precipitation reactions in al–4.0cu–0.3mg (wt.%)...

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Acta Materialia 56 (2008) 2147–2160

Precipitation reactions in Al–4.0Cu–0.3Mg (wt.%) alloy

S.P. Ringer a,b,*, B.T. Sofyan c, K.S. Prasad a,d, G.C. Quan e

a Australian Key Centre for Microscopy and Microanalysis, The University of Sydney, New South Wales 2006, Australiab ARC Centre of Excellence for Design in Light Metals, The University of Sydney, New South Wales 2006, Australiac Department of Metallurgy and Materials Engineering, University of Indonesia, Kampus UI Depok 16424, Indonesia

d Defence Metallurgical Research Laboratory, Kanchanbagh, Hyderabad 500 058, Indiae Wescast Industries Inc., Brantford, ON, N3T 5L8, Canada

Received 16 August 2007; received in revised form 17 December 2007; accepted 31 December 2007Available online 21 March 2008

Abstract

Analytical transmission electron microscopy was used to characterize precipitation processes in a model Al–4.0Cu–0.3Mg (wt.%)alloy aged at 200 �C. The evolution of microstructure was more complex than previously reported and involved processes common bothto the binary Al–Cu and ternary Al–Cu–Mg systems. We report precipitation of a novel orientation of the common intermediate phaseh0, which we have designated h0II. Isomorphous with the well-known body-centred tetragonal Al2Cu phase that occurs in Al–Cu-basedalloys, we observed the following novel orientation: f001gh0II

==f1 10ga, h100ih0II==h001ia. A trace of Mg was associated with the h0II ori-

entation and we suggest that this is a consequence of the role that Mg atom clusters play during nucleation. For the first time, we alsoreport the occurrence of the r phase (Al5Cu6Mg2) in a new orientation analogous to that of h0II, designated as rII:f001ga==f001grII

andh110ia==h010irII

. These results are discussed in terms of microstructural design for age hardening in Al–Cu–Mg base alloys, much ofwhich involves the h0 and r precipitate phases.� 2008 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: TEM; Age hardening; Precipitation; Microstructure; Cluster assisted nucleation

1. Introduction

Alloys of the Al–Cu–Mg system were among with firstheat-treatable, high-strength Al alloys developed for struc-tural applications, and are widely used because of theirstrength, ductility and superior creep strength at elevatedtemperatures [1]. The composition Al–4.0Cu–0.3Mg(wt.%) has received considerable attention as it has servedas the base composition for several light alloy developmentprojects. For example, Al–Cu–Mg–Ag alloys have beeninvestigated extensively as scientific model systems forunderstanding novel clustering and precipitation processes[2,3] as well as the development of high-strength, creep-

1359-6454/$34.00 � 2008 Acta Materialia Inc. Published by Elsevier Ltd. All

doi:10.1016/j.actamat.2007.12.046

* Corresponding author. Address: Australian Key Centre for Micro-scopy and Microanalysis, The University of Sydney, New South Wales2006, Australia. Tel.: +61 2 9351 4215.

E-mail address: [email protected] (S.P. Ringer).

resistant alloys [4,5]. Moreover, Li additions to the Al–Cu–Mg–Ag alloys have resulted in a family of alloysknown as the Weldalite series, which include compositionsthat have strengths that rate amongst the highest of all con-ventionally produced Al alloys [6]. Therefore, an under-standing of the precipitation processes in the base Al–Cu–Mg ternary composition is fundamentally importantin developing a detailed understanding of the effects ofalloying.

The nature of the strengthening precipitate phases inindividual Al–Cu–Mg alloys is determined by the solutelevels and the ratio of Cu:Mg, as well as the thermome-chanical processing parameters. There is generally a signif-icant increase in the hardness and strength levels attainedwith an increase in Cu:Mg ratio [7]. Fig. 1 shows the hard-ness–time curves for the Al–4.0Cu–0.3Mg (wt.%) basealloy under consideration here. Hardening appears toproceed in a single stage, although the rate of hardening

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Fig. 1. Hardness–time curves of the experimental alloy aged at 150, 200and 250 �C. AQ: as-quenched.

2148 S.P. Ringer et al. / Acta Materialia 56 (2008) 2147–2160

is higher after approximately 0.4–0.5 h at 200 �C. Alsonoteworthy is the relatively gradual decrease in hardnessafter attainment of the peak value, since a slow over-ageingrate is often associated with good creep properties.

