Microstructural evolution and grain growth mechanism of pre-twinned magnesium alloy during annealing

Ye Jin Kim ,Jong Un Lee ,Young Min Kim ,Sung Hyuk Prk,*

aSchool of Materials Science and Engineering,Kyungpook National University,Daegu 41566,Republic of Korea

b Implementation Research Division,Korea Institute of Materials Science,Changwon 51508,Republic of Korea

Abstract The microstructural evolution and underlying grain growth mechanism of a {10-12}-twin-containing Mg alloy during annealing are investigated through quasi in situ electron backscatter diffraction measurements of a rolled AZ31 alloy subjected to precompression along the rolling direction (RD).The precompressed material shows a partially twinned structure consisting of a twinned region and a residual matrix region,and this structure changes to an almost twin-free structure consisting of grown grains with serrated grain boundaries in twinand matrix-originated regions after annealing at 250°C for 1h.In addition,the average grain size almost doubles and the internal strain energy decreases significantl after annealing.These microstructural variations are induced mainly by grain growth through the strain-induced migration of high-angle grain boundaries without the movement of twin boundaries.The twinned region of the precompressed material has higher stored strain energy than the residual matrix region because the crystallographic orientation of the former region is favorable for basal slip and because of the occurrence of the dislocation transmutation reaction in the twins.For reducing the total strain energy accumulated in the precompressed material,the residual matrix region—having lower stored strain energy—preferentially grows while consuming the twinned regions formed in the surrounding grains during annealing.As a result,the area fraction of grains with a matrix texture increases whereas that of grains with a twin texture decreases after annealing.The twin texture intensity increases significantl and this texture becomes more concentrated along the RD because the highly RD oriented twins tend to grow during annealing on account of their fairly low stored strain energy.

Keywords: Magnesium;Twin;Annealing;Grain growth;Strain energy.

1.Introduction

Environmental issues such as global warming and abnormal weather phenomena are increasingly becoming signifi cant problems worldwide;in the transportation industry,these problems have been addressed through tightening of regulations on the carbon dioxide emissions and fuel efficien y of vehicles and placing greater emphasis on the weight reduction of vehicles.From this viewpoint,Mg alloys have recently attracted great attention in the transportation industry because they have lower density and higher specifi strength than other commercially available structural metals such as steels and Al alloys.When cast Mg alloys are subjected to hot metal forming processes (e.g.,rolling,extrusion,and forging),casting defects such as blow holes and pores are generally eliminated and dynamic recrystallization occurs;consequently,wrought Mg alloys have much higher mechanical properties than their cast Mg counterparts[1-3].Therefore,wrought Mg alloys can be applied to not only interior components of a vehicle but also body and chassis components requiring higher strength;such application leads to greater weight reduction of vehicles.However,the applicability of wrought Mg alloys to the fabrication of automotive components is still limited,one of the reasons for which is their low ductility and formability at room temperature (RT) because of their strong basal texture.

In an effort to weaken the basal texture of wrought Mg alloys,many studies have attempted to control the recrystallization behavior during hot forming through the addition of alloying elements to pure Mg or commercial Mg alloys.The addition of rare-earth (RE) elements such as Gd,Ce,Y,and Nd has been reported to cause weakening,tilting,or splitting of the basal texture of wrought Mg alloys [4-12].Activation of non-basal slip and contraction twinning during hot deformation is known to be easier in RE-containing Mg alloys than in RE-free Mg alloys,and this activation results in a uniform formation of shear bands in the material [13-18].Since the grains recrystallized in the shear bands have off-basal orientations,an RE texture that is tilted～25° from the normal direction (ND) is formed in rolled or extruded Mg sheets containing RE elements [14,19].However,the addition of RE elements causes undesirable increases in the cost and specifi gravity of the material,which,in turn,lead to reductions in the cost competitiveness and weight reduction effect of the fina Mg products.For these reasons,various other methods for modifying the texture of wrought Mg alloys have been explored,one of which is the induction of static recrystallization through cold forming and subsequent heat treatment[20-22].Grain boundaries,shear bands,deformation twins,and second phases are the main nucleation sites for recrystallization in deformed Mg alloys [23-26].Because of the hexagonal close-packed (HCP) crystal structure and relatively low stacking fault energies of Mg alloys,dislocations easily accumulate in the material during cold forming,especially at grain boundaries.Accordingly,new recrystallized grains are predominantly formed at the grain boundaries in the coldformed alloys during their subsequent annealing.However,the crystallographic orientations of recrystallized grains formed at the grain boundaries are similar to those of the initial parent grains,so the nucleation of new grains at the grain boundaries is not effective in inducing a texture change [27,28].In contrast,recrystallized grains that nucleate at shear bands—which refer to locally deformed regions with high shear strain—have relatively random orientations;however,these recrystallized grains cannot grow beyond the shear band region [29,30].Therefore,it is difficul to bring about a significan change in material texture via static recrystallization at grain boundaries and shear bands.

