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Influence of equal-channel angular pressing on precipitationin an Al–Zn–Mg–Cu alloy

G. Sha a,b,*, Y.B. Wang c, X.Z. Liao c, Z.C. Duan d, S.P. Ringer a,b, T.G. Langdon d,e

a ARC Centre of Excellence for Design in Light Metals, The University of Sydney, NSW 2006, Australiab Australian Key Centre for Microscopy and Microanalysis, The University of Sydney, NSW 2006, Australia

c School of Aerospace, Mechanical and Mechatronic Engineering, The University of Sydney, NSW 2006, Australiad Departments of Aerospace and Mechanical Engineering and Materials Science, University of Southern California, Los Angeles, CA 90089-1453, USA

e Materials Research Group, School of Engineering Sciences, University of Southampton, Southampton SO17 1BJ, UK

Received 27 February 2009; received in revised form 7 March 2009; accepted 7 March 2009Available online 6 April 2009

Abstract

Processing by equal-channel angular pressing (ECAP) affects the morphology of g precipitates in an Al–Zn–Mg–Cu (Al-7136) alloy. Itis shown by transmission electron microscopy that ECAP changes the orientation of precipitates and this influences the atomic config-uration and the interfacial energy at the g/a-Al interfaces. Consequently, g precipitates adopt an isotropic growth mode and evolve intoequiaxed particles. A three-dimensional atom probe analysis demonstrates that large g precipitates formed in different numbers of ECAPpasses are of similar composition. The coalescence of smaller precipitates, rather than the fragmentation of larger precipitates, dominatesthe precipitate evolution.� 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Aluminium alloy; Equal-channel angular pressing; Precipitation; Three-dimensional atom probe; Transmission electron microscopy

1. Introduction

Ageing treatments and alloy composition engineeringare used widely to control the type, size and number den-sity of nanoscale precipitates with the objective of achiev-ing optimum precipitation hardening. In practice, themorphology of nanoscale precipitates is seldom a suitablevariable for manipulation using conventional ageing treat-ments, although the same precipitation phases, such as theg phase in Al–Zn–Mg alloys [1,2] and the h phase in Al–Cualloys [3], may have multiple morphologies. During ageingtreatments, the precipitation process follows a fixed trans-formation sequence dominated by the precipitation kineticsand thermodynamics so that the precipitates generally have

a simple orientation relationship with the matrix and thisdetermines their unique morphology, such as rods or plate-lets [2]. Manipulating the orientation and morphology ofnanoscale precipitates has long been a challenge for mate-rial scientists such that even microalloying was demon-strated recently to affect the precipitate morphologies andorientations [4].

Processing by severe plastic deformation (SPD) is an effec-tive procedure for producing metals and alloys having ultra-fine-grained (UFG) and/or nanocrystalline structures [5–7].Of the various SPD techniques currently available, equal-channel angular pressing (ECAP) is especially attractivebecause it can be scaled-up to produce bulk UFG materialsthat may be used for structural applications [5,6,8]. Earlyresearch on the processing of Al and Al alloys by ECAP gen-erally focused on the process of grain refinement [9–12].More recently, there has been a growing interest in control-ling the precipitation microstructures of Al alloys andthereby achieving a combination of strengthening from both

1359-6454/$36.00 � 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

doi:10.1016/j.actamat.2009.03.017

* Corresponding author. Address: Australian Key Centre for Micros-copy and Microanalysis, The University of Sydney, NSW 2006, Australia.Tel.: +61 2 90369050; fax: +61 2 93517682.

