This article was downloaded by: [Pennsylvania State University]On: 12 August 2014, At: 06:59Publisher: Taylor & FrancisInforma Ltd Registered in England and Wales Registered Number: 1072954 Registeredoffice: Mortimer House, 37-41 Mortimer Street, London W1T 3JH, UK

Philosophical Magazine LettersPublication details, including instructions for authors andsubscription information:http://www.tandfonline.com/loi/tphl20

Stimulation of the Ω phase in anAl–1.1 at.% Cu–0.5 at.% Mg alloy by aduplex ageing treatment involvinginitial natural ageingG. B. Winkelman , K. Raviprasad & B. C. Muddlea School of Physics and Materials Engineering , Monash University ,Victoria 3800, Australiab School of Physics and Materials Engineering , Monash University ,Victoria 3800, Australia E-mail:Published online: 23 Aug 2006.

To cite this article: G. B. Winkelman , K. Raviprasad & B. C. Muddle (2005) Stimulation of the Ωphase in an Al–1.1 at.% Cu–0.5 at.% Mg alloy by a duplex ageing treatment involving initial naturalageing, Philosophical Magazine Letters, 85:4, 193-201, DOI: 10.1080/09500830500157611

To link to this article: http://dx.doi.org/10.1080/09500830500157611

PLEASE SCROLL DOWN FOR ARTICLE

Taylor & Francis makes every effort to ensure the accuracy of all the information (the“Content”) contained in the publications on our platform. However, Taylor & Francis,our agents, and our licensors make no representations or warranties whatsoever as tothe accuracy, completeness, or suitability for any purpose of the Content. Any opinionsand views expressed in this publication are the opinions and views of the authors,and are not the views of or endorsed by Taylor & Francis. The accuracy of the Contentshould not be relied upon and should be independently verified with primary sourcesof information. Taylor and Francis shall not be liable for any losses, actions, claims,proceedings, demands, costs, expenses, damages, and other liabilities whatsoeveror howsoever caused arising directly or indirectly in connection with, in relation to orarising out of the use of the Content.

This article may be used for research, teaching, and private study purposes. Anysubstantial or systematic reproduction, redistribution, reselling, loan, sub-licensing,systematic supply, or distribution in any form to anyone is expressly forbidden. Terms &

Conditions of access and use can be found at http://www.tandfonline.com/page/terms-and-conditions

Dow

nloa

ded

by [

Penn

sylv

ania

Sta

te U

nive

rsity

] at

06:

59 1

2 A

ugus

t 201

4

Philosophical Magazine Letters,Vol. 85, No. 4, April 2005, 193–201

Stimulation of the : phase in anAl–1.1 at.%Cu–0.5 at.%Mg alloy by

a duplex ageing treatment involving initial natural ageing

G. B. WINKELMAN*y, K. RAVIPRASAD and B. C. MUDDLE

School of Physics and Materials Engineering,Monash University, Victoria 3800, Australia

(Received 4 November 2004; in final form 24 February 2005)

Initial natural ageing of a solution-treated and quenched Al–1.1 at.%Cu–0.5 at.%Mg alloy for as little as 2 h at 22� 2�C, followed by artificialageing at temperatures as low as 150�C, is observed to significantly enhance thenumber density of �-phase precipitate particles present, compared with samplesaged immediately at 150�C. The observation is significant because the metastable� phase is a particularly effective strengthening phase in such alloys, and theresults suggest potential for introducing the � phase into Al–Cu–Mg alloys, inthe absence of expensive microalloying additions, by modifying the ageingprocess. The formation of the � phase as a product of the duplex ageingtreatment is attributed to pre-precipitate clustering of Mg atoms and vacanciesthat is essential to the nucleation of the � phase, and to a thermal treatment paththat sustains such clustering and minimizes alternative paths for relief of Mgsupersaturation in solid solution.