Twinning,which is an important deformation mechanism in HCP metals,is capable of accommodating plastic deformation,reducing the effective grain size,and changing the crystal orientation [31-34].In Mg,which has ac/aratio of 1.623—close to the ideal value of 1.633—{10-12} twinning and {10-11} twinning predominantly occur in thec-axis extension and contraction stress modes,respectively;moreover,lattice reorientations of 86.3° and 56° occur in the {10-12}twins and {10-11} twins,respectively [35].Therefore,the basal texture of Mg alloys can be modifie via introduction of such deformation twins through cold forming and subsequent induction of recrystallization in the twinned region with changed orientations through heat treatment.Several studies have been conducted with the aim of weakening the texture of Mg alloys via static recrystallization at twins,with particular focus on the recrystallization behavior at {10-11} contraction and {10-11}-{10-12} double twins during annealing [36,37].When {10-11} contraction and {10-11}-{10-12} double twins are formed during deformation,large strain is accumulated in them because their lattice orientations are favorable for basal slip and because of the low mobility of twin boundaries [38].Accordingly,these {10-11} contraction and {10-11}-{10-12} double twins provide a high driving force for recrystallization,and static recrystallization preferentially occurs at these twins during subsequent annealing.Since the recrystallized grains formed at these twins have relatively random orientations,they can induce texture weakening.However,these grains do not grow beyond the twin boundaries and are easily consumed by the growth of recrystallized grains nucleated at the surrounding grain boundaries[37,39].Consequently,the texture variation caused by the recrystallization at {10-11} contraction and {10-11}-{10-12}double twins is insignificant {10-12} Extension twins are most commonly observed in Mg alloys and{10-12}twinning acts as a main deformation mechanism at low temperatures below～200°C.Since {10-12} extension twins are formed at low stress levels and they grow easily owing to the low critical resolved shear stress (CRSS) and high mobility of twin boundaries,most grains of wrought Mg alloys with a strong basal texture can be transformed such that they consist mainly of the {10-12} twinned region through the application of～6% plastic deformation underc-axis extension stress conditions;this transformation consequently results in a drastic change in the material texture [31,40,41].However,few in-depth studies have been conducted on the microstructural changes that occur in deformed Mg alloys containing {10-12} twins during subsequent annealing treatment.Therefore,this study systematically investigates the variations in the microstructural characteristics of a wrought Mg alloy with initial{10-12} twins during its annealing and the underlying mechanisms of these variations.To this end,numerous {10-12}twins are introduced into a rolled AZ31 alloy through compression at RT,and the characteristics of the twinned structure and their evolution during annealing are analyzed viaquasi in situelectron backscatter diffraction (EBSD) observations.

2.Experimental procedure

The material used in the present study was a commercial hot-rolled AZ31 Mg alloy (Mg-3Al-1Zn-0.5Mn,wt%)with a thickness of 20mm.The rolled alloy was homogenized at 400°C for 10h to remove residual stress and to induce uniform solute distribution;the homogenized alloy is hereafter referred to as the as-rolled material.For precompression,a rectangular bar with dimensions of 60mm(length)×30mm (width)×20mm (thickness)—which correspond to the rolling direction (RD),transverse direction (TD),and ND,respectively—was machined from the as-rolled material.For introduction of {10-12} twins into the machined bar,it was compressed to a plastic strain of 6.0% along the RD by means of an Instron 8516 testing machine at a strain rate of 10-3s-1and RT;the compressed sample is hereafter referred to as the precompressed material.

