twoc

.Ga

in the coarse-grained alloy are born less relaxed with respect to surrounding polycrystalline aggregate than those in the ne-grained alloy.

leads to a high level of anisotropy in mechanical properties[46]. Deformation twinning strongly inuences plasticbehaviour through: (i) stress relaxation the redistributionof local micro-stresses [79]; (ii) a rapid change in texture

twinning.Acoustic emission (e.g. Refs. [1821]) and neutron dif-

fraction (e.g. Refs. [7,10,11,22]) have been used separatelyto study deformation twinning in various materials (e.g.Mg, Ti, Zr, CuGe alloys). Since both these techniquesfocus on a dierent signal coming from the deformed

* Corresponding author. Tel.: +61 2 9717 3488; fax: +61 2 9543 7179.E-mail address: ondrej.muransky@ansto.gov.au (O. Muransky).

Available online at www.sciencedirect.com

15This is believed to lead to lower reversal stresses acting on twin grains in the coarse-grained alloy upon unloading and thus to less unt-winning and thus to a smaller pseudoelastic-like hysteresis.Crown Copyright 2009 Published by Elsevier Ltd. on behalf of Acta Materialia Inc. All rights reserved.

Keywords: Magnesium; Deformation twinning; Twin nucleation; Neutron diraction; Acoustic emission

1. Introduction

Twinning is known to be an important deformationmechanism in hexagonal-close-packed (hcp) metals becauseof an insucient number of independent slip systems [1,2].Its unidirectional nature, in conjunction with the strongdeformation textures typically seen in such materials [3],

which accompanies the sudden reorientation of the crystallattice [5,7,1012]; (iii) the introduction of barriers to dislo-cation movement [1317]; and (iv) transformation of latticedislocations [13]. It is thus of great importance, particularlyfor engineering applications, to understand the collabora-tive twin-slip dislocation mechanisms by which hcp metalsundergo plasticity and to control the eects of deformationAustralian Nuclear and Technology Organization, PMB 1, Menai 2234, NSW, AustraliabSchool of Engineering, Deakin University, Pigdons Rd., Geelong 3217, Australia

cLos Alamos National Laboratories, Los Alamos, NM 87545, USAd ISIS Facility, CCLRC Rutherford Appleton Laboratory, Chilton, Didcot OX11 0QK, UK

Received 21 September 2009; received in revised form 27 October 2009; accepted 29 October 2009

Abstract

Neutron diraction and acoustic emission were used in a single in situ experiment in order to study the deformation twinning of twoZM20 Mg alloys with signicantly dierent grain sizes at room temperature. The combination of these two techniques facilitates thedistinction between twin nucleation and twin growth. It is shown that yielding and immediate post-yielding plasticity in compressionalong the extrusion direction is governed primarily by twin nucleation, whereas plasticity at higher strains is presumably governed bytwin growth and dislocation slip. It is further shown that, in the ne-grained alloy, collaborative twin nucleation in many grains dom-inates yielding, whereas twin nucleation in the coarse-grained alloy is progressive and occurs over a larger strain range. In addition, it isshown that, despite twin nucleation stresses increasing with decreasing grain size, roughly the same overall volume fraction of twins isformed in both ne and coarse parent grains. This conrms the diculty of the alternative deformation modes and suggests a negligiblesuppressive eect of grain size on twinning in the case of the strongly textured ne-grained alloy. The current results also show that twinsInvestigation of deformationcoarse-grained ZM20 Mg all

diraction and a

O. Muransky a,*, M.R. Barnett b, D

Acta Materialia 58 (2010) 15031359-6454/$36.00 Crown Copyright 2009 Published by Elsevier Ltd. on behdoi:10.1016/j.actamat.2009.10.057inning in a ne-grained andy: Combined in situ neutronoustic emission

. Carr a, S.C. Vogel c, E.C. Oliver d

www.elsevier.com/locate/actamat

17alf of Acta Materialia Inc. All rights reserved.

atematerial, they can provide independent and complemen-tary information on the deformation processes and, in par-ticular, on deformation twinning. It has been shown that,in general, acoustic emission (AE) is more sensitive to twinnucleation than to twin growth [23]. In contrast, neutrondiraction detects the overall twin population withoutregard to twin nucleation or twin growth. However, themost striking dierence is in the time resolution: neutronmeasurement provides a data point (diraction pattern)every few minutes (in the present case 7 min), whereasAE can provide data points at a frequency of 10 ms. Thepresent work combines these two techniques into a singlein situ experiment in an attempt to gain additional insightinto deformation twinning in magnesium alloys.

Because deformation twinning is strongly sensitive tograin size [16,2426], the role played by grain size in theformation and evolution of deformation twins is of partic-ular interest. In strongly textured extruded magnesiumalloys, this manifests itself in a stronger HallPetch eectin compression, where tensile {10.2} twinning is prolic[5,6,16], unlike in tension, where twinning is usually scant[2730]. In the current work, in situ neutron diraction dataare analysed together with the AE data in order to comparetwinning evolution in ne-grained (FG) and coarse-grained(CG) ZM20 Mg alloy samples. The current experiment wasparticularly designed to separate the eects of twin nucle-ation from twin growth.

2. Experimental techniques

2.1. Neutron diraction technique

Time-of-ight (TOF) in situ neutron diraction measure-ments were performed using a third-generation ENGIN-Xstressstrain diractometer at the ISIS spallation neutronsource at the Rutherford Appleton Laboratory in the UK[31,32]. The instrument is equipped with a horizontal loadframe which is mounted on a positioning table in such away that its loading axis is oriented at 45 to the incidentbeam (Fig. 1a). Two detector banks placed on either sideof the load frame, at 90 to the incident beam, allow simul-taneous collection of the full diraction patterns with scat-tering vectors (Q) aligned axially (parallel) (Q||) andradially (perpendicular) (Q\) to the loading axis. The detec-tors coverage in theQ direction is half the detectors angularspread in the horizontal plane (i.e. 8) and nearly equal tothe angular coverage out of this plane (i.e. 18) [32]. (seeENGIN-X axial and radial detector windows in theENGIN-X pole gures in Fig. 3). The axial (Q||) detectorbank thus sees grains with {hk.l}|| plane-normals orientedparallel to the loading axis, while the radial (Q\) detectorbank registers only a portion of the grains whose {hk.l}\plane-normals are perpendicular to the loading axis. Inaddition, it is necessary to emphasize that grains with thesame lattice planes registered by the axial (Q||) detector bank

1504 O. Muransky et al. / Acta Mare all oriented more-or-less equally to the load axis, whilethe grainswith the same lattice planes registered by the radial(Q\) detector bankmight be oriented dierently with respectto the loadaxis. In short, the radial {00.2}\ reection is likelyto be averaged over the axial {10.0}||, {11.0}|| and all othergrain orientations with stress axes perpendicular to the{00.2} plane-normal. Owing to this ambiguity in the radialdiraction reections, the discussion herein focuses only onthe axial diraction data (radial data are discussed in com-parison with EPSC simulations in a concurrent paper [33]).