1. Introduction

The 2000 series aluminium alloys of the Al–Cu(–Mg) system are an important classof precipitation-hardening aluminium alloys. When artificially aged at medium tem-peratures (e.g. 150�C), ternary Al–Cu–Mg alloys from appropriate fields of theternary phase diagram [1] can exhibit the equilibrium precipitate phases � (Al2Cu),S (Al2CuMg) or T (Al6CuMg4). Each of these equilibrium phases is observed toform through the nucleation and growth of appropriate precursor precipitatephases, whose identity depends upon the alloy composition and both the tempera-ture and exposure time during artificial ageing. Microalloying is known to stimulatenew metastable precipitate phases or modify equilibrium precipitation. A notableexample [2–6] is the trace addition (about 0.1 at%) of Ag to Al–Cu–Mg alloys witha high Cu-to-Mg ratio (e.g. the well-studied Al–1.7 at.%Cu–0.3 at.%Mg).Microalloying with Ag comprehensively changes the precipitation process by

*Corresponding author. Email: graham.winkelman@materials.oxford.ac.ukyNow at Department of Materials, University of Oxford, Parks Road, Oxford OX1 3PH, UK

Philosophical Magazine Letters

ISSN 0950–0839 print/ISSN 1362–3036 online # 2005 Taylor & Francis Group Ltd

http://www.tandf.co.uk/journals

DOI: 10.1080/09500830500157611

Dow

nloa

ded

by [

Penn

sylv

ania

Sta

te U

nive

rsity

] at

06:

59 1

2 A

ugus

t 201

4

stimulating the formation of a uniformly distributed metastable precipitate phase,designated � (Al2Cu), which occurs as thin, hexagonally shaped plates parallelto f111g� [2–6]. This is in contrast with a very sparse to negligible distributionof � plates observed in the equivalent Ag-free ternary Al–Cu–Mg alloy (see,for example [7]). The marked stimulation of the � phase results in alloys withimpressive mechanical properties and thermal stability [8], and this has led to exten-sive study of the nucleation, growth and crystallography of the � phase in thequaternary alloy.

Various proposals for nucleation of an apparently uniform distribution ofprecipitates of the � phase in the quaternary alloy have emerged [4, 9–11], includingthe suggestion [4] that � forms from the precursor Mg3Ag, although this has notbeen shown conclusively. It has also been proposed [10] that the addition of Ag andMg to Al–Cu alloys reduces the stacking-fault energy, thereby increasing the numberdensity of f111g� stacking-faults upon which � may nucleate. However, evidencethat the nucleation of the � phase is not classically defect stimulated in quaternaryalloys [12, 13] does not support this.

Microalloying elements Mg and Ag interact preferentially with each other,and also have strong affinity for vacancies [14–16]. Such interactions may lead tothe formation of intermediate phases or a modified defect structure, both ofwhich may potentially assist in the profuse nucleation of the � phase in the qua-ternary alloy. The suggestion that co-clusters of Ag and Mg act as precursors of the� phase [9, 11] has recently been confirmed by three-dimensional atom probe fieldion microscopy [17], in which it has been shown that such Ag–Mg co-clusters formafter as little as 5 s ageing at 180�C following quenching. As ageing continues,evolution of the � phase is observed at the sites of these pre-precipitate clusters[17]. The process of assisted nucleation appears to be aided by the tendency of Mgatoms to form disc-shaped clusters of minimum elastic strain energy parallel tof111g� [18].

The � phase does not appear to nucleate at dislocations loops, since the additionof Ag is known to inhibit the formation of dislocation loops [11]. Similarly, thosedislocations introduced by plastic deformation prior to ageing do not aid thenucleation of the � phase, as no increase in the number density or volume fractionof �-phase precipitates is observed following cold working and ageing ofAg-containing alloys [13]. In fact, plastic deformation acts to increase the concen-tration of heterogeneously nucleated �0 precipitates at the expense of �.