Previous investigations on the precipitation processes inthis alloy were mostly done using X-ray diffraction. Sincethe Al–4.0Cu–0.3Mg (wt.%) composition lies in thea + h+S region of the ternary phase diagram [8], precipita-tion processes common to both the binary Al–Cu and theternary Al–Cu–Mg system are considered to dominatethe development of microstructure in this alloy. The mainprocesses in the decomposition of the supersaturated solidsolution (SSSS) of these alloys include

SSSS! Clusters! GP zones! h00 ! h0 ! h

! GPB zonesþ S! S

! X! h

The h phase is a body-centred tetragonal (bct) (I4/mcm,a = 0.607 nm, c = 0.487 nm) equilibrium form of Al2Cu,and details of the transformations between Guinier–Pres-ton (GP) zones and the metastable h00 and h0 have been wellstudied [1,9]. The h0 phase is stable over a wide range oftemperatures and times and plays a significant role on thehardening process. It also possesses a bct structure (I4/mcm, a = 0.404 nm, c = 0.580 nm) [10–12], forming as rect-angular or octagonal plates on the {100}a planes of thematrix phase. In the case of the ternary alloy, the equilib-rium S phase (Al2CuMg) is widely regarded as a base-cen-tred orthorhombic structure (Cmcm, a = 0.400 nm,b = 0.923, c = 0.714 nm) [13], which forms as laths parallelto matrix h100ia directions, possessing a {210}a habitplane. A number of metastable allomorphs of the S phasehave been mentioned in the literature, such as S00 and S0,however, these are ignored here since the structural andchemical characteristics that distinguish these from S phase(e.g. [14]) do not seem to require a distinct nomenclature, atleast in the alloy studied here.

The X phase (Fmmm, a = 0.496 nm, b = 0.859,c = 0.848 nm) [15], which forms as thin hexagonal-shapedplates on matrix {111}a planes, is an important strength-ening precipitate in creep-resistant Al–Cu–Mg–Ag alloys[4] and certain Al–Cu–Mg–Ag–Li alloys [6]. This phasehas also been observed in Ag-free compositions such asAl–4.0Cu–0.3Mg [16].

The precipitate r (Al5Cu6Mg2, Pm3, a = 1.999 nm),which seems to be an equilibrium phase in regions of highCu and Mg concentrations in the ternary phase diagram,was first reported in an Al–Cu–Mg alloy by Schuelleret al. [17] in their investigation of an Al–4.3Cu–2Mg/SiCwhisker cast composite after T7 heat treatment. Theseauthors recognized that this was nanoscale precipitationof the intermetallic phase Al5Cu6Mg2, first investigatedby Samson [18], who reported the stoichiometry and crys-tal structure. They also reported that the precipitateoccurs with a cube-on-cube orientation relationship suchthat

f001ga==f00 1grh100ia==h100ir

Since that work, there have been numerous papersreporting the significance of r in precipitation processesin Al–Cu–Mg alloys [19,20]. The potential for r phase toincrease the temperature range for application of precipita-tion-hardenable Al–Cu–Mg alloys and composites was wellstudied by Schueller et al. [17]. The work of these authorsplaced a lot of emphasis on the high thermal stability ofr phase, which they reported could extend as high as250 �C [17]. They also noted that the optimum size and dis-persion of this precipitate had not yet been successfullydetermined to maximize the strength increment due to pre-cipitation hardening.

In a study, of precipitation processes in an Al–4.0Cu–0.3Mg (wt.%) alloy, transmission electron microscopy(TEM), selected-area electron diffraction (SAED), micro-beam electron diffraction (MBED) and atom probe fieldion microscopy (APFIM) were used to show that the bin-ary Al–Cu alloy precipitation reaction is dominant at lowertemperatures [2]. Arising from an interest in the relativecompetition between these precipitation reactions, we haveused TEM, SAED and MBED to carefully characterize thestructure, morphology and orientation of precipitationwhen this important base alloy is aged at 200 �C.

2. Experimental procedure

A scientific alloy of nominal composition of Al–4.0Cu–0.3Mg (wt.%) was prepared from elemental components ofhigh purity (>99.99%) by induction melting and chill cast-ing into a cast iron mould. The ingot was homogenized at500 �C for 100 h, scalped and hot rolled to 5 mm strip forbulk Vickers hardness testing (5 kg load) and to 0.5 mmsheet for the preparation of thin foils for electron micros-copy. Samples were solution treated for 1 h at 525 �C in

S.P. Ringer et al. / Acta Materialia 56 (2008) 2147–2160 2149

a salt bath, cold water quenched and immediately aged inoil baths for various times at 200 �C.