E-mail address: g.sha@usyd.edu.au (G. Sha).

www.elsevier.com/locate/actamat

Available online at www.sciencedirect.com

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grain refinement and precipitation hardening [13–15]. Theevidence to date suggests that processing by ECAP has a rel-atively complicated effect on precipitate evolution. Thus,conducting ECAP at room temperature generally suppressesprecipitation in Al alloys because no precipitation wasreported in as-quenched samples during the process [14,16]and there was a dissolution of pre-existing h0 precipitates inthe matrix after eight passes of ECAP [13]. On the otherhand, when the ECAP is conducted at higher temperatures,such as 473 K, precipitation is promoted in an Al–Zn–Mgalloy [17]. Experiments on an Al–Cu–Mg alloy showed thatone pass of ECAP at room temperature initiated a rotationof pre-existing large h0 precipitates, whereas four passes ledto the fragmentation of h0 plates and the formation of spher-ical nanoparticles [13]. In an Al–Zn–Mg–Cu (Al-7034) alloy,processing by ECAP at 473 K produced spherical precipi-tates by fragmentation of the pre-existing larger platelet pre-cipitates [18]. To date, only fragmentation of the largerprecipitates has been identified as a mechanism dominatingprecipitate evolution for materials containing pre-existinglarge precipitates [13,18], and accordingly it is not clear atpresent whether other mechanisms may be involved and, ifso, whether they are equally important in precipitate evolu-tion. The results to date have established that the high strainsimposed by SPD are effective in altering the precipitate orien-tations within the matrix. This potentially may provide aneffective method for manipulating the precipitate morphol-ogy and obtaining unique microstructures that are signifi-cantly different from those formed by conventional ageingtreatments. Thus, a more complete understanding of the pre-cipitation behaviour during ECAP is an essential prerequi-site for successfully manipulating the precipitationmicrostructures and attaining optimum properties in thesematerials.

This present investigation was initiated to examine theeffect of ECAP on Al-7136 alloy containing a fine and uni-form dispersion of precipitates. The specific objective wasto influence the early stages of the precipitation processand thereby alter the orientations of any small precipitates,thus effectively manipulating their morphologies. Trans-mission electron microscopy (TEM) and atom probetomography (APT) were employed to investigate the evolu-tion of precipitation, the precipitation sequence and theorientation relationships between the precipitates and thematrix. Quantitative APT analysis was used also to providechemical information on the precipitates and the matrixand the spatial distribution of precipitates during ECAP,and to yield unique information on both the structureand the chemistry of the precipitate evolution [19,20].The new information from this research provides addi-

tional insight into the precipitation mechanism occurringin SPD processing.

2. Experimental material and procedures

The Al-7136 alloy was received as extruded rods havingthe composition shown in Table 1. Prior to ECAP, the rodswere cut into short billets having lengths of 64 mm. Toevaluate the effect of an ageing treatment, some of these bil-lets were aged at 473 K for various times up to a maximumof 40 min. The remaining samples were processed by ECAPin the as-received condition at a temperature of 473 K fordifferent numbers of passes up to a maximum of eight.Since a time of �2 min was needed to heat the sample tothe pressing temperature in the ECAP die, it is estimatedthat each ECAP pass corresponds to an effective thermalduration time of �5 min. Pressing through eight passes istherefore equivalent to a time at temperature of �40 min.

The ECAP processing was performed using a hydraulicpress of 150 tonnes capacity operating at a ram speed of�7 mm s�1. The pressing used a solid die having anL-shaped channel bent through an angle of U = 90� withinthe die and with an additional angle of W � 20� represent-ing the outer angle of curvature at the intersection of thetwo parts of the channel. Samples were taken throughone, two, four and eight passes, corresponding to imposedstrains of �1, �2, �4 and �8, respectively [21]. In repeti-tive pressings, each sample was rotated by 90� in the samesense between each pass in the procedure termed route BC

