarch wiresfor orthodontics, etc. which require simple functions.

However, there are manyother potential applicationswhich demandmore precise and/or repeating func-tions. In order to apply Ti-Ni alloys for such pur-poses, further improvement is still necessary. This is

another reason for the small numberof products inthe market at prescnt. Somcof the key directions forthe further improvement will be the addition of thethird elements, alloy purification, developing a favor-able texture, and so on.

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