Thin foil specimens for TEM were prepared using stan-dard techniques and studied in conventional LaB6 analyti-cal microscopes operating at 120 and 200 kV. Some energydispersive X-ray spectroscopy (EDXS) was undertaken in aVG HB 501 dedicated scanning transmission electronmicroscope operating at 100 kV and using a cold field emit-ter. The chemical composition of the phase was carried outusing EDXS in a dedicated scanning transmission electronmicroscope with a nominal probe size of less than or equalto 1 nm. Diffraction simulations of precipitate crystal struc-tures were performed using the software Diffract and Desk-top Microscopist.

3. Results

3.1. Precipitation sequence

The sequence of precipitation in the Al–4Cu–0.3Mg(wt.%) base alloy at 200 �C is provided in Fig. 2. TheSAED patterns are correctly oriented with respect to thebright-field (BF) images and so provide a basis for traceanalysis of the precipitate structures. The BF TEM micro-graphs were recorded near the h001ia orientation and pro-vide an overview of the general microstructure togetherwith the relevant typical h001ia SAED patterns. The BFimage recorded from the specimen aged for 20 min revealsa mottled contrast in a matrix that is otherwise devoid ofany precipitation. Correspondingly, no diffraction effectsthat could be associated with decomposition were observedin the corresponding SAED.

Fig. 2b reveals that a complex precipitate microstructurehas developed after ageing for 10 h. On the one hand, fineand uniform dispersion of precipitates exhibiting {100}a

traces were observed, with two orthogonal variants clearlyresolved in Fig. 2b, suggesting h0 platelets; this was con-firmed in the corresponding SAED pattern. Furthermore,Fig. 2b contains evidence of precipitates faceted along the{11 0}a planes (labelled as 1). This phase seems to havebeen previously unreported and a detailed characterizationof this precipitation is provided below. Nucleation of this{11 0}a precipitate was concomitant with a steep increasein the hardening rate, towards peak hardness (Fig. 1). Alsoobserved at this stage of ageing were S phase and X phaseprecipitates. Contrast from both of these phases, respec-tively, is available by {210}a precipitate traces (labelledS) and from the X plate-shaped precipitates exhibitingthickness-fringes inclined to the electron beam (arrowed,labelled 2).

An h001ia BF TEM image and corresponding SAEDpattern of the peak hardness microstructure is providedin Fig. 2c. Image trace and diffraction effects consistentwith a fine and uniform dispersion of h0 precipitates wereagain observed. Other precipitates labelled as ‘‘1” were alsoobserved as described above. Finally, S phase precipitateswere observed end-on and, as in the case of Fig. 2b, thick-

ness-fringe contrast from X phase precipitates was visible.Reflections at 1=3 and 2=3 gf2 2 0ga were observed in the cor-responding SAED pattern, consistent with diffraction fromX phase precipitates that are oriented at an inclination tothe electron beam [15].

A gradual precipitate coarsening reaction was observedto take place on subsequent ageing. The h001ia BF TEMimage of the over-aged microstructure is provided inFig. 2d after ageing for 200 h at 200 �C. The alloy micro-structure retains uniform dispersions of both h0 and Xphase precipitates. Also observed with increasing frequencyin these over-aged microstructures were S phaseprecipitates.

3.2. Identification of {110}a precipitate

As mentioned, TEM images of the microstructure of thebase alloy following ageing at 200 �C for 10 h revealed sig-nificant precipitation of a phase with habit planes parallelto the {110}a planes (Fig. 2b, arrowed and labelled 1).Two variants of this phase are apparent from the h001iaorientation. The composition, crystal structure and mor-phology of this phase were investigated in detail.

