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Scripta Materialia 68 (2013) 171–174

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Grain refinement due to complex twin formation in rapid hot forgingof magnesium alloy

. Yunping Li,a,⇑ Shuo Wu,a Huakang Bian,b Ning Tang,b Bin Liu,c Yuichiro Koizumia

and Akihiko Chibaa

aInstitute for Materials Research, Tohoku University, Sendai 980-8577, JapanbGraduate School of Engineering, Tohoku University, Sendai 980-8577, Japan

cState Key Lab for Powder Metallurgy, Central South University, Changsha, Hunan, People’s Republic of China

Received 19 May 2012; revised 10 September 2012; accepted 8 October 2012Available online 16 October 2012

Twinning behavior and twin interactions were investigated to characterize grain refinement in magnesium alloys subjected torapid hot forging. Formation of f1 0 �1 2g tensile twins with various twin variants and their subsequent mutual interaction resultedin new boundaries at low strain. With further straining, fragmentation of f1 0 �1 1g-f1 0 �1 2g double twin bands by the formation ofmore complex twins in the interior, accompanied by a drastic decrease in the stored energy, led to grain refinement.Crown Copyright � 2012 Published by Elsevier Ltd. on behalf of Acta Materialia Inc. All rights reserved.

Keywords: Magnesium alloy; Grain refinement; Hot forging; Twinning

Numerous studies have shown that the grainrefinement occurring during hot working of Mg alloysis closely related to twinning activity [1–6], which is espe-cially obvious in low-temperature and/or high-strain-rate regimes where twinning becomes the dominantdeformation mechanism [4–6]. Dynamic recrystalliza-tion (DRX) of Mg alloy has been investigated underconditions of warm deformation by Yin et al. [4], andat room temperature (RT) by Kelley et al. [1] in whichthe nucleation of new grains through DRX, driven bythe distortion energy accumulated from twinning wasreported. The influence of twinning on grain refinementin AZ31B alloys was analyzed by Yu et al. [5] by study-ing a combination of the refined grain size with andwithout the influence of twinning, along with the Ze-ner–Hollomon (Z) parameter; they found that twinningfavors the formation of finer grains. Recently, Ma et al.[6] suggested the idea of twin-induced grain refinementusing electron backscatter diffraction (EBSD) analysis,which is conceptually different from that proposed byYin et al. [4]. Ma et al. found that a large number of re-fined grains were surrounded by rough and incoherentf1 0 �1 2g tensile twin boundaries, which were graduallyrotated into general grain boundaries by external stress.

1359-6462/$ - see front matter Crown Copyright � 2012 Published by Elsevhttp://dx.doi.org/10.1016/j.scriptamat.2012.10.007

⇑Corresponding author. Tel.: +81 22 215 2452; fax: +81 22 2152116; e-mail: lyping@imr.tohoku.ac.jp

However, Ma et al. ascribed the refining mechanism tothe f1 0 �1 2g tensile twin formation, although prominentmisorientation angle distribution peaks also appeared atabout 30� and 38�, which are generally related to thef1 0 �1 1g-f1 0 �1 2g double twins as mentioned in a num-ber of earlier reports [4,5]. It is well known that thef1 0 �1 2g tensile twin is usually formed extensively atlow strains owing to the low critical resolved shear stress(CRSS) required and is prone to disruption by subse-quent deformation or to coalescence with further strain-ing [7]. In order to quantitatively investigate twinning-related grain refinement under hot working, extensiveresearch on the behaviors of tensile f1 0 �1 2g and othertwins, the roles of these twins in the formation of newgrain boundaries, and the orientation relationship be-tween the twins and refined grains are required. How-ever, few reports have addressed these aspectsspecifically. Hence, in this work, we analyzed by EBSD:(i) the twinning behavior, (ii) the twin interactions; and(iii) the specific orientation relationship between the newgrains and the twins at various strain levels in AZ31BMg alloys subjected to rapid hot forging. The resultsof our study revealed an interesting grain-refinementmechanism.