The TOF neutron diraction patterns (Fig. 1b) were col-lected in situ by stepwise compressive and tensile loadingup to 5% of macroscopic strain, using combined stressstrain control of the load frame. Mechanical loading ofthe sample was interrupted for 7 min to allow for the neu-tron diraction pattern acquisition at a constant level ofapplied stress (i.e. stress control: rappl < ryield), or strain(i.e. strain control: rapplP ryield). The collected full dirac-tion patterns were then analysed by applying a single-peaktting routine (Rawplot) [33,34] to follow: (i) the integratedintensity (Ihk.l) change; and (ii) the relative position (dhk.l)change of hk.l diraction peaks (dhk.l spacing) as a functionof applied load.

2.1.1. Peak intensities (Ihk.l)

It has been shown previously [7,10,11] that characteristicchanges in the integrated intensities of particular hk.l dirac-tion peaks can, upon mechanical straining, be used as a sen-sitive indicator of deformation twinning. In the case of the{10.2} tensile twinning mode, which is well known to bethe most active twinning mechanism in Mg and its alloys,a sudden reorientation of a parent hk.l grain by 86.3 occursduring twinning. This leads to a rapid intensity drop in theassociated parent hk.l diraction peaks. Because of thenearly 90 reorientation of the parent grains during tensiletwinning and the adopted diraction geometry on ENGIN-X (Fig. 1a), an intensity decrease in the parent hk.l dirac-tion peaks is accompanied by a concurrent intensity increasein the twin hk.l diraction peaks in the same detector bank[7,12] (Fig. 1b). However, owing to the diraction geometryon the ENGIN-X (Fig. 1a) only twin grains originating fromthe parent {10.0}|| and {11.0}|| grains can be directly detectedin the present experiment, i.e. (i) {10.0}|| twinning? {00.2}||[7,10]; and (ii) {11.0}|| twinning? {10.3}|| [12] (see Fig. 1b).The progress of twinning in other grain families is stillobserved through the decrease in intensity in the diractionpeaks associated with that particular parent-grain family.For example, the {20.1}|| grains undergo twinning, butnewly formed twins originating from these parent grainsend up out of the detectors coverage, therefore the progressof twinning is observed only as a diraction intensitydecrease in the {20.1}|| reection (see Fig. 1b).

2.1.2. Peak position (dhk.l)

The relative change in the peak position, commonly eval-uated as the elastic lattice plane strain (ehk.l = Ddhk.l/dhk.l) isproportional to the applied load and provides invaluable

rialia 58 (2010) 15031517information on microscopic intergranular stresses [3538](i.e. stresses borne by individual hk.l grain families in the

ateO. Muransky et al. / Acta Mpolycrystalline aggregate upon loading). In the present case,the positions of diraction peaks in the diraction patternsobtained prior to loading were taken as the stress-free dhk.lspacings, so that any potential initial residual stresses stem-ming from the manufacturing process were neglected. Inaddition, owing to the strong initial texture typically seenin the extruded hcp metals [3,7,39] the stress-free dhk.l spac-ings for {00.2}|| and {10.3}|| reections which were absentin the initial axial (Q||) diraction pattern (Fig. 1b) were esti-mated from the hcp stress-free lattice parameters (Fig. 1aand c). These were determined by the Rietveld analysis[34] on the initial (prior to loading) TOF diraction pattern(the same approach was used in Ref. [12]).

2.2. Acoustic emission technique

Two piezoelectric transducers were employed continu-ously during the test to detect AE waves propagatingtowards the surface of the samples. This continued through-

Fig. 1. (a) Schematic drawing of the diraction geometry adopted on ENGINacoustic data simultaneously in one single in situ experiment. A detail showsgrain and associated neutron diraction peaks. (b) Initial and nal axial (Q||){20.1}||, {11.0}||, {10.0}|| and twin {00.2}||, {10.3}|| grain families associated witFirst gure shows a low-energy AE wave typical for dislocation slip (E = 0.083energy AE wave typical for a twin nucleation (E = 0.263 pJ). Notice that tworialia 58 (2010) 15031517 1505out each of the 7 min temporary dwells required for neu-tron diraction pattern acquisition. The AE sensors wereattached to the sample using two waveguides (steel spacers)on each side of the specimen in order to avoid their directinteraction with the neutron beam (see the schematic draw-ing in Fig. 1a). Each sensor had a sensitivity of at least25 dB Ref V lbar1, with a wide-band, at frequencyresponse in the range 1001000 kHz. The raw AE signalswere passed through a 40 dB preamplier before being pro-cessed and stored. The sample rate was 20 107 sam-ples s1 and the incoming waveforms were ltered with a1001000 kHz band-pass lter. For each test, the couplingwas kept at a consistent quality by actively pinging the sen-sors with a step voltage and recording the transmitted wave.The dierence in amplitude response between each experi-ment was maintained to within a 2.5 dB range with respectto the initial test signal.

In order to detect continuous type emissions, theacoustic activity was sampled continuously by way of

-X and the principle of collecting the in situ neutron diraction and in situa parent {10.0}|| grain undergoing twinning together with its twin {00.2}||neutron diraction patterns showing dramatic intensity changes in parenth the propagation of the tensile twinning it the parent-grain families. (c)pJ) together with its Fourier transformation; second gure shows a high-waves can have the same median frequency fm but very dierent energy.

the Root-Mean-Square (RMS) signal voltage. In the cur-rent experiment, the RMS was calculated and recordedrepeatedly over 10 ms periods. The advantage of usingthe RMS transducer voltage is that it enables continuousmeasurement of a parameter which can be used for com-parative experiments [40]. The threshold was selectedsuch that external background noise (such as that fromthe hydraulic load frame) could be excluded from themeasurements. Several measurements were taken ofbackground noise prior to testing under load and while

3.2. Texture

Fig. 3 shows the initial (as-extruded) texture of bothalloys carried out on the HIPPO (high-pressure-preferredorientation) TOF powder diractometer at LANSCE inthe USA [43,44]. It is seen that both samples have a strongaxisymmetrical texture, with the majority of crystallites ori-ented with their basal {00.2} plane-normals perpendicularto the extrusion direction. There are quantitative texturedierences between the alloys. For instance, the FG alloy

h th

0 m

0 m

1506 O. Muransky et al. / Acta Materialia 58 (2010) 15031517the neutron ux was present to ensure that acousticactivity was detected only in the test specimens. Thebackground noise in the continuous RMS measurementswas 0.002 V.