3.2.1. Microanalysis

The VG scanning transmission electron microscopewas used to conduct X-ray microanalysis of the Al–4Cu–0.3Mg (wt.%) alloy aged for 10 h at 200 �C at a ser-ies of points across the precipitate and surroundingmatrix (Fig. 3a). Representative results are provided inFig. 3b for the whole spectrum and for the Cu Ka andMg K peak in detail (see inset, Fig. 3b). A precipitatehaving interfaces oriented at 45� to the h001ia directionis clearly shown in Fig. 3a. Five points were selectedacross the precipitate as the positions for scanning trans-mission electron microscopy microanalysis. The enlargedCu Ka peaks (Fig. 3b) indicate that the Cu content islow at position 1 (near the precipitate/matrix interface)and increases significantly inside the precipitate beforereducing again at position 5, in the adjacent matrix onthe other side of the precipitate. In respect of the MgK spectrum, it is interesting to note that there is a dis-tinct Mg K peak at position 1, near the precipitate/matrix interface. This Mg K count diminishes at position2 and is absent at position 3 (in the middle of the precip-itate). A distinct Mg K peak was recorded at position 4(at the precipitate/matrix interface), whereas no peak wasdetected at position 5 in the adjacent matrix. Careful cal-ibration and quantification of the peaks revealed that Mgwas present at positions 1 and 4, with the averageAl:Cu:Mg atomic ratio of �1.6:0.8:0.04. The Mg contentat these positions is considerably lower than that presentin the S phase, suggesting that the composition of thephase is much closer to that of the h0 (Al2Cu). However,the fact that Mg was detected at the precipitate/matrixinterface suggests that same enrichment of Mg is occur-ring there.

Fig. 2. Bright-field transmission micrographs and the corresponding SAED patterns showing the microstructural evolution of the Al–4Cu–0.3Mg (wt.%)alloy close to h001ia aged at 200 �C for (a) 20 min, (b) 10 h, (c) 30 h and (d) 200 h.


3.2.2. Crystallography

A system of MBED studies in conjunction with large-angle tilting experiments was undertaken to reveal the mor-phology and crystal structure of the precipitates exhibitingthe {110}a trace. Fig. 4a shows a BF TEM image of thebase alloy aged at 200 �C for 10 h recorded from near theh001ia zone axis. Trace analysis revealed various precipi-tates were imaged: (i) three {110}a precipitates (arrowedand labelled 1, 2 and 3); (ii) an edge-on h0; (iii) two end-on S precipitates (S1 and S2); (iv) a side-on S precipitate(S3); and (v) an inclined X precipitate.

Precipitate 1 has a rectangular cross-section, whereasprecipitates 2 and 3 possess nearly square cross-sections.

All possess four {11 0}a facets. The strong structure factorcontrast of the phase at this orientation suggests that theseprecipitates extended parallel to the beam direction. How-ever, the contrast for precipitate 1 is lower than that forprecipitates 2 and 3. Unfortunately, precipitates 2 and 3formed in contact with and overlapped adjacent S phaseprecipitates, making it impossible to record their unconvo-luted MBED patterns. An MBED pattern from precipitate1 is provided in Fig. 4b. This diffraction pattern is strikingbecause it is identical in symmetry and spacing to that ofthe h0 phase (I4/mcm, a = 0.404 nm, c = 0.580 nm) [10–12]. What is significant is the observation that the orienta-tion relationship between this phase and the a matrix was

Fig. 3. (a) Bright-field scanning transmission electron micrograph of a {110}a precipitate in the specimen aged at 200�C for 10 h close to h001ia.(b) Typical EDXS spectra from the five locations marked in (a). The insets provide enlargements of Mg K and Cu K regions of the spectra.


rotated 45� from the usual orientation of h0 in binary Al–Cu alloys and this was confirmed by computer diffractionsimulation in a 45� rotated orientation also provided inFig. 4b. Moreover, the edge-on h0 precipitate, labelled inFig. 4a yielded a MBED pattern as shown in Fig. 4c, whichmatches precisely with the diffraction simulation (also pro-vided in Fig. 4c) based on the above-mentioned tetragonalstructure in the usual orientation relationship between aand h0. The experimentally obtained MBED patterns ofthe end-on S1 and the side-on S3 precipitates are providedin Fig. 4d and e, respectively, and match well with the sim-

ulated diffraction pattern based on the Perlitz–Westgrenstructure (orthorhombic, Cmcm, a = 0.404 nm, b = 0.923nm, c = 0.714 nm) [13].

To identify the structure and morphology of the precip-itates exhibiting the {110}a traces, large-angle tiltingexperiments were extended to other low-index zone axes.Tilting the specimen as oriented in Fig. 4 through 45�towards the [01 1]a direction resulted in a low contrastimage, as shown in Fig. 5a. Precipitates 1, 2 and 3 wereinclined to the beam direction. The contrast suggested thatprecipitate 1 possessed a prismatic, or short-platelet

Fig. 4. (a) Bright-field transmission electron micrograph recorded close to h001ia from the sample aged for 10 h at 200 �C. The image includes an edge-onh0 platelet, two end-on S (S1 and S2), a side-on S (S3), an inclined X and three {110}a precipitates (1, 2 and 3). Also provided are microbeam electrondiffraction patterns and the corresponding simulated patterns recorded from (b) precipitate 1; (c) the h0; (d) the S1 precipitate and (e) the S3 precipitateshown in (a).