[22].Disks of 3 mm diameter were punched from thin strips

with thicknesses of �500 lm cut from the ECAP samples,ground carefully to a thickness of �100 lm and then twin-jet electropolished into thin foils using an electrolyte of30% nitric acid and 70% methanol at 20 V at a temperatureof �20 �C. Characterization by TEM was performed usinga Philips CM12 operating at 120 kV and a JEOL 3000F at300 kV. The tip samples for atom probe characterizationwere prepared using a two-step electropolishing procedurewith thin bars having cross-sections of 0.5 � 0.5 mm2. Thefirst step used an electrolyte of 25% perchloric acid in aceticacid at 15 V at room temperature and the second step usedan electrolyte of 4% perchloric acid in 2-butoxethynal at20 V. Characterization by APT was performed in a localelectrode atom probe at a specimen temperature of 20 Kunder a laser pulse at an energy of 1.0 nJ and with a targetevaporation rate of 1%. This laser energy is sufficiently highfor measuring the precipitate concentration consistently[23]. The maximum separation method [24,25] wasemployed to identify and analyse small solute-rich precipi-

Table 1The nominal composition of the Al–Zn–Mg–Cu (Al-7136) alloy.

Zn Mg Si Fe Cu Zr Mn Cr Ti Al

wt.% 9.4 2.5 0.12 0.15 2.5 0.2 0.05 0.05 0.1 84.93at.% 3.78 2.71 0.11 0.07 1.03 0.06 0.02 0.03 0.05 92.14

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tates inside an analysed volume. The identifiable soluteswere set as Zn, Mg and Cu, and a precipitate is definedas containing at least 10 solute atoms, each of which isno further than 0.5 nm from the next nearest solute atom.

3. Results

3.1. Microstructure of the as-received alloy

The as-received material contained densely distributedultrafine features with a size of �5 nm and dark contrast,as shown in Fig. 1a. Several spherical Al3Zr particles, withdiameters of �25 nm, are marked with black arrows in theTEM micrograph. The selected area diffraction pattern(SADP) along a <111>a-Al zone axis (shown in Fig. 1b)contains weak diffractions from precipitates. The weak dif-fraction at 1/3 of {422}a-Al, marked with a white arrow, isfrom GPII zones [26] and the weak diffractions at 1/3 and2/3 of {220}a-Al, marked with black arrows, probablycome from the fine g0 precipitates [26,27]. The relativelyhigher intensity of GPII zone diffractions indicates thatthere is a significant GPII zone presence in the microstruc-ture. The sharp diffraction spots at ½{220}a-Al, markedwith the two large hollow black arrows in Fig. 1b, are fromthe Al3Zr particles in the matrix.

A three-dimensional atom probe characterization of theas-received material confirmed that small solute-rich fea-tures with a size <4 nm. Atom maps are provided inFig. 2 and reveal precipitation consistent with that

observed by TEM, although the size estimated from theTEM micrographs was slightly larger due to the straineffect of coherent precipitates [28,29]. After the atoms inthe matrix were removed, only closely aggregated soluteswere visible and the precipitates were then observed in avery high number density of �2.08 ± 0.07 � 1024 m�3,where they are enriched with Mg, Zn and Cu (Fig. 2b–d).

3.2. The precipitation microstructures of thermally aged

samples

After a short ageing of the as-received material at 473 Kfor 5 min (Fig. 3a), the microstructure is dominated by fineplatelets on {111}a-Al with lengths of �3–14 nm, thick-nesses of �0.6–1.2 nm and aspect ratios of length to thick-ness in the range from �3:1 to �12:1. After ageing for20 min (Fig. 3b), the platelet precipitates are �3–25 nmlong, �0.6–4 nm thick and with aspect ratios in the rangefrom �3:1 to �16:1. After ageing for 40 min (Fig. 3c), somelarge precipitates are thicker and others are longer, withlengths of �25–40 nm, thicknesses of �1.2–5 nm andaspect ratios in the range from �5:1 to �33:1.

The corresponding <110>a-Al SADPs of Fig. 3a–c,shown in Fig. 3d–f, indicate that two types of precipitates,g0 and g, coexist in the microstructures. It is known that gprecipitates can have up to 11 variants in Al–Zn–Mgalloys, designated g1–g11, with each having different orien-tation relationships with the matrix [1,2]. According to theorientation relationship observed in the SADPs in

Fig. 1. (a) TEM micrograph and (b) SADP of as-received Al–Zn–Mg–Cu (Al-7136) alloy.