Cylindrical samples of AZ31 alloy (8 mm diame-ter � 12 mm high) were cut from an extruded rod. Com-pressive tests were carried out along the extrusion

ier Ltd. on behalf of Acta Materialia Inc. All rights reserved.

(a)

n

{10-12} twin (b)

172 Y. Li et al. / Scripta Materialia 68 (2013) 171–174

direction (ED) in vacuum at 250 �C using a computer-aided Thermecmaster-Z hot-forging simulator. Thestrain rate used was 10 s�1. The specimens were heatedto the target temperature at a rate of 5 �C s�1 by high-frequency induction. As soon as the samples had beencompressed to the final strain level, they were quenchedto RT with He (4 MPa). Analysis of the microstructurewas carried out at the center of the compressed samplesby EBSD, using data acquisition software (TSL-OIM5.0). Textures of the samples were scanned by the EBSDequipment over a fixed area (200 lm � 130 lm) with astep size of 0.35 lm.

To analyze the grain- and twin-boundary evolutionwith respect to strain, misorientation angle distributionsat various strain levels were analyzed by EBSD, asshown in Figure 1. Only boundaries with misorientationangles >5� were considered in the current study, becausethe boundary angles <5� measured by EBSD were ob-served to be drastically affected by the polishing statecompared to the boundaries with higher angles. Rota-tion axis distributions of some of the prominent peaksare inserted in the corresponding figures. The as-re-ceived sample exhibited a random distribution of misori-entation angles, without a dominant rotation axis(Fig. 1a). In the deformed sample, distinguished bound-ary misorientation peaks in the ranges 5–10�, 29–31�,37–39�, 53–61� and 86–88� about the h2 �1 �1 0i axis;29–31� about the h0001i axis; and 53–61� about theh1 0 �1 0i axis were observed. For the specimens de-formed to a true strain of 0.08, a maximum peak rangeof the boundary misorientation existed around 86–88�about the h2 �1 �1 0i axis, and the other peaks were muchweaker. With the increase in strain, the maximum peakrange shifted from 86–88�/h2 �1 �1 0i to the other peaks.After straining to �0.15, the frequencies of the bound-ary misorientation peaks around 29–31�, 37–39� and53–61� increased to their maximum values (Fig. 1c).The boundary misorientation gradually became ran-

Figure 1. Number fractions of the misorientation angle in AZ31 Mgalloy. Misorientation axis distributions in the crystal coordinate systemare shown for some significant angular ranges for (a) as-receivedsamples and for samples deformed to true strains of (b) 0.08, (c) 0.15,(d) 0.4 and (e) 0.7.

domly distributed at a strain of 0.7 (Fig. 1d, e). Mostof the peaks that occurred at lower strains disappearedat a strain of 0.7 (Fig. 1e), and no noticeable rotationaxis in the above-mentioned angular ranges were de-tected, except for the rotation angle around 29–31�,where the two rotation axes (h0001i and h2 �1 �1 0i) werestill significantly prominent.

Combining the rotation axis distributions obtained atvarious strains and the previously reported results [8], itwas confirmed that the local peak in the range 86–88�about the h2 �1 �1 0i axis can be ascribed to thef1 0 �1 2gh1 0 �1 0i tensile twin. In addition, the peaksaround 5–10�/h2 �1 �1 0i and 53–61�/h1 0 �1 0i boundarieswere found to be related to two separate f1 0 �1 2g twins(f1 0 �1 2g/f1 0 �1 2g) with various twin variants [8]. Also,both the peaks around 29–31�/h2 �1 �1 0i and around 37–39�/h2 �1 �1 0i boundaries were generated from thef1 0 �1 1g-f1 0 �1 2g double twins [9]. The peak aroundthe 53–61�/h2 �1 �1 0i boundary was formed because ofthe f1 0 �1 1gh1 0 �1 0i compressive twin [8]. A preferencefor h0001i was also observed in the rotation angle range29–31� (Fig. 1b–e) throughout the deformation. How-ever, it was not confirmed theoretically whether thiswas related to the twins or slips. From the above results,it can be inferred that by increasing strain to �0.15, themisorientation peak for the f1 0 �1 2g twins increased totheir maximum value at a strain of �0.08, and vanishedgradually. This was accompanied by the strengtheningof some peaks such as those at the 5–10�/h2 �1 �1 0i and53–61�/h1 0 �1 0i boundaries, due to the interaction be-tween the f1 0 �1 2g twin boundaries with various twin-ning variants (Fig. 1a–e).