morphology. Precipitates 2 and 3 seemed to also have aprismatic morphology, although the interference with theS phase made their contrasts less clear than that for precip-itate 1. This orientation of precipitate 1 did not diffractclearly and so the MBED pattern reveals only a-matrixeffects (Fig. 5b). The h0 precipitate remained edge-on in thisorientation and gave rise to a MBED pattern, as providedin Fig. 5c. Inspection reveals that this matches well with thecorresponding simulated patterns. The X phase was alsoviewed edge-on at this orientation. The MBED pattern ofprecipitate S1 is provided in Fig. 5d, which matches wellwith the corresponding simulation. In this h01 1ia direc-

tion, the trace of the S3 precipitate was perpendicular tothe trace of the edge-on h0 precipitate.

Further high-angle tilting towards the h111ia zone axisrevealed that precipitates 1, 2 and 3 were bound by two par-allel planes parallel to {110}a (Fig. 6a). This is consistentwith the prismatic morphology assumed above where thetwo parallel planes are the broad {01 0}a planes of the pre-cipitates. The diffuse line contrast in precipitate 1 (arrowed,Fig. 6a) is one edge of the prism. The adjacent S precipitateshowed similar morphology at this orientation. Signifi-cantly, the MBED pattern recorded from precipitate 1(Fig. 6b) could also be indexed using the h0 structure in a

Fig. 5. (a) The same field of view as in Fig. 4a recorded from the h101ia zone axis. Microbeam electron diffraction patterns and the correspondingsimulated patterns as recorded from (b) precipitate 1; (c) the h0 precipitate and (d) the S3 precipitate shown in (a).


rotated orientation relationship. When Fig. 6a was recorded,the sample had already been under constant exposure to theelectron beam for some time, causing some radiation dam-age to the specimen. Therefore, the h0 and S3 precipitateswere barely revealed in this examination, and the focus ofthe investigation became centred on precipitate 1.

Further high-angle tilting towards the h112ia directionrevealed that the {110}a precipitates exhibited strong con-trast (Fig. 7a), verifying a prismatic morphology. TheMBED pattern from precipitate 1 (Fig. 7b) also providedexcellent matching with the simulation based on the h0

structure in a rotated orientation. The X phase was alsoedge-on at this particular orientation.

The diffraction patterns from the large-angle tiltingexperiments has so far been satisfactorily described as con-sistent with the h0 crystal structure in a novel orientationrelationship, namely

f110ga==f010gh0 ; h001ia==h100ih0We have designated this phase as h0II in order to convey

that this phase is, in fact, h0 but that it occurs in a novel ori-entation relationship. To confirm further the structure andsymmetry of this phase, careful tilting experiments wereperformed in the electron microscope to other low-indexzones. A compilation of the experimental MBED patternstogether with the stereographic projection and symmetries

Fig. 6. (a) The same field of view as in Fig. 4a recorded from the h111ia zone axis. (b) Microbeam electron diffraction pattern recorded from precipitate 1arrowed in (a) and the corresponding simulated pattern, where (�) a; (�) precipitate 1; and ( ) double diffraction.

Fig. 7. (a) The same field of view as in Fig. 4a recorded from the h121ia zone axis. (b) Microbeam electron diffraction pattern recorded from precipitate 1arrowed in (a) and the corresponding simulated pattern, where: (d) a; (s) precipitate 1; and ( ) double diffraction.


is provided in Fig. 8. All the MBED patterns recordedexperimentally matched with the simulations based onthe h0II structure.

3.3. Precipitation of r phase

We also observed precipitation of r phase during ageingat 200 �C and significantly we also report the formation of

this phase in two different orientations. The BF TEMimage provided in Fig. 9a was recorded near h00 1ia andreveals the trace and contrast well-known from r phase,labelled rI [17]. Also visible are numerous platelets of h0

indeed in the edge-on orientation with traces parallel to{001}a as well as in the plane of the image. Also visibleare S phase precipitates. Attention is drawn to the cuboi-dal-like precipitates that exhibit strong diffraction contrast

Fig. 8. Summary of the microbeam electron diffraction patterns and the stereographic projection showing the orientation relationship between the {110}a

phase (h0II) and the a matrix.