Fig. 2. Element maps of an as-received Al–Zn–Mg Cu alloy sample: (a) Mg map; (b), (c) and (d) are Mg, Cu and Zn maps after removing the matrixatoms.

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Fig. 3d–f, the g precipitates in the thermally aged samplesbelong only to g2. After ageing for 5 min, the intensity ofthe g0 diffraction (marked by a white arrow in Fig. 3d) ismuch stronger than the g diffraction (marked by a blackarrow). This indicates that fine g0 precipitates are highlydominant in this sample. After ageing for 20 min, the dif-fraction intensity of g is stronger than for g0, as observedby the weak diffractions near {111}a-Al (marked witharrows in Fig. 3e). This implies the formation of a largenumber of g precipitates. After ageing for 40 min, g plateletdiffraction dominates, as shown in Fig. 3f, thereby indicat-ing an additional increase in the g volume fraction.

3.3. The precipitation microstructure in the ECAP samples

The sample taken through one pass of ECAP containedplatelet precipitates primarily having low aspect ratios of63:1, as shown in Fig. 4a. For this sample, the largest precip-

itate in the matrix was�33 ± 0.6 nm long and�12 ± 0.6 nmthick, with an aspect ratio close to�3:1. After two passes ofECAP there was a significant change in the precipitate micro-structure, most of the precipitates now being equiaxed. Afteradditional processing through four and eight passes, the pre-cipitates remain equiaxed but their average size increasessteadily with increasing number of passes, as shown inFig. 4b and c. Thus, the diameters of the largest precipitatesincrease from �37 ± 2 nm after four passes to �66 ± 4 nmafter eight passes.

The indexing of weak diffractions in a <1 10>a-Al SADPfor the one-pass sample (Fig. 5a) confirmed that theseplatelets are primarily g and g0 precipitates. However, theg precipitates are dominant because the diffraction inten-sity of g0 at near 2/3 of {220}a-Al appears weaker thanthe intensity of the g diffractions. Most of the g precipitatesare of the g2 type and only a few belong to the g3 type,according to the much weaker diffraction intensity of g3

Fig. 3. TEM bright-field (BF) images and the corresponding SADPs for samples thermally aged at 473 K for up to 40 min: (a), (b) and (c) are BF imagesand (d), (e) and (f) are SADPs after ageing for 5, 20 and 40 min, respectively.

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(marked with a white arrow in Fig. 5a). One important spe-cial feature of the g reflections after one pass is that, asindicated by the arrows, they are arced in a circumferentialdirection. This arcing is due to a rotation of the g precipi-tates relative to the matrix [13]. Interestingly, the reflectionsof Al3Zr close to ½{220}a-Al and ½{002}, as marked bythe bulky black arrows in Fig. 5a, also deviate slightly inthe lateral direction from their typical reflection positions,implying that the severe plastic deformation imposed inone pass of ECAP probably changes the correspondingplane spacing due to an elastic lattice distortion and evenfurther rotates the Al3Zr precipitates slightly from theiroriginal orientation in the matrix. It appears that, since itis difficult for the dislocations to shear through the precip-itates, there is an accumulation of dislocations at the pre-

cipitate/matrix interfaces and this is effective in alteringthe precipitate/matrix orientation.

After four passes of ECAP (Fig. 5b), most weak diffrac-tion spots of the precipitates scatter randomly, with manyspots situated on the white ring pattern of the g phase butaway from typical g2 and g3 reflections (Fig. 5c). This indi-cates that these precipitates are g phase with random orien-tations in the matrix. Thus, the accumulated strainimposed by severe plastic deformation significantly altersthe orientations of the g precipitates away from their initialpreferred orientations. The remainder of the diffractionspots situated between the rings of the g phase (marked withred rings in Fig. 5d) are covered primarily by the diffractionrings of the S-phase with a unit cell in the space group Cmcm,a = 0.400 nm, b = 0.923 nm and c = 0.714 nm [30]. It isknown that the S-phase is a high-temperature phase in Alalloys with high Cu content [31] that is situated predomi-nantly at the grain boundaries. It is clear from these observa-tions that four passes of ECAP cause significantfragmentation of the large high-temperature precipitates ofthe S-phase and increase their appearance within the matrix.