Figure 2a illustrates the microstructure after slightcompression (to a strain of 0.08) in terms of the inverse

g

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t2

t2

t31

t3

t1

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A

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D50 µm

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Figure 2. (a) Inverse pole figure map of the AZ31B alloy after hotcompression to a true strain of 0.01. (b) A magnified image of therectangular region represented in (a), indicating the formation ofvarious kinds of boundaries by the interaction of f1 0 �1 2g tensile twinswith various twin variants, in which compression direction is parallelto ED. (c) Schematic misorientation axis distribution of the newboundaries formed by the interaction of f1 0 �1 2g tensile twins withvarious twin variants.(d) Point-to-origin profiles along the straightlines A–B and C–D represented in (b).

(a)

30 µm

0 4

1 2

3 4 5

6

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M (b)

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Boundary(15-180o) (2-15o) 87o<2110> 57o<2110>

30, 38o<2110> 60o<1010>

-- -- -

-- {1012} /{1012}

{1012} {1011}

{1011}-{1012}

- - - -

- -

Figure 3. (a) Inverse pole figure map of the AZ31B alloy after hotcompression to a true strain of 0.15. (b) A magnified image of therectangular portion shown in (a) represented as a boundary map,indicating the formation of twinning bands and the role played by theirsubsequent fragmentations in the grain refinement process. (c) Thecorresponding kernel average misorientation map of (b), in which thecompression direction is indicated by the arrow. (d) Schematic inversepole figure of the new grains formed through band formation andsubsequent fragmentation.

Y. Li et al. / Scripta Materialia 68 (2013) 171–174 173

pole figure (IPF) map. The microstructure was charac-terized by the formation of a large number of lenticularf1 0 �1 2g twins. Since the f1 0 �1 2g twin is responsiblefor the dominant deformation behavior at low strainlevels, new boundary formations related to the activityof this twin are very important. A local area (enclosedby a rectangle in Fig. 2a) that included a grain withf1 0 �1 2g twins of various twin variants was considered.The magnified figure of the enclosed area is shown inFigure 2b by IPF map, for identifying the orientationof the parent grains (g), the f1 0 �1 2g twins with twovariants inside the parent grains (t1 and t3), and the twinwithin the t1 twin band (t2). Owing to the complex for-mation of the {1012} twins, four kinds of special bound-aries were generated, which are represented in Figure 2cby a schematic misorientation axis distribution. In Fig-ure 2c, (i) 87�/h2 �1 �1 0i is generated by the f1 0 �1 2g ten-sile twin from the boundaries of g/t1, g/t3, t1/t2; (ii) 60�/h1 0 �1 0i is generated by the f1 0 1 2g/f1 0 �1 2g doubletensile twins from the boundary of g/t2; (iii) 48.3�/h0111i is generated from the interface between thef1 0 �1 2g tensile twin (t3) and the f1 0 �1 2g/f1 0 �1 2gdouble twin (t2); and (iv) 60.4�/h1 0 �1 0i is formed be-cause of the interaction between the two f1 0 �1 2g tensiletwins with various twinning variants(t1/t3). Boundaries(ii) and (iv) possess almost the same rotation angleand axis, although their formation mechanisms are com-pletely different. The Schimid factors (m) correspondingto these twins at various conditions were determinedand are summarized in Table 1. Except for the t1 vari-ant, the m of other variants were close to zero. Thenon-Schimd behavior of both t2 and t3 twins is in con-trast to the formation of the t1 twin, which agrees rea-sonably well with the Schmid predictions. Theoccurrence of t2 and t3 twins with low m could be attrib-uted to the lower CRSS of this mode compared to theother twin or nonbasal slipping systems [10]. In the pres-ent case, the CRSS for this kind of twin has been re-ported to fall in the range 30–50 MPa [10], which ismuch lower than the stress at the strain of 0.01(�120 MPa). In addition, from the misorientation pro-files along the straight lines in both the interior (C-D)and exterior (A-B) of the t1 twin (Fig. 2d), it can be in-ferred that the misorientation generated during defor-mation was accumulated much more in the exteriorthan that did in the interior of twin t1 (Fig. 2b). This im-plies that the formation of these low m twins (t2, t3) pos-sibly correlated to the high stress accumulated by thebasal or nonbasal slips in the nontwinned area. Fromthe misorientation angle distributions shown in Figure 1,the boundaries of types (ii) and/or (iv) contribute to theformation of a new boundary of 53–61�/h1 0 �1 0i beforethe strain level of 0.15 is reached, and the frequency ofoccurrence of boundaries of types (ii) and/or (iv) wasobserved to be much higher than that of type (iii). Thiswas because no prominent misorientation angle distri-