and which have edge planes parallel to {11 0}a. For heuris-tic purposes, these are labelled rII. One such precipitate islabelled in Fig. 9a and two other examples are arrowed inthe image, but not labelled. Fig. 9b is a MBED patternrecorded from the precipitate labelled rI in Fig. 9a. It isconsistent with the archetypal diffraction pattern from rin the usual cube-on-cube orientation relationship.Included in Fig. 9b is the corresponding MBED simulationbased on the crystal structure proposed by Samson [18] andthe cube-on-cube orientation relationship reported bySchueller et al. [17]. The experimental and simulated pat-terns match precisely. Fig. 9c provides the MBED patternfrom the precipitate labelled rII. It could be indexed in aself-consistent manner using the structure of r proposed

by Samson [18] in a new orientation relationship rotated45� from that previously reported. The new orientation issuch that

f001grII==f001ga

h001irII==h110ia

A system of MBED studies in conjunction with large-angle tilting experiments was undertaken and results areprovided in Fig. 10 for all major low-index zone axes viaa rII/a stereogram. All experimental MBED patterns werematched precisely by computed diffraction simulationusing the rII structure and orientation used to providethe match attained in Fig. 9.

Fig. 9. (a) Bright-field transmission electron micrograph of the specimen after ageing for 200 h at 200 �C recorded near h001ia of r phase. (b,c)Microbeam electron diffraction patterns and the corresponding simulated patterns recorded from rI and rII precipitates, respectively.


We have used the projected traces and image projectionsto reconstruct the rII morphology and the three crystallo-graphically equivalent variants are sketched in real space inFig. 11, summarizing the precipitate morphology, orienta-tion and symmetry properties.

4. Discussion

According to the present experimental observations,decomposition of the SSSS in the Al–4Cu–0.3Mg alloyduring ageing appears more complex than previouslyreported [2,4,7,16,20] and involves the following precipita-tion processes:

SSSS! Cluster stage! GP zones! h00=h0 þGP zones

! h0 ! h! GPB zonesþ S

! GPB zonesþ f110ga faceted phaseþ S! S

! X! h! r

4.1. Precipitation of h0I and h0II

Microanalysis using scanning transmission electronmicroscopy EDXS of the h0I and h0II precipitates did not dis-tinguish a difference in their composition, with results con-verging on Al2Cu. The trace and symmetry analysis also

Fig. 10. A schematic illustration of the identification and morphology of r phase with a novel orientation relationship with respect to the a matrix.(a) Stereogram and (b) real space.


confirms that their crystal structures are identical (tetrago-nal, I4/mcm, a = 0.404 nm, c = 0.580 nm) [10–12]. How-ever, the h0II form occurs in a different orientationrelationships with the matrix, such that f1 10ga==f010gh0 ;h001ia==h100ih0 .

In fact, the possibility of h0 forming in this orientationwas proposed previously by Silcock et al. [9] in a detailedX-ray diffraction study of binary Al–Cu alloys. Theseresearchers observed an irregularity of the (1.01 nm)a spotin X-ray diffraction patterns and proposed that a smallproportion of h0 could hold an alternative orientation rela-

tionship as described here for h0II. However, they could notconfirm the presence of h0 with this orientation in the bin-ary Al–Cu alloys. Phenomenologically, it would be seenthat the presence of Mg is required for the nucleation ofthis orientation.

Based on our large-angle tilting experiments, the mor-phology of the h0II phase can be described as a prism orshort platelet. The precipitate morphology was consideredin further detail by examining the symmetry characteristicsof the proposed morphology. By using intersection pointgroup symmetry analysis [21], and by assuming that the

Fig. 11. Schematic illustration of the proposed morphology for the h0II

phase.


precipitate possesses the proposed tetragonal structure, it isrevealed that the shape of the h0II precipitates may be bestdescribed by combinations of a {110}a prism parallel tothe h1 10ia 2-fold axis, together with a pinacoid normalto the h110ia 2-fold axis. This is illustrated in Fig. 11.The observed morphology is thus consistent with the sym-metry requirements imposed by the tetragonal structuralmodel.

To verify further the proposed morphology of the h0II

phase, the misfit of this form was assessed across the prin-cipal h0 axes and results are summarized in Table 1. It isclear that the misfit at the ½100�h0II

and ½001�h0IIdirections

is much lower than that for the ½010�h0IIdirection. This

may lead to a preferred growth direction along ½100�h0II

and ½001�h0II. The cube planes will form at directions normal

to ½010�h0II, and this will end up in the shapes of short plate-

lets or prisms. Inspection of the TEM images provided herereveals that this assessment is consistent with the observedprecipitate morphology.