3. Experimental material

In the present study, commercially available alloy,ZM20 (Mg2Zn0.5Mn), was extruded by two slightlydierent thermo-mechanical routes in order to produceFG and CG samples. The thermo-mechanical (extrusion)parameters are shown in Table 1. In short, the dierencein the grain size is caused by the dierent forming (extru-sion) input-heat, which leads to dierent recrystallizationbehaviour of the material during the extrusion process[19,41,42] and thus to dierent average grain sizes.

3.1. Microstructure

The micrographs on the section parallel to the extrusiondirection in Fig. 2 show that the CG alloy displays a fullyrecrystallized microstructure, whereas the FG alloy is onlypartially recrystallized. Using the linear intercept method(1000 grains), the grain size of the alloys studied wasdetermined as 114.3 48.7 lm (mean standard devia-tion) and 17.1 10.3 lm for the CG alloy and FG alloy,respectively. The grain size statistics shown in Fig. 2 (seethe bottom histograms depicting d3 f as the y-axis, andd as the x-axis, where d is the mean equivalent circle diam-eter in each histogram bin, and f the frequency) point to thebimodal character of the microstructure of the FG alloy.This is due to large grains extending in the extrusion direc-tion. Such recovered grains are often seen in commerciallyextruded material which did not undergo fullrecrystallization.

Table 1Extrusion thermo-mechanical parameters for both tested alloys together witcompression along the extrusion direction.

Alloy Extrusion details

ZM20

FG (1) as-cast; (2) hom. 450 C/8 h; (3) ext.450 C, 1 mm s1, 4, 3(4) ext.280 C, 1 mm s1, 4, 15 mm, AC

CG (1) as-cast; (2) hom. 450 C/8 h; (3) ext.450 C, 1 mm s1, 4, 3

(4) ext.480 C, 1 mm s1, 4, 15 mm, ACexhibits more grains with a {10.0} plane-normal orientedin the extrusion direction, whereas the CG alloy has moregrains with a {11.0} plane-normal oriented in the extrusiondirection. Nevertheless, the overall character of the initial(ex-extruded) texture is the same in both tested alloys.(The initial texture is discussed in more detail, togetherwith the in situ neutron diraction results, in Section 4.2.)

3.3. Test specimens

Cylindrical test specimens for the in situ compression( = 10 mm, l0 = 20 mm) and tensile ( = 10 mm,l0 = 20 mm) testing were machined from the as-extrudedrods (/ = 15 mm) parallel to the extrusion direction.Therefore, the majority of the crystallites (grains) are ori-ented with their {00.2}\ plane-normals perpendicular tothe loading (extrusion) axis, diracting into the radial(Q\) detector bank, and with their {10.0}||, {11.0}|| or{20.1}|| plane-normals parallel the loading (extrusion) axis,diracting into the axial (Q||) detector bank (see Fig. 3).

4. Results and discussion

4.1. Macroscopic stressstrain response

The stressstrain curves obtained in both compressionand tension along the extrusion axis during the stepwisein situ loading are shown together with the AE data inFig. 4. The yield stresses are summarized in Table 1. Asexpected, owing to the strong initial texture and unidirec-tional nature of the deformation twinning, which is pro-fusely active only in compression, both tested alloysexhibit tensioncompression asymmetry. Furthermore,the general nature of the macroscopic work-hardening issignicantly dierent in tension and compression. This

eir as-extruded grain sizes and the yield stresses determined in tension and

Grain size (lm) Yield stress (MPa)

Mean SD Rp0.2 Com. Rp0.2 Ten.

m, AC; 17.1 10.3 84 139

m, AC; 114.3 48.7 42 135

ateO. Muransky et al. / Acta Mhas been previously attributed to the action of deformationtwinning [5,6,12,16]. It is further seen in Fig. 4 that theyield stress in compression increases dramatically withdecreasing grain size, whereas there is only a subtle increasein the yield stress in tension with decreasing grain size. Theobserved lack of the HallPetch eect in tension is in con-tradiction with the previous results in Refs. [4548], whereyield stress dependence on the grain size with respect to theHallPetch relation was clearly shown, although, in accor-dance with work by Kurzydlowski and Bucki [49], the over-lapping of grain sizes owing to the bimodal grain sizedistribution (see Fig. 2) can lead to lower yield strength,as would be expected from the HallPetch relation. There-fore, the authors believe that the observed anomaly is dueto the bimodal microstructure of the FG alloy (Fig. 2) anda possible eect of the slight dierence in the initial texturebetween the alloys [45,46]. However, this emphasizes a fargreater eect of the grain size on the ow stress in compres-sion (i.e. dierence in slip-controlled yielding and twin-con-trolled yielding). It becomes clear from the current resultsthat the HallPetch slope must be signicantly greater incompression along the extrusion direction, where macro-

Fig. 2. Initial microstructure of both tested ZM20 Mg alloys together with graiof the as-extruded microstructure: (a) FG alloy; (b) CG alloy. The statistical hThe disparity in the grain size measurements in FG histograms for the rightmgrains.rialia 58 (2010) 15031517 1507scopic yielding is known to be governed by profuse defor-mation twinning. This conrms previous observations[26,50,51] that the grain size impacts on twinning somehowmore that it does on slip. In addition, one can see that theFG alloy displays a small Luders-like (zero work-hard-ening) plateau in compression, which is not seen in anyof the other stressstrain curves.

4.2. Neutron diraction results

The evolution of hk.l diraction intensities (Ihk.l) and thedevelopment of hk.l elastic lattice plane strains (ehk.l) in thegrains oriented with their {hk.l} plane-normals in theextrusion/loading direction are shown: (i) as a function ofapplied compressive engineering stress in Fig. 5; and (ii)as a function of elapsed time in Fig. 6. (Owing to the lim-ited amount of twinning in tension along the extrusiondirection [27,28,30], the tensile data have been excludedfrom this discussion.) For clarity, Figs. 5 and 6 depict onlythe diraction reections associated with: (i) the parent-grain families, i.e. {10.0}||, {20.1}||, {11.0}|| the grainswhich are the most favourably oriented for the tensile

n size statistics which shows the heterogeneity (uniformity) in the grain sizeistograms are normalized in order to see the grain size distribution clearly.ost bin stems from longitudinal measurements of the fully recrystallized

ate1508 O. Muransky et al. / Acta Mtwinning in compression along the {10.0}, {20.1}, {11.0}plane-normals; and (ii) newly formed twin grain families,i.e. {00.2}||, {10.3}||the grains (twins) formed in thedeformation process as a consequence of twinning in theparent grains. Notice that, owing to the nature of the extru-sion texture (Fig. 3), the {10.0}||, {20.1}||, {11.0}|| grain fam-ilies are also the most populated, therefore the plasticity inthem dominates that of the whole polycrystalline aggregate[5,28,12].