Fig. 6a, b and c shows the relevant Mg maps for samplesexperiencing one, four and eight passes in ECAP, respec-tively, and it is apparent that the precipitate microstructuredevelops significantly with increasing numbers of passes.After one pass (Fig. 6a), the precipitates exhibit a wide sizerange of �3–30 nm in their longest dimension and thesmall precipitates, having sizes in the range of �3–5 nm,are mostly blocky, with low aspect ratios of less than 1.6.Most of the larger precipitates in this condition appearelongated or are platelets with thicknesses of �5–10 nmand lengths or diameters of �15–40 nm. By comparisonwith the small solute-rich features of the as-received sampleshown in Fig. 2, the precipitates grow significantly afterone pass of ECAP and their number density decreases to�1.2 ± 0.1 � 1023 m�3, which compares with the as-received value of �2.08 ± 0.07 � 1024 m�3. In the sampletaken through four passes (Fig. 6b), only nine precipitateswere observed in the analysed volume, giving a numberdensity of �1.6 ± 0.5 � 1022 m�3, which is nearly one orderof magnitude lower than in the one-pass sample and twoorders of magnitude lower than for the as-received condi-tion. All precipitates are equiaxed after four passes andthe largest precipitates have a diameter of �20 nm, whichis smaller than the maximum diameter of �30 nm observedby TEM on a larger scale in Fig. 4b. In addition, the small-est precipitates are �8 nm in diameter, which is larger thanthe diameter of �5 nm observed in the one-pass sample.Interestingly, no small debris with a size smaller than�5 nm was observed in the microstructure.

In the eight-pass sample shown in Fig. 6c, the precipitatenumber density drops significantly to�2.3 ± 1.3 � 1021 m�3

and there is further growth of large precipitate to �40 nm.Examination by TEM (Fig. 4c) gave a similar precipitateevolution with a largest size of �70 nm and a smallest sizeof �15 nm. Again, no small debris <5 nm in size wasobserved in the analysed volume.

Fig. 4. TEM BF micrographs after (a) one pass, (b) four passes and (c)eight passes of ECAP.

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The chemical compositions of the larger precipitatesformed at different numbers of passes are listed in Table2, and it is apparent that these precipitates are enrichedwith Mg, Cu and Zn. The precipitates are not alumin-ium-free, with 7–8 at.% Al, 8–11 at.% Cu, 42–47 at.% Znand 34–39 at.% Mg, and a ratio of (Zn + Al + Cu)/Mg � 2. This contrasts with the 8.45 at.% Al, 59.05 at.%Zn and 32.46 at.% Mg of g precipitates in the Al-7108alloy (T7) with a trace Cu concentration of 0.007 at.%measured by 3DAP [32] and Zn/Mg = 2 for equilibriumg in the ternary Al–Zn–Mg system [1]. It is probable thatAl and Cu primarily occupy some of the Zn sites in thetypical g phase (Zn2Mg) and they have certain solubilitiesin the precipitation phase in the quaternary Al–Zn–Mg–Cusystem. The solute concentrations in the matrix decreasewith increasing numbers of passes, as shown in Fig. 7a,and this matches the decrease in their number densityshown in Fig. 7b. This suggests that the precipitate evolu-tion is in the growth regime.