While combined additions of Mg and Ag are most effective in promoting theformation of the � phase, observation of the phase in ternary Al–Cu–Mg alloysclearly indicates that Ag additions are not essential. For example, Garg et al. [7]detected the presence of � among the more dominant �0- and S0-phase distributions,following quenching and artificial ageing at 250�C and 190�C of an Al–1.7 at.%Cu–0.5 at.%Mg alloy. The phase was not observed at the lower artificial ageingtemperature of 130�C, suggesting that nucleation of the � phase in ternary alloysis favoured at higher ageing temperatures. In addition, the � phase has also beendetected in small fractions in the commercial Al–Cu–Mg alloys 2024 and 2124, bothof which are Ag free [19], and in several other experimental ternary alloys [20, 21],although the origin of its presence was not discussed. Recently, it has been shownthat the number density of the � phase may be increased in ternary alloys by the

194 G. B. Winkelman et al.

Dow

nloa

ded

by [

Penn

sylv

ania

Sta

te U

nive

rsity

] at

06:

59 1

2 A

ugus

t 201

4

introduction of dislocations via cold working [22, 23], in contrast with the behaviourof quaternary alloys. These workers suggested that, in the absence of Ag, the Mg wasfree to aid in the nucleation of � in the presence of matrix dislocations.

Available evidence suggests that any process promoting clustering of Mg atoms,and the development of planar Mg(–Cu) clusters parallel to f111g� planes, may beeffective in promoting nucleation of the � phase. The role of Ag as an additive toternary alloys might be rationalized on the basis that it accelerates the aggregation ofMg atoms by retaining vacancies, and therefore aiding Mg diffusion. In the absenceof Ag, higher ageing temperatures favour formation of the � phase, perhaps becausethey aid the diffusion of Mg atoms. In the present work, it is demonstrated that aninitial period of natural ageing of quenched supersaturated ternary solid solution,followed by artificial ageing at 150�C, also stimulates an increased presence of the �phase. The initial stage of ageing at ambient temperature ensures, among otherthings, that the very earliest stages of decomposition of the supersaturated solidsolution will occur in the presence of an excess of quenched-in vacancies over anextended period of time following solution treatment. This may favour the aggrega-tion processes, such as clustering of Mg atoms that precede precipitation and appearessential to the formation of the � phase. The work presented here considers morecarefully than previously [7, 20] the origin of the � phase.

2. Experimental procedures

An alloy of composition Al–1.13 at.%Cu–0.56 at.%Mg was prepared in air by themelting of pure aluminium followed by the addition of elemental components ofhigh purity. Following casting and solidification in a graphite-lined steel bookmould, the alloy was solution treated, air cooled and rolled–annealed–rolled to athickness of 500 mm. From this material, 3mm discs were punched and placed insmall (15mm� 15mm) wire mesh baskets to be solution treated for 1 h at 525�Cbefore quenching into water. The alloy specimens were either artificially aged insilicone oil at 150�C immediately or subjected to natural ageing at 22� 2�C for aperiod of 2 h prior to subsequent elevated-temperature ageing for periods up to480 h. Alloy discs were thinned in a twin-jet Tenupol 5 polishing unit using a solutionof 33 vol.% HNO3 and 66 vol.% CH3OH cooled to �25�C. The defect andprecipitate microstructures were characterized from thin foil samples using aPhilips CM20 (200 kV) transmission electron microscope operating as necessary inbright-field, dark-field and diffraction modes.

3. Results and discussion

Electron micrographs representative of microstructural evolution during elevated-temperature ageing for 0 s, 60 s and 72 h at 150�C immediately following solutiontreatment and quenching are shown in figures 1a, b and c respectively. They are to becompared directly with the microstructures observed in samples naturally aged for2 h at room temperature prior to artificial ageing for equivalent periods at 150�C,shown in figures 1d, e and f respectively. All micrographs were recorded with the

Stimulation of � phase in Al–1.1 at.%Cu–0.5 at.%Mg 195

Dow

nloa

ded

by [

Penn

sylv

ania

Sta

te U

nive

rsity

] at

06:

59 1

2 A

ugus

t 201

4

electron beam parallel to h001i�. The microstructure of the as-quenched alloy wasessentially featureless (figure 1a), apart from a random and extremely sparse distri-bution of small dislocation loops and other quenched-in defects. The microstructurefollowing 60 s ageing at 150�C (figure 1b), contained a uniform distribution ofdislocation loops lying parallel to f011g� with a Burgers vector of 1

2h011i�. Such

loops form as a result of the thermally activated diffusion and aggregation ofquenched-in vacancies in the alloy. No precipitation was detected in the volumesbetween these loops at this stage. Following ageing for 72 h at 150�C (figure 1c), themicrostructure contained evidence of both �0- and S-phase precipitates, with boththese forms of precipitate forming preferentially in the vicinity of the dislocationloops. There is no attempt here to establish the identity of these phases in detail;they are readily recognizable from their form and crystallography, and electrondiffraction patterns were entirely consistent with the identification reported.

The microstructure following quenching and 2 h natural ageing at 22� 2�C(figure 1d), was featureless in conventional electron microscopy. However, following60 s ageing at 150�C (figure 1e), a dense distribution of dislocation loops had formed.This implies that the vacancy concentration in the solution-treated-and-quenchedalloy following 2 h natural ageing remained well in excess of the equilibrium vacancyconcentration at 150�C. The dislocation loops appeared to be larger in size andsignificantly fewer in number compared with the samples that were not exposed tonatural ageing.

Figure 1. Al–1.1 at.%Cu–0.5 at.%Mg alloys artificially aged at 150�C for (a), (d) 0 s,(b), (e) 60 s and (c), (f) 3 days immediately following quenching and (a)–(f) following roomtemperature natural ageing for 2 h. (B¼ h001i�). The arrows in (f) indicate the location of aphase not present in the absence of pre-ageing.

196 G. B. Winkelman et al.

Dow

nloa

ded

by [

Penn

sylv

ania

Sta

te U

nive

rsity

] at

06:

59 1

2 A

ugus

t 201

4

Also observed in the microstructure following 60 s artificial ageing at 150�C wasa uniform distribution of extremely fine precipitates that were not detectable in theabsence of natural ageing. Following extended ageing (e.g. 240 h), it was noted thatthere was an increased density of precipitates of the S phase between the dislocationloops in the specimens exposed to natural ageing, when compared with the speci-mens artificially aged immediately following quenching. This increase in precipita-tion between the dislocation loops appears likely to be related to the distribution ofvery fine features in these locations following natural ageing and 60 s at 150�C,although the linkage is not established in detail here.

After ageing for 72 h at 150�C, following 2 h natural ageing, the microstructureagain exhibited extensive precipitation of �0 and S phases, both of which formedpreferentially in association with the dislocation loops (figure 1f). In addition, a thirdprecipitate species was observed in sufficient number density for it to be regarded as asignificant constituent of the microstructure. Taking the form of thin platelets, par-ticles of this third precipitate phase could be observed (indicated by arrowheads infigure 1f) in association with the precipitate network, at f011g� dislocation loops.This third species of precipitate, formed only in those specimens exposed to initialnatural ageing, was further investigated to verify its identity.

The microstructures typical of samples, aged naturally for 2 h and then artifi-cially for 72 h and 240 h at 150�C, are shown in figures 2a and b respectively. Theelectron beam is, in each case, parallel to f011g�. Those additional precipitates,distinguishable from the �0- and S-phases, have a thin plate-shaped form of highaspect ratio; there are two variants of the phase parallel to the electron beam, withtraces parallel to those of f111g� planes, and two variants that are inclined to theelectron beam and appear as hexagonal plates in the projected image. One such plateis highlighted in outline in figure 2b. Figure 3a is a bright-field image, recorded withthe electron beam direction corresponding to h011i�, of a sample aged naturally for

Figure 2. Al–1.1 at.%Cu–0.5 at.%Mg alloy aged artificially at 150�C for (a) 72 h and (b)240 h following 2 h natural ageing at room temperature. Both microstructures contain addi-tional plate-like precipitates parallel to f111g� matrix planes (B¼h011i�).