It is also appropriate to make some remarks about thenucleation of the h0II form. The scanning transmission elec-tron microscopy EDXS analysis revealed that Mg wasclearly detected in association with the h0II orientationalvariant. Given the fact that h0 seem never to be observedin this orientation in binary Al–Cu alloys, it seems likelythat the presence of Mg is essential for the nucleation ofh0 in this particular orientation relationship. To assess thispossibility, the misfit between h0II and the matrix was com-pared with the misfit between h0 and the matrix (Table 1). Itis clear that the h0I phase possesses a large misfit with the

Table 1Misfit between the h0I and h0II orientations and the a matrix

h0I h0II

Orientationrelationship

Misfit(%)

Orientationrelationship

Misfit(%)

½100�h0==½100�a 0 ½100�h0II==½001�a 0½001�h0==½001�a 43.4 ½001�h0II==½110�a 1.4½010�h0==½010�a 0 ½010�h0II==½�110�a �41.4

matrix only along the c-direction of the structure. In theother directions, the h0I form is highly coherent with thematrix. On the other hand, the h0II phase has a large misfitalong [010]h0II as well as a significant misfit along [00 1]h0II.This misfit in two orthogonal directions suggests that theremay exist a higher nucleation barrier for the nucleation ofthe h0II orientation than for h0I. Some means of accommoda-tion of this misfit would be needed to favour the h0II orien-tation to nucleate in the relatively uniform dispersion aswas observed.

Clustering of Mg atoms, which have an atomic volume�29% larger than Al atoms, during early stages of ageingmay give rise to an extensional distortion of the a matrix.This distortion does not relieve the misfit of h0, which isalso extensional, but seems to be appropriate in relievingthe compressive misfit of the h0II phase. This feature appearsto explain why pre-precipitate clustering of Mg atoms mayfacilitate nucleation of the h0II phase.

4.2. Precipitation of rI and rII

The present identification of the r phase in Al–1.7Cu–0.3Mg (at.%) base alloy is the first report of this phase inan Al–Cu–Mg alloy with this solute level and ratio. Previ-ous work had referred to precipitation of this phase in Al–Cu–Mg alloys with compositions in higher solute regionsof the Al–Cu–Mg phase diagram such as the a + S anda + S + T phase fields. A further point of interest in thepresent work was the observation of two orientation rela-tionships for r phase, which we have designated as rI

and rII,

f001grI==f00 1ga; h100irI==h100iaand

f010grII==f1 00ga; h00 1irII==h001iaAs with the observation of multiple orientations of h0, it

was considered important to refrain from introducingentirely new precipitate designations. Indeed it is theauthors’ view that there is an excess of precipitate phasedesignations in the Al alloy literature. Where the precipi-tate composition and crystal structure are identical, thereseems to be no need to introduce a new designation. Evenslight variations in lattice parameters, which do not alterthe symmetry of the precipitate crystal should not requireunique designations. Since the phases may exhibit multipleorientations, symmetry usually requires that the phasesexhibit a variety of morphologies. Therefore, the subscriptsI and II have been used in this work only where it wasthought to be important to distinguish the orientation.

The r phase is widely regarded as an equilibrium precip-itate. The observation of multiple orientation and mor-phologies for equilibrium precipitates in Al alloys isincreasingly common. The h phase (Al2Cu) phase in Al–Cu-based alloys may form in 159 orientations relative tothe a-matrix [22]. Similarly, Gjonnes and Simensen [23]identified nine unique orientations for the MgZn2 phase


in Al–Zn–Mg alloys, which would produce a similar largenumber of total orientations when all of the orientationvariants are considered. Following the initial observationthat the S phase (Al2CuMg) exhibits 2–4� rotations aroundh001ia such that the habit plane of the lath is not alwaysparallel to {210}a, Kovarik et al. [14] recently proposedthat these orientational variants follow an invariant lineanalysis. The present observations of rI and rII seem con-sistent with these characteristics. Perhaps the only distin-guishing feature of r phase precipitation is that both theorientations are rational, whereas the phases mentionedabove exhibit one or more irrational orientations. Thismay be a consequence of the cubic structure, and the latticeparameter matching. It also seems plausible that the r pre-cipitate may exist in as yet unobserved irrational orienta-tions. We also note the reports of Mg segregation to a/Al3Sc precipitate–matrix interfaces in a Al–Sc–Mg alloyand the suggestion that this segregation modifies Al3Sc pre-cipitate morphology [24]. We propose that similar mecha-nisms are operating in these cases, i.e. Mg-rich solutecluster formation from the quench and during the earlystages of ageing leading to a modified nucleation kineticsand, in certain circumstances such as here, thermodynam-ics so that the nature (orientation, morphology) and disper-sion (size, number density, etc.) of the precipitates ischanged.