Table 1. Schmid factors m calculated for the twins t1, t2 and t3 shownin Figure 2b.

Twin t1 t2 t3

Schmidt factor, m 0.46 0.022 0.031

bution peak at the corresponding angle for a boundaryof type (iii) (about 48.3�) could be observed at all strains(Fig. 1a–e). This is also considered to be related to thelow frequency of the concurrent occurrence of both t2in t1 and t3 in matrix g, with low m in both cases.

With the increase in strain to 0.15, f1 0 �1 2g twinsgrew gradually and covered most of the matrix, resultingin the formation of new boundaries owing to the inter-action of these twins with various variants, and an over-all rotation of the basal plane into the compressiondirection by �87� (Fig. 3a) was observed. In additionto the boundaries formed by f1 0 �1 2g twins in the pre-ceding stage, a few twin bands with widths ranging from3 to 7 lm were formed in the rotated matrix, and subse-quently fine grains were formed by their fragmentation.In this strain regime, the misorientation angle distribu-tion, as shown in Figure 1, also indicated that complextwinning related to the generation of peaks of 29–31�,37–39� and 53–6� (Fig. 1) occurred. To investigate thisfurther, a selected area of Figure 3a that included severalof these bands was magnified as shown in Figure 3b andc, which represent the boundary map and kernel averagemisorientation (KAM) map, respectively. EBSD con-firmed that most of the boundaries of these bands pos-sessed a rotation relationship of 30� or 38�/h2 �1 �1 0idue to the ð1 0 �1 1Þ-ð1 0 �1 2Þ double twin. Along the inte-rior of the double twin bands, many fresh boundarieswith specific rotation relationships such as 57�/h2 �1 �1 0i from (1011) compressive twin, 60�/h1 0 �1 0ifrom the f1 0 �1 2g/f1 0 �1 2g twin boundary, as well as87�/h2 �1 �1 0i from the f1 0 �1 2g tensile twin wereformed, and played a key role in the fragmentation pro-cess (Fig. 3a). These bands were observed to be formedextensively inside the matrix, and intersected with eachother, which was accompanied by fragmentation at highstrains (Fig. 4a). With further straining, the matrix wasobserved to be refined homogeneously into fine grains

Figure 4. Inverse pole figure maps of the AZ31B alloy after hotcompression to true strains of (a) 0.4 and (b) 0.7. In (b) thecorresponding grain size distribution at a true strain of 0.7 is inserted.

174 Y. Li et al. / Scripta Materialia 68 (2013) 171–174

(Fig. 4b). The grain size distribution of the refinedmicrostructure (Fig. 4b) was obtained by EBSD. Thedistribution of the refined grains was relatively uniform,with a mean diameter of about 5.54 lm, which was closeto the widths of the previous double twin bands. Thisobservation further supports the hypothesis that thegrain refinement in the current situation can be ascribedto the fragmentation of these double twin bands.