Stimulation of � phase in Al–1.1 at.%Cu–0.5 at.%Mg 197

Dow

nloa

ded

by [

Penn

sylv

ania

Sta

te U

nive

rsity

] at

06:

59 1

2 A

ugus

t 201

4

2 h at room temperature and then aged artificially for 24 h at 150�C. Figure 3b is thecorresponding selected-area electron diffraction pattern from the volume of materialrepresented in figure 3a. The pattern exhibits characteristic diffraction maxima andstreaks consistent with the presence of the �0 and S phases [11, 24]. Additionalprecipitate maxima (arrowed) are observed at 1

3h220i� and 2

3h220i�. In the centred

dark-field image in figure 3c, recorded using a reflection located at 23h220i�, selected

precipitate plates, distinguishable from the �0 and S phases, are revealed in brightcontrast. These additional diffraction maxima coincide with the positions of preci-pitate reflections reported in the presence of the � phase (metastable Al2Cu) [6],which is known to form as thin hexagonal plates of high aspect ratio parallelto f111g�.

Although not the focus of the work presented here, it is interesting to considerany differences in alloy hardness inferred by the change in microstructure associatedwith pre-ageing. Microhardness curves were obtained for both immediately aged andpre-aged alloy Al–1.1 at.%Cu–0.5 at.%Mg and have been presented previously [20].In the case of the alloy aged immediately at 150�C, the alloy experienced a rapidincrease in vickers hardness to about 80HV within 60 s, followed by a hardnessplateau and then a second rise to a maximum hardness of about 103HV. It wasfound that the pre-aged alloy hardened significantly during pre-ageing for 2 h(88HV), before losing much of the hardness gained during subsequent elevated-temperature ageing (hardness reversion). However, following extended elevated-temperature ageing, the vickers hardness increased to 105HV. Hence, the maximumhardness of alloys following 2 h preageing plus artificial ageing is approximatelyequivalent to that exhibited by non-pre-aged alloys. Hence the presence of the �phase, in the number density observed in the present work, does not significantlyincrease the maximum hardness available. However, pre-ageing for 24 h followed byartificial ageing leads to an increase of approximately 20% in the maximum hardnessabove the non-pre-aged case. Nevertheless, since pre-ageing for 24 h introduces intothe microstructure both an increase in the number density of both � phase and auniformly distributed S phase between the dislocation loops, The origin of theincrease in hardness is not clear. Further work is planned to investigate theseimportant effects.

Figure 3. Al–1.1 at.%Cu–0.5 at.%Mg alloy specimens naturally aged for 2 h at roomtemperature and artificially aged for 240 h at 150�C (B¼h011i�): (a) bright-field image;(b) selected-area electron diffraction pattern; (c) centred dark-field image using precipitatereflection at 2

3h220i�.

198 G. B. Winkelman et al.

Dow

nloa

ded

by [

Penn

sylv

ania

Sta

te U

nive

rsity

] at

06:

59 1

2 A

ugus

t 201

4

The precipitate plates of � phase were invariably observed to be closely asso-ciated with rods or laths of S phase extended parallel to h100i� and apparentlynucleated heterogeneously in association with the vacancy condensation loops.Examples of the close association of these two precipitate phases are provided infigure 4. There was no evidence that the � phase itself nucleated directly on thedislocation loop, since the appearance of � phase was always associated with the Sphase and apparently somewhat removed from the initial location of the loop itself.The nature of any interaction between the � and S phases has not been examined indetail and is beyond the scope of this preliminary report. The large aspect ratios ofboth S-phase laths and rods and �-phase plates almost inevitably guarantee that thetwo will intersect if they form in the same volume of the microstructure, and theassociation noted here does not necessarily imply an influential interaction betweensuch pairs of precipitates.