4.2.1. Twinning and overall twin volume fraction

First, the discussion focuses on the evolution of the inte-gral intensities of the parent-grain and twin-grain dirac-tion reections. In order to follow one-to-one ippingbetween the parent and twin grains in the instances of com-pressive straining, the axial (Q||) peak intensities are pre-sented as reduced by the 4th power of the dhk.l spacingand the 2nd power of the Fhk.l the structure factor [52].(Note that there is no change in the multiplicity betweenthe parent and twin diraction peaks considering one-to-one parent-to-twin ipping [53], and the eect of hk.l-dependent attenuation [32] is neglected.) These reducedTOF intensities (rIhk.l) are proportional to the volumeoccupied by particular hk.l grains, oriented such that theyfull the diraction condition in the axial (Q||) detectorbank [52].

Figs. 5c, d and 6c, d clearly identify the quantitative dif-ferences in the initial (as-extruded) texture between the FG

Fig. 3. Initial (as-extruded) texture of (a) the FG alloy and (b) CG alloy measudetector banks coverage.rialia 58 (2010) 15031517and CG alloys. This is consistent with the texture measure-ment. However, the initial {10.0}|| and {11.0}|| diractionintensities prior to loading in the CG are more-or-less iden-tical in Figs. 5 and 6d, despite a noticeable dierence in thepole densities seen in the {10.0} and {11.0} pole gures inthe extrusion direction in Fig. 3b (centre of the pole g-ures). This experimental inconsistency can be attributedto the diraction averaging over a large detector area onENGIN-X (see 16 36 ENGIN-X detector windowsdepicted in the pole gures in Fig. 3).

For a simpler comparison of the initial texture betweenthe tested samples, the diraction peak intensities (rIhk.l) ofall the parent grains studied are summed in Fig. 7 (RrIpar-ent = rI10.0 + rI20.1 + rI11.0). Here, it becomes clear that theinitial volume fraction of the parent grain families isroughly equal for both FG and CG alloys. As there isroughly the same volume fraction of parent grains, thisenables a direct comparison of twinning activity betweenthe alloys. One must, however, recognize that dierent par-ent grains (when twinned) will contribute by a slightly dif-ferent fraction to the overall macroscopic strain. Thiscontribution can be estimated by: etwin = Rc0mhk.lfhk.l,where etwin is the twinning strain, c0 is the tensile twinshear strain (=0.13), mhk.l is the Schmid factor, and fhk.lis the volume fraction of the twinned parent hk.l grains[54].

It can be seen in Figs. 5 and 6 that the diraction inten-sities of the parent-grain {10.0}||, {20.1}||, {11.0}|| reec-

red on HIPPO. Pole gures show the axial (Q||) and radial (Q\) ENGIN-X

ateO. Muransky et al. / Acta Mtions in both samples start to decrease in the vicinity ofyielding (see P1, which marks the rst neutron diractionmeasurement >0.2% engineering strain) implying the oper-ation of tensile twinning in them. The onset of twinningis, however, seen somewhat better in the appearance of thetwin {00.2}||, {10.3}|| diraction reections, which areabsent prior to loading and are thus a very sensitive indica-tor of twin initiation. In general, it is clear that the decreasein diraction intensity of the parent-grain reections isaccompanied by a concurrent intensity increase in dirac-tion reections associated with the new twin-grain orienta-tions. Indeed, there is a one-to-one intensity exchangebetween the parent-grain {10.0}|| and twin {00.2}|| dirac-tion reections. This reects the reorientation of the{10.0}|| grains into the new twin {00.2}|| orientation duringtwinning [7,10,12,22]. There is not, however, a one-to-oneintensity exchange between the parent {11.0}|| and twin{10.3}|| reections, because some of the twinned {11.0}||grains end up out of the axial (Q||) detector coverage [12].

Despite the dierent Schmid factors (mhk.l) of the parentgrains (m10.0 = 0.5, m20.1 = 0.45, m11.0 = 0.39), twinningseems to begin in all parent grain orientations more-or-lesssimultaneously (see Figs. 5 and 6). This suggests a lack ofalternative deformation mechanisms. It is, however, clear

Fig. 4. Macroscopic stressstrain curves of FG (left-hand side) and CG (right-h(stepwise) compression (a and b) and tensile (c and d) loading along the extruover each data acquisition interval of 7 min. The in situ stressstrain curves aThe same AE data are plotted in the bar chart in inserts showing the AE actirialia 58 (2010) 15031517 1509from the relative decrease in the parent diraction intensi-ties in Figs. 5c, d and 6c, d that twinning then proceeds atdierent rates in dierently oriented parent grains. In fact,it can be seen that the twinning proceeds in each parent-grain family in accordance with its Schmid factor, i.e. thefastest rate is observed in {10.0}|| grains, and the slowestin {11.0}|| grains. In addition, it is shown in Fig. 5c and dthat the fractional decrease in the parent-grain {11.0}||,{20.1}||, {10.1}|| diraction intensities is 69%, 59% and58%, respectively, in the FG sample, and 69%, 64% and61%, respectively, in the CG sample. This also agrees wellwith Schmid factors, and it shows that a slightly higherproportion of the coarse parent grains twinned duringstraining.

For a more convenient comparison of the twinningactivity between the FG and CG samples, one can observein Fig. 7 that, on average, the parent-grain families in bothalloys twin at more-or-less the same rate. However, it canbe seen that at strains

ate1510 O. Muransky et al. / Acta Ma compressive strain of 5%, 65% of the volume fractionof the parent orientations was reoriented by tensile twin-ning in both samples. This might seem to be in contrastwith the widely accepted view that small grains suppresstwinning, but it needs to be understood with respect tothe strong initial texture which eectively suppresses alter-native deformation modes (the second-order hc + ai pyra-midal slip [27,28]). It appears, somewhat similarly, fromthe work by Jain et al. [48] that the twin volume fractionis not strongly aected by the grain size at room tempera-ture for sizes greater than 30 lm (see Fig. 8 in Ref. [48]).