4. Discussion

4.1. Precipitation during ECAP of the Al–Zn–Mg–Cu alloy

After one pass of ECAP at 473 K, the precipitationmicrostructure evolves from a mixture of fine coherentGP zones and g0 precipitates in the as-received alloy to amixture dominated by platelet g but with some remainingg0. In addition, only the g precipitates are present after fourand eight passes. This observation shows that g is stable at473 K in the alloy, which agrees with differential scanningcalorimetry measurements showing the thermal stabilityof g in an Al–Zn–Mg (Al-7050) alloy in the temperaturerange from 473 to 510 K [33,34]. Among the 11 variantsof g precipitates designated as g1–g11 [1,2], g2 platelets aredominant after one pass according to the SADP shownin Fig. 5a. In practice, it should be noted that g2 and g0

have similar orientation relationships with the matrix andsimilar morphologies, and it is generally believed that g2

Fig. 5. SADPs of samples after (a) one pass and (b) four passes of ECAP at 473 K, (c) the SADP of the four-pass sample superimposed on white artificialdiffraction rings of the g precipitate and (d) the SADP of the four-passes sample superimposed on red artificial rings from the S-phase. (For interpretationof the references to color in this figure legend, the reader is referred to the web version of this article.)

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platelets are formed by direct transformation of g0 [2]. Onthe basis of the precipitation observed in different ECAPsamples, it is concluded that precipitation during one passof ECAP probably follows the same sequence as in a con-ventional ageing treatment (i.e. GP zones ? g0 ? g).

4.2. The effect of ECAP on the precipitation kinetics

The precipitation microstructures formed in ECAP aresignificantly different from those formed in thermal ageingbased on the TEM and APT characterizations. According

Fig. 6. Mg maps of samples after (a) one pass, (b) four passes and (c) eight passes of ECAP.

Table 2The compositions of large precipitates formed after ECAP passes in the Al–Zn–Mg–Cu (Al-7136) alloy.

Zn (at.%) Mg (at.%) Cu (at.%) Al (at.%) (Zn + Cu + Al)/Mg

One pass 45.4 ± 0.5 38.6 ± 0.4 8.6 ± 0.1 7.4 ± 0.1 1.6Four passes 42.6 ± 0.5 38.4 ± 0.3 11 ± 0.2 8 ± 0.2 1.6Eight passes 49.5 ± 0.4 34.6 ± 0.3 8.1 ± 0.1 7.7 ± 0.1 1.9

1

10

100

1000

10000

Num

ber d

ensi

ty x

1021

, m-3

0-pass1-pass4-pass8-pass

0.00.51.01.52.02.53.0

Mg Cu Zn

Solu

te c

onte

nt, a

t% as-recieved1-pass4-pass8-pass

)b()a(0-p 1-p 4-p 8-p

Fig. 7. (a) Matrix solute concentration evolution and (b) precipitate number density evolution after different numbers of ECAP passes.

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to the relative diffraction intensities observed near 2/3 of{02 2} in the SADPs shown in Figs. 3d and 5a, the equilib-rium g precipitates are dominant after one pass of ECAPwhereas metastable intermediate g0 precipitates are domi-nant after thermally ageing for 5 min despite the fact thatone pass of ECAP has a similar effective thermal duration.Furthermore, the size of the g precipitates in the ECAPsample is much larger than in the thermally aged sample.The larger g plates after one pass of ECAP are up to�12 nm thick and �33 nm long, with an aspect ratio<3:1, whereas the g platelets in the sample thermally agedfor 5 min are <1.2 nm thick and <14 nm long, with aspectratios of �3:1–12:1. As the precipitates are g2-type, theyshould be in the form of hexagonal or rounded plates[35]. Based on these precipitate dimensions, the volumeof a large precipitate after one pass of ECAP is estimatedas more than 50 times larger than in the sample thermallyaged for 5 min. Thus, from the size differences of the largestprecipitates formed under these two conditions, precipita-tion under ECAP is estimated to occur approximately 50times faster than in conventional thermal ageing at473 K. This is in agreement with an earlier report suggest-ing that SPD promotes the precipitation process [17].