Both the S phase and the �0 phase have been observed to nucleate preferentiallyon the dislocation loops that form in Al–Cu–Mg alloys, presumably partly becausein doing so they effectively eliminate a segment of the defect. However, both phasesform on different but well-defined regions of an individual dislocation loop [25],implying that the local nature of the dislocation strain field and the accommodationthat it affords the precipitate nuclei are also important factors in heterogeneousnucleation. In the case of the intermetallic S phase, the composition is Al2CuMg,and the formation of S will also be facilitated in the regions of Mg enrichmentthat arise from the segregation of Mg that accompanies vacancy aggregation anddislocation loop formation.

While the � phase does not contain Mg as a significant constituent, effectivenucleation of metastable � appears to be intimately linked to the presence of Mgatom (or Mg atom–vacancy) clusters, and in this sense the nucleation processes of

Figure 4. Electron microscopy image showing typical examples (indicated by white arrow-heads) of the frequent association of the � phase and the S phase (B¼ h011i�) in Al–1.1 at.%Cu–0.5 at.% Mg naturally aged for 2 h at room temperature and artificially aged for 240 hat 150�C.

Stimulation of � phase in Al–1.1 at.%Cu–0.5 at.%Mg 199

Dow

nloa

ded

by [

Penn

sylv

ania

Sta

te U

nive

rsity

] at

06:

59 1

2 A

ugus

t 201

4

the S and � phases are to be seen as competitive. It is thus not surprising to find thatthe two phases form in similar volumes of the microstructure, in which Mg concen-trations are locally increased. That there is a significantly increased fraction of the �phase formed in those samples given an initial short natural ageing treatment, com-pared with those subjected to immediate artificial ageing, would seem to imply thateither the pre-precipitate clusters essential to nucleation of � are significantly moredense or the clustering more stable in the samples subjected to duplex ageing. It waspreviously shown that the introduction of dislocations by cold work stimulates theformation of � phase [22, 23]. We have shown here that, in the absence of naturalageing prior to artificial ageing, the presence of dislocation loops alone is insufficientto promote �-phase formation. Further work is necessary to discern the impact ofboth dislocations and natural ageing on �-phase promotion.

It is tempting to speculate that, during natural ageing, there is a preliminaryclustering of Mg atoms and vacancies that establishes distributions of both speciesthat are initially more stable during subsequent artificial ageing. There is indirectevidence in support of this insofar as the rate of nucleation and thus number densityof dislocation loops are diminished in samples that are aged artificially followingnatural ageing. There is also evidence of very-fine-scale precipitation in matrixvolumes between the loops that is not detectable in those samples subjected toimmediate artificial ageing. This would seem to imply that the critical solute andvacancy concentrations required for nucleation are being maintained locally in amore stable fashion during the very early stages of artificial ageing. Enhanced sta-bility of Mg atom–vacancy aggregates combined with reduced site densities forheterogeneous nucleation of S phase may be sufficient conditions to permit increasedcompetitive nucleation of the � phase.

4. Summary and conclusions

The present results provide an interesting example of the manner in which a duplexthermal treatment schedule, typically involving a low-temperature (natural) ageingtreatment followed by a higher-temperature (artificial) ageing treatment, offers thepotential for usefully modifying the microstructure of often well-developed alloys[26]. In this case, the duplex ageing treatment generates significant volume fractionsof a third strengthening precipitate phase in an Al–Cu–Mg alloy conventionallystrengthened by combinations of the �0 and S phases. The form and crystallography,and limited electron diffraction evidence, indicate that this precipitate is the meta-stable � phase, which has only previously been observed in sparse quantities insuch alloys subjected to higher temperature (not less than 190�C) ageing treatmentsand in alloys microalloyed with Ag. The observation is significant for the followingreasons.

(i) The metastable � phase has been demonstrated to be a particularly effectivestrengthening phase, and a key constituent of high strength alloys with super-ior thermal stability.

(ii) The results suggest potential for introducing the � phase into Al–Cu–Mgalloys in the absence of Ag by modifying the ageing process.