Finally, it can be seen in Fig. 7c that, upon unloading,the (sum of) parent-grain diraction intensities (RrIparent)increases, whereas the (sum of) twin-grain diraction inten-

Fig. 5. (a and b) Macroscopic strainstress curves recorded during the in situ (hand side) and the CG alloy (right-hand side) together with the AE data collecteintegral intensities in the diraction reections associated with the parent-grain10 ms on the linear scale. (e and f) Development of the elastic lattice strains in alevery 10 ms on the linear scale. The same data are shown as a function of elarialia 58 (2010) 15031517sities concurrently decrease. The same can be seen from theindividual diraction reections in Fig. 6c and d. (Notethat the unloading points as a function of engineering stresswere omitted for greater clarity in the drawn graphs, i.e.Fig. 5c and d, Fig. 7d). This decrease in the volume fractionof twins upon unloading was previously attributed to spon-taneous untwinning. This is thought to be driven by thereversal stresses acting on the freshly twinned grainsupon unloading, thus leading to a pseudoelastic-like eect[12]. It is seen in Fig. 7 that higher volume fraction of twinsundergo untwinning in the FG alloy (see the insert inFig. 7c). This is supported by a larger macroscopicspring-back of the FG alloy upon unloading (see Fig. 7aand b). This nding is in agreement with previous studies

stepwise) compression along the extrusion direction of the FG alloys (left-d every 10 ms shown on the linear scale. (c and d) Evolution of the reducedfamilies and twin grain families together with the AE data collected everyl studied parent and twin grain families together with the AE data collectedpsed time in Fig. 6.

by Caceres et al. [58] and Mann et al. [59], who showed thatthe smaller the grain size, the larger the pseudoelastic-likeeect. Further clarication and explanation of this phe-nomenon are provided in the following section.

4.2.2. Twinning and internal stressesAnother point of interest in the current neutron dirac-

tion experiments is the development of the elastic latticestrains (ehk.l) reecting the internal stresses in the parentand the newly formed twin grains. When analysing the dif-fraction elastic lattice strains during a twinning event, oneneeds to keep in mind that, due to the reection analysis

method, a partially twinned {10.0}|| parent grain is detectedas (i) a twin in the twin {00.2}|| reection and (ii) simulta-neously the remaining part of the same parent grain isdetected in the parent {10.0}|| reection (see Fig. 1). Theparent-grain reections are thus averaged over the twinned(relaxed) parent grains and simultaneously over the as yetnot twinned (stressed) parent grains. However, the twinreections are averaged only over the twinned parentgrains. The authors are mindful that, as the deformationproceeds, diraction averaging in the twin-grain reectionsgoes through (i) freshly twinned parent grains, and (ii)previously twinned parent grains, which are known to

ringAEoutpdevasso

O. Muransky et al. / Acta Materialia 58 (2010) 15031517 1511Fig. 6. (a and b) Evolution of the engineering (applied) stress and strain dualloys (left-hand side) and the CG alloy (right-hand side) together with thethe extensometer (strain) data is only the reading from the second Instronreading from the second Instron output was collected by the AE controllingd) Evolution of the reduced integral intensities in the diraction reections

the AE data collected every 10 ms with the linear scale. (e and f) Developmentogether with the AE summed up over a 10 s time increment and shown withthe in situ (stepwise) compression along the extrusion direction of the FGdata collected every 10 ms shown with the log scale. (Note that the noise inut, and it is not the controlling strain input which is shown in Fig. 4. Theice in order to synchronize the AE data with the stressstrain data.) (c andciated with the parent-grain families and twin grain families together with

t of the elastic lattice strains in all studied parent and twin grain familiesthe linear scale.

accommodate elastic strain at a higher rate than the sur-rounding grains [7,9,12]. This is of great importancebecause of the redistribution of local micro-stresses associ-ated with twinning [8,9,12]. It is seen in Figs. 5e, f and 6e, fthat newly formed {00.2}||, {10.3}|| twins are nucleated in arelaxed state with respect to their parent grains and the sur-rounding polycrystalline aggregate (lower internal strainsseen at the onset of twinning in twin grain families).

This nding agrees with previous studies on extrudedAZ31 Mg alloy [7,1012]. Clausen et al. [9] recently mea-sured a tensile strain of roughly +2000 le within newlynucleated (formed) twins at the onset of twinning in anextruded AZ31 alloy (see Fig. 7a in Ref. [9]). Such a rever-sal of stresses in twinned grains was not seen in the exper-iments performed in the present work (this includes anadditional eight other unreported measurements carriedout on four dierent alloys, including AZ31). This dier-ence could be ascribed either to an unlucky strain incre-ment in our in situ experiments, where newly formed twinsare detected for the rst time only after they already accu-mulated a small amount of the compressive strain. This

would mean that the moment was missed when the neutronaveraging goes solely through freshly nucleated twins.Alternatively, the observed dierence could be attributedto micro-slip associated with twin nucleation. This immedi-ately relaxes high local stresses caused by the mistbetween the twined grains and the surrounding aggregate.Note that if a {00.2}|| twin is generated within a {10.0}||parent grain without any additional plasticity (micro-slip),the {10.0}|| parent and its {00.2}|| twin grain will be in ten-sion along the loading axis with respect to its surrounding.This is ultimately due to a sudden geometrical shorteningof the twinned parent grain (Fig. 1a) and the constraintsimposed by the surrounding polycrystalline aggregate.

From comparison of the lattice strains at the onset oftwinning in the {10.0}|| parent grains in Figs. 5, 6e and fit is clear that twinning begins in the ne {10.0}|| parentgrains at a lattice strain of 1610 le and in the coarse par-ent grains with the same orientation at a signicantly lowerlattice strain of 670 le. Taking into account the elasticconstant of 45 GPa, one can estimate the internal stressin the ne and coarse {10.0}|| parent grains at the onset