The formation of the ð1 0 �1 1Þ-ð1 0 �1 2Þ double twinbands and their subsequent fragmentation resulted inlower KAM value than that in the matrix (Fig. 3c).However, the decrease in KAM value was also observedin the neighboring locations, as indicated by the arrowin Figure 3c. These areas exhibiting low KAM weremostly surrounded by straight low-angle boundaries(2–15�), which were almost parallel to f1 0 �1 1g as indi-cated by the black straight lines shown in Figure 3c. Thiscan be ascribed to the presence of slips in f1 0 �1 1g, withthe complex interactions of these ha + ci slips possiblyleading to a local rotation of the lattice during the defor-mation, resulting in an orientation favoring double twinformation and subsequent fragmentation (Fig. 3b). Inthis sense, the spread of the double twin bands is likelyto be attributed to the successive accumulation of slipsin f1 0 �1 1g.

The formation of refined grains by complex twinningwithin the double twin bands tended to lower the polar-ity of the texture. This could be observed from the ori-entation map of the six numbered grains (Fig. 3b). Mrepresents the matrix. Grains 1–5 were formed by thefragmentation of the double twin bands surrounded byvarious kinds of twin boundaries, and grain 6 was di-rectly formed by the ð1 0 �1 1Þ-ð1 0 �1 2Þ double twin.The formation of new grains due to the fragmentationof the double twin bands tended to be more randomthan both the matrix and band (Fig. 3c).

Twinning-induced DRX has been described in a fewstudies [4,11]. In general, this can be explained on thebasis of the twin providing a suitable site for nucleationeven at RT [1] due to the distortion energy stored be-cause of the higher CRSS required during twin forma-tion. This nucleation-growth process is generallyobserved at low or intermediate strain rates, as longerdurations are necessary for atomic diffusion or latticereorganization. Nucleation and growth of new grainsby diffusion at RT inside the ð1 0 �1 1Þ twin band was ob-served in pure Mg [1]. However, there were very fewnucleated grains. According to Chao et al. [12], in theheavily cold-drawn (�61%) AZ31 alloy, even after

annealing for 3 min, only a few recrystallized grainswere observed. This implies that the duration of heattreatment was insufficient for obtaining fully refinedmicrostructures by the usual nucleation mechanism.Nucleation and growth of new grains may not domi-nantly occur in the current situation, because the com-pression was carried out at high temperatures, wherethe driving force for nucleation is much lower than thatin the cold-worked sample. In addition, inclusive of thecooling stage, the maximum period of the hot-workingprocess is restricted to only several seconds, which isconsidered to be insufficient for both grain nucleationand subsequent growth, throughout the sample. Whenthe alloys used in the present study were compressedat high strain rates, neither discontinuous DRX by thebulging of grain boundaries nor the continuous DRXoccurred by the accumulation of prismatic slips to formhigh-angle grain boundaries. A high strain rate is con-sidered to inhibit the diffusion-controlled discontinuousDRX owing to the short duration of treatment, therebysignificantly enhancing twin formations. Grain refine-ment by the formation of these twinning bands and theirfragmentation is expected to be an energy-decreasingprocess, owing to the accumulation of pyramidal slips,which produced gradual local lattice rotation and ener-getically and mechanically reduced the distortion energyby the selective formation of complex twins.

In summary, during hot-working process at highstrain rates, the formation of specific boundaries atlow strains was closely related to the complex formationof f1 0 �1 2g tensile twins with various variants. At highstrains, fragmentations of the f1 0 �1 1g-f1 0 �1 2g doubletwin bands by the formation of complex twin bound-aries played a key role in the grain refinement process.It was found that both the double twin band formationand their fragmentation are processes that reduce thestored energy, and are the result of the activity of thepreceding hc + ai slips in the f1 0 �1 1g pyramidal planes.

This research was supported by the Grant-in-Aid for Young Scientists B (24760591), Japan Societyfor the Promotion of Science.

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