200 G. B. Winkelman et al.

Dow

nloa

ded

by [

Penn

sylv

ania

Sta

te U

nive

rsity

] at

06:

59 1

2 A

ugus

t 201

4

The mechanism underlying the formation of the � phase as a product of theduplex ageing treatment is not fully understood. However, there is increasing evi-dence that pre-precipitate clustering of Mg atoms or, more particularly, Mg atomsand vacancies is essential to the nucleation of the metastable � phase. Heat treat-ment schedules that sustain such clustering and minimize alternative paths for reliefof Mg supersaturation in solid solution will be those schedules effective in promotingformation of the � phase.

Acknowledgements

This work was completed while GBW was at Monash University. KR is grateful forthe provision of a Logan Research Fellowship (Monash University). The work waspartially supported by the Australian Research Council.

References

[1] I.J. Polmear, Light Alloys: Metallurgy of the Light Metals, edited by R.W.K.Honeycombe and P. Hancock (Edward Arnold, London, London, 1981), p. 25.

[2] J.T. Vietz and I.J. Polmear, J. Inst. Metals 94 410 (1966).[3] R.J. Chester, PhD thesis, Monash University, Victoria, Australia (1983).[4] J.A. Taylor, B.A. Parker and I.J. Polmear, Metal Sci. 12 478 (1978).[5] B.C. Muddle and I.J. Polmear, Acta Metall. 37 777 (1989).[6] K.M. Knowles and W.M. Stobbs, Acta Crystallogr. B 44 207 (1988).[7] A. Garg, Y.C. Chang and J.M. Howe, Scripta Metall. Mater. 24 677 (1990).[8] I.J. Polmear and M.J. Couper, Metall. Trans. A 19 1027 (1988).[9] K. Hono, T. Sakurai and I.J. Polmear, Scripta Metall. Mater. 30 695 (1994).[10] S. Kerry and V.D. Scott, Metal Sci. 18 289 (1984).[11] S.P. Ringer, K. Hono, I.J. Polmear and T. Sakurai, Acta Mater. 44 1883 (1996).[12] A. Garg and J.M. Howe, Acta Metall. Mater. 39 1925 (1991).[13] S.P. Ringer, B.C. Muddle and I.J. Polmear, Metall. Mater. Trans. A 26 1659 (1995).[14] A. Somoza, A. Dupasquier, I.J. Polmear, P. Folegati and R. Ferragut, Phys. Rev. B 61

14464 (2000).[15] Y. Nagai, T. Honma, Z. Tang, K. Hono and M. Hasegawa, Phil. Mag. A 82 1559 (2002).[16] Y. Nagai, M. Murayama, Z. Tang, T. Nonaka, K. Hono and M. Hasegawa, Acta Mater.

49 913 (2001).[17] L. Reich, M. Murayama and K. Hono, Acta Mater. 46 6053 (1998).[18] I.S. Suh and J.K. Park, Scripta Metall. Mater. 33 205 (1995).[19] L.M. Wang, H.M. Flower and T.C. Lindley, Scripta Mater. 41 391 (1999).[20] G.B. Winkelman, K. Raviprasad and B.C. Muddle, Mater. Sci. Forum 396 1037 (2002).[21] Q. Cui, G. Itoh and M. Kanno, J. Japan Ins. Metals 59 492 (1995).[22] N. Unlu, B.M. Gable, G.J. Shiflet and E.A.J. Starke, Metall. Mater. Trans. A 34 2757

(2003).[23] N. Unlu, B.M. Gable, G.J. Shiflet and E.A.J. Starke, Mater. Sci. Forum 396–402 801

(2002).[24] S.P. Ringer, T. Sakurai and I.J. Polmear, Acta Mater. 45 3731 (1997).[25] G.B. Winkelman, PhD thesis, Monash University, Victoria, Australia (2004).[26] R.N. Lumley, A.J. Morton and I.J. Polmear, Acta Mater. 50 3597 (2002).

Stimulation of � phase in Al–1.1 at.%Cu–0.5 at.%Mg 201

Dow

nloa

ded

by [

Penn

sylv

ania

Sta

te U

nive

rsity

] at

06:

59 1

2 A

ugus

t 201

4