(sts of

1512 O. Muransky et al. / Acta Materialia 58 (2010) 15031517Fig. 7. (a and b) Macroscopic strainstress curves recorded during the in situCG alloy. (c and d) The evolution of the sum of reduced integral intensitie

twin-grain reections (RrItwin = rI00.2 + rI10.3). Insert shows the intensity decreaparent-grain reection upon unloading.epwise) compression along the extrusion direction of the FG alloys and thethe studied parent-grain reections (RrIparent = rI10.0 + rI20.1 + rI11.0) and

se in the twin-grain reections and the concurrent intensity increase in the

of twinning: 75 MPa in ne {10.0}|| grains and 30 MPain coarse grains. Note that this is aected by (i) the initialtexture, (ii) the heterogeneity in the grain size, and (iii)the grain-to-grain interactions. It is, nevertheless, clear thatthe twin nucleation stress is considerably higher in the FGalloy. Since the suppressing eect of grain size on twinningis due to the increasing stress required for twin nucleation(i.e. twin nucleation stress) [2426], one would expect sig-nicantly less twin activity in the FG alloy. From Fig. 7cand d it is, however, clear that roughly the same overallvolume fraction of twins is formed in both tested alloys.This implies that, even though the twin nucleation stressis increasing with decreasing grain size, it is still lower thanthe stresses required for competitive slip (e.g. the second-order hc + ai pyramidal slip [27,28]). Therefore, there iseectively no twin suppression in the FG alloy, just a

from the internal redistribution of stresses upon twinning[8,9,12] by decreasing the mist between the twinned grainsand the surrounding polycrystalline aggregate (see Fig. 9Cin Ref. [12]). Larger stress relaxation of twinning grains inthe FG alloy suggests larger mists between twinned grainsand surrounding aggregate, and thus less micro-slip associ-ated with twins formation. This is most likely due to thesmall size of twins and thus small local stresses. Addition-ally, larger stress relaxation of freshly formed twin grainsin the FG alloy means that upon unloading they areexposed to higher reversal stresses than twins in the CGalloy. These reversal stresses acting on twin grains uponunloading provides driving force for spontaneous untwin-ning and thus give rise the pseudoelastic-like hysteresiseect [12]. Hence, there is a higher driving force for untwin-ning [12] upon unloading and thus a stronger pseudoelas-

O. Muransky et al. / Acta Materialia 58 (2010) 15031517 1513higher ow stress (see Fig. 7d). (Of course, if the twinnucleation stress becomes higher than the stress for com-petitive slip, twinning will be suppressed.)

It is further interesting to notice the dierence in thestress relaxation upon twinning in the ne and coarse{10.0}|| parent grains. In Figs. 5e, f and 6e, f, it can beclearly seen that the dierence in the elastic lattice strains(De) between the parent {10.0}|| and their twin {00.2}||grains at the onset of twinning (i.e. when the rst twinsare nucleated) is 1000 le in the FG alloy and only450 le in the CG alloy. This suggests a smaller stressrelaxation upon twinning of coarse parent grains. In short,it means that twins in coarse grains are formed somehowless relaxed with respect to the surrounding polycrystallineaggregate. (This seems to be conrmed by the results ofAydiner et al. [55], who studied the AZ31 Mg alloy withgrain size 100 lm [56].) The lower stress relaxation oftwins in the CG alloy can be rationalized by easiermicro-slip in the coarse fully recrystallized grains. Addi-tionally, larger twins are more easily stabilized by micro-slip, owing to the higher local stresses in their immediatevicinity [57]. This micro-slip associated with twin nucle-ation can eectively relax extensive local stresses resultingFig. 8. Optical micrograph of the CG alloy after compression to a strain oshadows represent grains that have totally reoriented, and the shadows revetic-like eect in FG alloys (see insert in Fig. 7c andprevious experimental results in Ref. [58,59]). Moreover,small twins are thermodynamically less stable [58,59],and they are less stabilized by micro-slip [51]. This alsomakes them more vulnerable for untwinning uponunloading.

Finally, it can be seen in Fig. 6e and f that, at higherstresses the applied load is gradually transferred towardstwin {00.2}|| grains which (on average) accommodate elas-tic strain at a higher rate than other grain families. This isdue to the hard plastic orientation of the newly formed{00.2}|| twins, which thus behave as fast-growing hardinclusions [12] and gradually bear a higher proportion ofthe applied load. The same, however, does not apply tothe {10.3}|| twins because of the favourable orientationfor basal slip of hai dislocations [12]. This behaviour seemsto be fairly consistent between the present samples and thepreviously reported results on AZ31 alloy in Refs. [7,8,12].

4.3. Acoustic emission results

The continuously measured AE data, sampled every10 ms, are presented as RMS voltage and are shownf 5%. Twins can be seen in all grains. The grains containing faint twinal locations of twin impingement.

atetogether with the macroscopic stressstrain curves as afunction of engineering strain in Fig. 4, and with the neu-tron diraction data as a function of engineering stress inFig. 5, and elapsed time in Fig. 6. Notice that, in orderto facilitate the analysis, the AE data are shown (i) witha logarithmic scale in Fig. 6a and b, (ii) with a linear scalein Fig. 6c and d, and (iii) the data are then summed over10 s time increments and shown with the linear scale inFig. 6e and f.

4.3.1. Twinning vs dislocation slip and AE activity

Let us rst compare AE data recorded during compres-sion and tension loading along the extrusion direction. It isclear from Figs. 4 and 5 that AE activity rises signicantlyand reaches a local maximum in compression as well as intension in the vicinity of yielding (i.e. at the onset of plas-ticity). This has been previously attributed to (i) a suddenincrease in dislocation movement, (ii) massive dislocationmultiplication, and (iii) deformation twinning (if present)[39,60]. The most striking observation is, however, themagnitude (power) of the AE signal, which is 10-fold incompression compared with tension. It stands to reasonthat this is due to the strong eect of the deformation twin-ning on the AE signature [18,20,23,60]. The following dis-cusses in detail only the compression along the extrusiondirection which is known to be governed predominantlyby deformation twinning (see the discussion of the neutrondiraction results in Section 4.2).

4.3.2. Twinning and AE activityIn general, it can be seen from the present AE data

(Figs. 46) that the overall trend in AE activity is similarin both alloys, though it does appear that there are moreevents occurring in the elastic region in the CG alloy. Thismight point to earlier/easier activation of dislocation slip,which leads to AE activity well before yielding. Withincreasing applied load, the AE activity increases gradually(more dramatically in the FG alloy, see inserts in Fig. 4aand b) reaching a peak at the onset of macroscopic plastic-ity (this is best seen in Fig. 6e and f). Fig. 5 shows that thepeak of the AE activity coincides nicely with the onset ofdeformation twinning, manifested by the decrease in thediraction intensity of the parent-grain reections andthe concurrent increase in diraction intensity of reectionsassociated by newly formed twin orientations.

After an initial rise in the AE activity close to the yield,the AE signal subsequently declines rapidly at strains>0.5% (0.8% in the CG alloy), before reaching a more-or-less continuous low-level at strains >1.52% (seeFig. 4a and b). This post-yielding decrease in the AE activ-ity was previously attributed to (i) the increasing number ofdislocations and thus decreasing ight path and freelength for dislocation movement [61] and (ii) exhaustionof twinning activity. The latter is apparently in conict withthe current results, because as can be seen from the neutron

1514 O. Muransky et al. / Acta Mdiraction results, the twinning continues to operate at thesame rate at higher strains (Figs. 57). The observeddecline in the AE activity cannot be thus simply attributedto an exhaustion of twinning activity. The present neutrondiraction acoustic emission results, thus clearly show, inagreement with Refs. [20,23], that the AE measurement ismuch more sensitive to twin nucleation than to twingrowth.