The high density of mobile dislocations plays an impor-tant role for precipitate evolution during ECAP. First, thestrain effect of mobile dislocations, in combination with theincreased interfacial areas due to dislocations shearingthrough the coherent precipitates, effectively promotes thedissolution of small metastable precipitates [36]. Secondly,the dislocations serve as fast diffusion paths for soluteatoms in the matrix, thereby permitting rapid diffusion tothe stable g precipitates to assist in their growth. Thirdly,severe shear deformation may bring two well-separatedprecipitates closer together and promote their coalescence.In an earlier report, a kinetic Monte Carlo simulation in a7xxx Al alloy without the presence of mobile dislocationsidentified coalescence as an important mechanism for lar-ger GP zone growth [37]. Thus, if a metastable precipitateis located extremely close to a neighbouring stable g precip-itate, the metastable precipitate may merge with the g pre-cipitate and this will lead to the growth of the large gplatelet at the expense of the small coherent precipitates.As a result, the growth of large precipitates will becomeproportional to the decrement in the number of smallcoherent precipitates, and this is consistent with the signif-icant drop in the precipitate number density shown inFig. 7b. Processing by the alternative SPD process ofhigh-pressure torsion is known to increase the concentra-tion of vacancies and this consequently increases the diffu-sion coefficient of solute atoms within the matrix [38]. Ifthere is a similar increase in vacancy concentration inECAP, this will assist the rapid development of the precip-itation microstructure in the Al–Zn–Mg–Cu alloy.

Fragmentation of the larger platelets plays a major rolein precipitate evolution and this process is responsible forthe spheroidization of precipitates during ECAP [13,18].If fragmentation of the small precipitates also occurs dur-

ing ECAP, the growth of the small precipitates will berestrained and essentially no small platelets will have thecapability of growing into larger precipitates. This is incontradiction with the many larger platelets, more than�25 nm long, observed after one pass and shown inFig. 4a. Accordingly, it is concluded that the fragmentationof small platelets <25 nm is not important in the precipitateevolution occurring during ECAP. Careful examination byAPT confirmed that all large platelets have a hexagonalshape, which is a characteristic feature of the perfect g platereported elsewhere [3,39]. Fig. 8 shows seven pairs of close-neighbour precipitates out of 80 precipitates inside an ana-lysed volume of the sample taken through one pass wherethese precipitates are probably undergoing coalescence.Indeed, if the growth of large precipitates is predominantlyvia coalescence rather than by the dissolution of smallerprecipitates, no small solute cluster debris is expected toexist at points remote from these larger precipitates. Thisconclusion is fully consistent with the present APT andTEM observations.

Fig. 8. Multiple view examinations of seven pairs of close-neighbourprecipitates adjacent to each other within the sample processed throughone pass of ECAP and shown in Fig. 6a: (a), (b) and (c) show threeexamples of close-neighbour small precipitates; (d), (e), (f) and (g) show asmall precipitate touching a larger precipitate.

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4.3. The effect of ECAP on g precipitate orientation and

morphology

Processing by ECAP alters the orientation of the g pre-cipitates in the matrix so that it deviates from the typicalmorphology observed in thermally aged samples. A singlepass of ECAP produces a limited rotation of the larger pre-cipitates and this rotation is about ±2� away from the ini-tial orientation, based on the curved reflections of g in theSADP after one pass in Fig. 5a. The severe shear deforma-tion arising from multiple passes of ECAP moves the g pre-cipitates further away from their original orientationsaccording to the SADP in Fig. 5b. This change in the pre-cipitate orientation introduces an important consequencefor the morphological evolution of the g precipitates.