The observations of multiple orientations of h0 is inter-esting because there is less, if any, precedence for multipleorientations of metastable precipitates. In most cases,metastable intermediate precipitates have rather limitedregions of phase stability. They usually possess rational ori-entation relationships, which involve limited misfit strainsin one or more directions. Therefore, variations in orienta-tion are difficult because, unlike equilibrium precipitateswhere volume free energy dominates, the contribution fromstrain energy may play a significant role. The chemical con-tribution arising from the nucleating defect is also likely tobe significant, though these considerations are beyond thescope of our paper. It is significant that the nucleation ofX (Al2Cu) phase in Al–Cu–Mg(–Ag) alloys has some par-allels to the h0I/h

0II precipitation observed here, since Mg

and Mg–Ag clusters stimulate the formation of an interme-diate form of Al2Cu along {111}a planes (X) rather thanalong {001}a planes. This is a case where the intermediateprecipitate forms with the same composition but differentorientation and crystal structure as a result of specificnucleation events. That situation may be only a slight var-iation of the h0I/h

0II situation here in that the crystal struc-

tures of h0 and X are obviously different, whereas h0I/h0II

are distinguished only on the basis of orientation.These results show that the potency of h0 as a strength-

ening agent may also now be considered in terms of a sec-ond significant orientation. Similarly, the question of theeffectiveness of r in providing strength and particularlycreep-resistance now may be considered in terms of thenew rII orientation. We also point out that the proposedaction of the Mg-rich solute clusters in assisting the nucle-

ation of precipitate phases in new orientations and newmorphologies represents a significant extension to the con-cept of cluster assisted nucleation [25], since the examplesusually discussed in that context relate either to the forma-tion of wholly new precipitate phases with different crystal-lography and often different chemistry or the formation ofknown phases in refined dispersions.

5. Conclusions

(1) The precipitation sequence in the Al–4Cu–0.3Mg(wt.%) alloy during isothermal ageing at 200 �C wasfound to be as follows:

SSSS! Clusters! GP zones! h00=h0 þGP zones! h0

! h! GPB zonesþ S

! GPB zonesþ f110gafaceted phaseþ S! S! X

! h! r

(2) The precipitation processes in this alloy involve amixture of those observed in the binary Al–Cu andthe ternary Al–Cu–Mg systems. Since this alloy hasa high Cu:Mg ratio, the precipitates common in bin-ary Al–Cu alloys dominate the microstructure.

(3) A new orientational and morphological variant of theh0 (Al2Cu, tetragonal I4/mcm, a = 0.404,c = 0.508 nm) phase was observed in this alloy nearthe peak aged condition at 200 �C. This form of h0

possesses a short platelet or prismatic morphologyand occurs in a novel orientation: f110ga==f01 0gh0 ,h00 1ia==h10 0ih0 . This form is designated as h0II,where the subscript is used only when it is necessaryto distinguish the precipitate orientations. A smallamount of Mg was observed in association with thisphase. Since the h0II orientation exhibits a higher mis-fit with the a-matrix, the Mg atom clustering is pro-posed to play a significant role in accommodatingthis misfit, and thereby favouring nucleation of h0 inthis orientation.

(4) The cubic r phase (Al5Cu6Mg2, Pm3, a = 0.831 nm)was observed in the over-aged microstructures ofthe base alloy aged at 200 �C. To our knowledge, thisis the first report demonstrating the occurrence of thisphase in such a dilute alloy.

(5) We have observed the precipitation of r phase in anew orientation. In addition to the previouslyobserved cube-on-cube orientation, we report a 45�rotated variant designated rII whereby: {110}a//{100}rII, and h001ia//h001irII.

Acknowledgements

The authors are grateful for funding support from theAustralian Research Council, which partly sponsored thiswork. Bondan T. Sofyan and Quan Guangchun are grate-


ful for Australian Aid scholarships, used to complete Ph.D.and M.Eng.Sc. thesis work at Monash University, part ofwhich related to this investigation. K. Satya Prasad isgrateful for the DMRL, Hyderabad, India for leave to takeup a visiting scientist position at The University of Sydney.The authors are grateful for scientific and technical inputand support from the Australian Microscopy and Micro-analysis Research Facility (AMMRF) node at The Univer-sity of Sydney.

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precipitation reactions in al–4.0cu–0.3mg (wt.%)...

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