4.3.3. Macroscopic stress relaxation and AE activity

Finally, a brief discussion is given of the AE activityduring the temporary stopovers (7 min) during whichthe applied strain (rapplP ryield) was maintained at a con-stant level in order to allow the acquisition of neutron dif-fraction data. It can be seen in Fig. 6 that in both testedsamples the AE activity drops rapidly during these tempo-rary stopovers, which are associated with macroscopicstress relaxation in the plastic region. One can notice inFig. 6 and b (notice the logarithmic scale) that the AEactivity decreases during the stress relaxation as a functionof time and then reappears during subsequent loading.More importantly it is seen that low-level AE activityremains during the stress relaxations at strains

atepart by the fact that twins can be autocatalytic in their for-mation [24]. That is, twins in one grain can lead to thenucleation of twins in another via an increase in the localstress at the grain boundary due to the twin impingingon the grain boundary. Such a phenomenon is apparentin Fig. 8, where twins are frequently seen to be connectedto other twins in neighbouring grains. (Note that all grainscontain either distinct twins or faint markings in the shapeof twins which are either unetched twins or low angleboundaries marking where twins have impinged on eachother.) One might expect the autocatalytic nature of twinnucleation to be more pronounced in the FG alloy becauseof a higher grain boundary area per unit volume. Thisseems to be conrmed in the present results by the appar-ently more massive character of the onset of twinning inthe FG alloy; see the discussion in Section 6. Fig. 8 alsoprovides some evidence that twin growth can be coopera-tive in neighbouring grains. In such cases, the length ofconnection along the grain boundary between twins inneighbouring grains is considerable. This could occur bythe transmission of twinning dislocations through favour-ably oriented grain boundaries.

The drop in twin activity with strain following yielding,as detected by AE (Figs. 5 and 6), is not likely to be simplydue to the replacement of volume by twinned material. It ismore likely that the twin nucleation frequency dropsbecause of the consumption of easily activated nucleationsites. Given that the stresses required to initiate twinningcan be expected to vary over the grains in a sample, thegrains in which twinning is easiest will twin rst. Withincreasing twin nucleation, the stresses needed to activatesubsequent sites becomes progressively greater. This eectcan be expected to be accentuated by the elastic relaxationthat accompanies twin formation (see Section 4.2). In otherwords, regions of relaxed material exist around the initia-tion point of a twin within which subsequent twin nucle-ation is less likely because of the relaxed stresses [8,9]. Asthe number of twins increases, this gives way to a dampen-ing of twin formation owing the relaxation of local stressesand the consumption of easy nucleation sites.

6. Yielding by twinning and the grain size eect

The direct quantitative comparison of AE activitybetween the FG and CG alloys is complicated because ofthe potential eects of grain size on the AE signals smal-ler twins in the FG alloy are likely to produce weaker AEsignals. Nevertheless, it is clear from Fig. 6e and f thatthere is a noticeably stronger AE activity at the onset ofplasticity (twinning) in the FG alloy, despite the possibledepleting eect of the small grains on the AE signal [19].Therefore, if one considers (i) the signicantly dierentgrain size between the FG and CG alloys, and (ii) the sim-ilar overall volume fraction of twins measured by the neu-tron diraction in both samples, it suggests that a higher

O. Muransky et al. / Acta Mnumber of twins nucleate at the onset of plasticity in theFG alloy. This is simply due to the constraints which thegrain size places on the twin size [26]. Thus, for the FGalloy, there will be a greater area of twin interface capableof growth. It is generally thought that twin growth occursunder lower stresses than those required for nucleation[54,64]. It is therefore appropriate to state that the relativeamounts of twin interface at yielding may also be animportant factor for understanding the stress and spreadof twinning following yielding.

Comparing the width of the AE peak in terms ofstrain range (see inserts in Fig. 4a and b) suggests that highAE activity remains slightly longer (at higher strains) inthe CG sample (up to 0.8%). It points to the progressivetwin nucleation in the CG alloy at the onset of plasticity.In contrast, twin nucleation in the FG alloy seems to occurcollaboratively in many parent grains at once, see the sharpAE peak in the vicinity of yield point in the insert inFig. 4a. Notice that the AE activity in the inserts inFig. 4 somewhat resembles the grain counting statistic dia-gram in Fig. 2. This points to a strong grain size eect onthe twin nucleation stress. The profuse twin nucleation inthe FG alloy at yield leads to a sudden decrease in thework-hardening, and subsequent (post-yielding) twinningoccur without additional increase in the applied stress(see the continuing twinning during the Luders-like pla-teaux in the FG alloy in Fig. 5c). The situation is slightlydierent in the CG alloy, where a wide grain size distribu-tion (48.7) and probably earlier/easier onset of disloca-tion slip prevents collaborative twinning in many grainsat the onset of plasticity. In addition, micro-slip associatedwith the twin nucleation in fully recrystallized coarse grainsseems to eectively eliminate local reversal stresses gener-ated during a twinning event (see Section 4.2) and thus pre-vents the softening eects of deformation twinning (localstress relaxation).

Turning now briey to the stronger eect of grain sizeon yield stress in compression compared with tension seenin the stressstrain curves, this has been previously attrib-uted to the reportedly higher HallPetch slope for yieldingwhen it is governed by deformation twinning [26,50,51]. Inthe light of the present results, the grain size seems to aectyielding particularly via (i) the readiness of micro-plasticslip activated during twin nucleation, and (ii) independentdislocation slip, which assists yielding, even though it is pri-marily governed by deformation twinning. The currentresults seem to oer further evidence of the twofold eectof grain size on deformation twinning [26], i.e. the grainsize (i) inuences the twin nucleation through micro-plasticslip and its accommodation near grain boundaries [24,25],and (ii) it is restricting the progress of twinning into neigh-bouring grains; however, this seems to be opposed by theautocatalytic character of twin nucleation and even twingrowth (see Section 5).

7. Summary

rialia 58 (2010) 15031517 1515The present work combined neutron diraction and AEin one single in situ experiment in order to investigate

1516 O. Muransky et al. / Acta Materialia 58 (2010) 15031517deformation twinning in a ne-grained (FG) and coarse-grained (CG) ZM20 Mg alloy. The results show the com-plementarity of the techniques used, where it becomesclear that the AE technique is more sensitive to twinnucleation than it is to twin growth. However, the neutrondiraction technique is sensitive only to an overall twinpopulation, without distinguishing between twin nucle-ation and growth. Hence, the combination of these tech-niques shed more light on the process of the twinnucleation and twin growth as well as on the rather com-plicated nature of yielding governed by deformationtwinning.