After one pass of ECAP the g platelets have lower aspectratios compared to the g platelets in the thermally agedsamples. In particular, the smaller precipitates have aspectratios of <2:1 whereas the large platelets have aspect ratiosfrom 3:1 to 4:1, as shown in Fig. 4a. Such morphologicalcharacteristics are consistent with the orientation changesthat occur in one pass of ECAP. It is reasonable to assumethat SPD produces the same orientation change in smallplatelets as in the larger platelets. Thus, the change of pre-cipitate orientation should increase the interfacial energybetween the precipitates and the matrix due to the changein the atomic configuration at the low energy interfacesbetween the g basal planes and the matrix. Since the smallprecipitates have larger relative surface areas, they are alsomore sensitive to the change in interface energy. It followsthat, by adopting low aspect ratios, these small platelets areable to reduce their surface area and accordingly smallplatelets with low aspect ratios are favoured thermodynam-ically. It is worth noting that the change of phase interfaceenergy due to the rotation of the g precipitates by about±2� is insufficient to cause isotropic growth but that furtherrotation after multiple passes of ECAP will promote theirisotropic growth. There is also a similarity to the diffu-sional rotation of precipitates observed when grain bound-aries bypass certain coherent precipitate dispersions [40].

The equiaxed g precipitates present after multiple passesof ECAP are due to their new orientations within thematrix. Since the new orientations deviate from the typicalorientation for g platelets, the atomic configurations at theinterphase interfaces of g/a-Al are significantly differentfrom those of typical platelets. The missing low-energyconfiguration interface between the basal plane of g and{11 1} of the a-Al matrix increases the interfacial energybetween g and the matrix. Thus, the g precipitates becomemore favourable for isotropic growth and they evolve intoan equiaxed morphology. Although fragmentation of thelarger precipitates may directly reduce the average precip-itate size, the increase in their number density and thechange in their aspect ratio are effective in maintainingan overall equiaxed morphology for these fragmented pre-cipitates. Furthermore, if the fragments retained their ori-ginal orientations, it follows they would maintain their

anisotropic growth mode and it would be difficult orimpossible for them to evolve into the larger equiaxed pre-cipitates. The growth of the large equiaxed precipitatesobserved in this investigation therefore provides a cleardemonstration that the change in the precipitate orienta-tion induced by SPD is directly responsible for the g pre-cipitates in isotropic growth and their equiaxedmorphology.

This experimental investigation has revealed a signifi-cant effect in conducting ECAP at 473 K on the evolutionof the precipitate microstructure in an Al–Zn–Mg–Cualloy. First, ECAP accelerates the precipitation kineticsto about 50 times faster than in conventional ageing treat-ments. Secondly, it affects the precipitation thermodynam-ics by promoting isotropic growth and forming equiaxed gprecipitates in the alloy. To control the precipitation micro-structure effectively for an optimization of the strengthen-ing effect arising from the precipitation of fine g0, furtherresearch is now needed to obtain a more complete under-standing of the effect of ECAP at lower temperatures onthe evolution of precipitates.

5. Summary and conclusions

1. Processing by ECAP of an Al–Zn–Mg–Cu (Al-7136)alloy at 473 K changes the precipitate orientations inthe matrix. The consequent change in interfacial energyat the g/a-Al interfaces promotes g precipitates in an iso-tropic growth and leads to equiaxed particles in multiplepass samples.

2. The precipitate evolution occurring during ECAP at473 K is about 50 times faster than in conventional age-ing treatments at the same temperature.

3. The high density of mobile dislocations produced byECAP promotes the dissolution of small metastable pre-cipitates and the formation of large g precipitates bycoalescence.

4. The major solute concentrations within the matrix decreasewith increasing numbers of ECAP passes, thereby suggest-ing that precipitate evolution is in the growth regime.

Acknowledgements

The authors are grateful for scientific and technical inputand support from the Australian Microscopy & Microanal-ysis Research Facility (AMMRF) node at The University ofSydney. This project is in part financially supported by theAustralian Research Council (ARC), including ARC Dis-covery Project (Grant No. DP0772880, Y.B.W. andX.Z.L.), and by the US Army Research Office (Grant No.W911NF-08-1-0201, Z.C.D. and T.G.L.).

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