Some of the specic conclusions are as follows.

1. It is shown that the most of the twin nucleation is hap-pening at the onset of plasticity at strains

[15] Salem AA, Kalidindi SR, Doherty RD. Acta Mater 2003;51:422537.[16] Barnett MR, Keshavarz Z, Beer AG, Atwell D. Acta Mater

2004;52:5093103.[17] Salem AA, Kalidindi SR, Semiatin SL. Acta Mater 2005;53:3495502.[18] Vinogradov A, Merson DL, Patlan V, Hashimoto S. Mater Sci Eng A

2003;34:15773.[19] Bohlen J, Chmelik F, Dobron P, Kaiser F, Letzig D, Lukac P, et al. J

Alloy Compd 2004;378:2149.[20] Lou XY, Li M, Boger RK, Agnew SR, Wagoner RH. Int J Plasticity

2007;23:4486.[21] Lu Y, Gharghouri MA, Taheri F. Nondestruct Test Eval

2008;23:14161.[22] Oliver EC, Daymond MR, Withers PJ. Mater Sci Forum 2005;495

497:15538.[23] Toronchuk JP. Mater Eval 1977(October):513.[24] Ecob N, Ralph B. J Mater Sci 1983;18:241929.[25] El-Danaf E, Kalidindi SR, Doherty RD. Metall Mater Trans A

1999;30:122333.[26] Barnett MR. Scripta Mater 2008;59:6968.[27] Agnew SR, Tome CN, Brown DW, Holden TM, Vogel SC. Scripta

Mater 2003;48:10038.[28] Agnew SR, Brown DW, Tome CN. Acta Mater 2006;54:484152.[29] Barnett MR. Mater Sci Eng A 2007;464:816.[30] Muransky O, Carr DG, Barnett MR, Oliver EC, Sittner P. Mater Sci

[39] Tome CN, Lebensohn RA, Kocks UF. Acta Metall Mater1991;39:266780.

[40] Wadley HNG, Scruby CB, Speake JH. International Metals Reviews,No. 2, Review 249; 1980. p. 4164.

[41] Swiostek J, Bohlen J, Letzig D, Kainer KU. In: Proceedings of the 6thinternational conference and exhibition on magnesium alloys andtheir applications; 2003. p. 278.

[42] Bohlen J, Yi SB, Swiostek J, Letzig D, Brokmeier HG, Kainer KU.Scripta Mater 2005;2:25964.

[43] Wenk H-R, Lutterotti L, Vogel S. Nucl InstrumMethods Phys Res A2003;515:57588.

[44] Matthies S, Pehl J, Wenk H-R, Lutterotti L, Vogel SC. J ApplCrystallogr 2005;38:46275.

[45] Wilson DV, Chapman JA. Philos Mag 1963;8:154351.[46] Rao GS, Prasad YVRK. Bull Mater Sci 1983;5(5):45964.[47] Wang YN, Huang JC. Mater Trans 2007;48:1848.[48] Jain A, Duygulu O, Brown DW, Tome CN, Agnew SR. Mater Sci

Eng A 2008;486:54555.[49] Kurzydlowski KJ, Bucki JJ. Acta Metall Mater 1993;41:31416.[50] Hull D. Acta Metall 1961;9:191204.[51] Meyers MA, Vohringer O, Lubarda VA. Acta Mater

2001;49:402539.[52] Windsor CR. Pulsed neutron scattering. London: Taylor & Francis;

O. Muransky et al. / Acta Materialia 58 (2010) 15031517 1517Eng A 2008;496:1424.[31] Dann JA, Daymond MR, James JA, Santisteban JR, Edwards L.

Physica B: Condens Matter 2004;350:e5114.[32] Santisteban JR, Daymond MR, Edwards L, James JA. J Appl

Crystallogr 2006;39:81225.[33] Muransky O, Barnett MR, Luzin V, Vogel SC. doi:10.1016/

j.msea.2009.10.018.[34] Larson A, Von Dreele RB. General Structure Analysis System

(GSAS). Los Alamos National Laboratory Report LAUR 86748;1994.

[35] Clausen B, Lorentzen T, Bourke MAM, Daymond MR. Mater SciEng A 1999;259:1724.

[36] Pang JWL, Holden TM, Wright JS, Mason TE. Acta Mater2000;48:113140.

[37] Daymond MR, Tome CN, Bourke MAM. Acta Mater2000;48:55364.

[38] Daymond MR. Acta Crystallogr 2005;61:c32.1981.[53] Oliver EC, Daymond MR, Withers PJ. Mater Sci Forum 2005;490

491:25762.[54] Partridge PG. Metall Rev 1967;12:16994.[55] Aydiner CC, Bernier JV, Clausen B, Lienert U, Tome CN, Brown

DW. Phys Rev B 2009;80:16.[56] Tome CN. Personal communication; 2009.[57] Caceres CH. Personal communication; 2009.[58] Caceres CH, Sumitomo T, Veidt M. Acta Mater 2003;51:62118.[59] Mann GE, Sumitomo T, Caceres CH, Griths JR. Mater Sci Eng A

2007;456:13846.[60] Dobron P, Bohlen J, Chmelk F, Lukac P, Letzig D, Kainer KU.

Mater Sci Eng A 2007;462:30710.[61] Heiple CR, Carpenter SH. J Acoust Emission 1987;6:177204.[62] Lamark TT, Chmelk F, Estrin Y, Lukac P. J Alloy Compd

2004;378:2026.[63] Chino Y, Kimura K, Mabuchi M. Mater Sci Eng A 2008;486:4818.[64] Partridge PG. Philos Mag 1965;12:104354.

Investigation of deformation twinning in a fine-grained and coarse-grained ZM20 Mg alloy: Combined in situ neutron diffraction and acoustic emissionIntroductionExperimental techniquesNeutron diffraction techniquePeak intensities (Ihk.l)Peak position (dhk.l)

Acoustic emission technique

Experimental materialMicrostructureTextureTest specimens

Results and discussionMacroscopic stressstrain responseNeutron diffraction resultsTwinning and overall twin volume fractionTwinning and internal stresses

Acoustic emission resultsTwinning vs dislocation slip and AE activityTwinning and AE activityMacroscopic stress relaxation and AE activity

Twin nucleation and twin growthYielding by twinning and the grain size effectSummaryAcknowledgmentsReferences