Fig.1.Microstructural characteristics of as-rolled material:(a) inverse pole figur (IPF) map,(b) (0001) and (10-10) pole figures and (c) grain orientation distribution. davg denotes the average grain size.

The microstructure and texture of the as-rolled material were analyzed by EBSD measurements with a step size of 1.3μm in a 670×536 μm2scanning area.In addition,the microstructural variations during annealing were systematically investigated by performingquasi in situEBSD measurements at the same position on the RD-TD plane of the precompressed material before and after annealing.The annealing treatment was performed at 250°C for 1h in a box furnace,after which the annealed sample was air-cooled.For thequasi in situEBSD measurements,a sample with dimensions of 2mm(length)×2mm (width)×1mm (thickness)—which was machined from the precompressed material—was progressively ground and polished with abrasive papers,3 μm and 1 μm diamond compound pastes,and a colloidal silica solution.After the firs EBSD measurement of the polished sample,it was annealed and then polished very slightly with only the colloidal silica solution to remove the surface oxide layer formed during annealing.Subsequent EBSD measurements of the annealed sample were performed in the same scanning area.Thequasi in situEBSD measurements were performed using a Hitachi SU-70 field-emissio scanning electron microscope equipped with a Hikari camera with a step size 1.0μm in a 550×600 μm2scanning area.The Tex-SEM Laboratories orientation imaging microscopy (TSL OIM) analysis software(version 7.0) was used to analyze the EBSD data.Only those EBSD data with a confidenc index greater than 0.1 were used for the analyses in order to ensure reliability of the analysis results.

3.Results and discussion

3.1.Initial microstructures of as-rolled and precompressed materials

Fig.1 shows the microstructural characteristics of the asrolled material.It has a twin-free microstructure comprising equiaxed recrystallized grains with an average grain size of 44.1μm and a strong basal texture with a maximum texture intensity of 14.2 multiples of random distribution(m.r.d.).Thec-axes of most grains are oriented almost parallel to the ND,whereas theira-axes are randomly oriented along the RDTD plane (Fig.1b).Fig.1c shows the distribution of the angleθbetween thec-axis of the grains and the RD,which reveals that for almost all grains (i.e.,area fraction of 99%),this angle is larger than 55° This result indicates that when precompression is applied to this material along the RD,{10-12} extension twinning can occur in all grains because they are in thec-axis extension stress state.

Fig.2 shows inverse pole figur (IPF) maps and (0002)and (10-10) pole figure of the precompressed material.Numerous {10-12} twins are formed by precompression,which leads to (a) grain refinemen (from 44.1 to 19.2μm) owing to the formation of twin boundaries and (b) significan texture variation (from the ND-oriented texture to the RDoriented texture) owing to lattice reorientation by twinning.Since {10-12} twinning causes lattice reorientation of 86.3°from the matrix,a twin texture with thec-axes oriented along the precompression direction (i.e.,RD) develops in the precompressed material (Fig.2a).Accordingly,the twinned region and residual matrix region of the precompressed material can be divided according to the angleθbetween thec-axis and the RD;a region with 0°≤θ≤45°is the twinned region,and that with 45°<θ≤90° is the residual matrix region.The area fractions of the twinned region and residual matrix region are 79% and 21%,respectively (Fig.2b and c).Some grains are completely twinned;that is,they consist entirely of {10-12} twins and do not have any residual matrix region.In contrast,other grains are partially twinned and they consist of both the twinned region and the residual matrix region(Fig.2b).This difference in twinning behaviors is attributed to the considerable variation in the extent of twin formation in an individual grain with its size and crystallographic orientation.Our previous study [42] investigated the twinning behavior of a rolled AZ31 alloy with a relatively wide grain size distribution of 10-80μm under compression to a plastic strain of 6% along the RD at RT;the study revealed that the grain size(GS),Schmid factor for {10-12} twinning (SFtwin),and twin fraction (TF) exhibit the relationTF/SFtwin=109.8+1.45·GS.According to this relation,the area fraction of the twinned region formed in a grain of the precompressed material examined in the present study can change from 30% to 100%depending on the size and orientation of the grain.Despite the different extents of twin formation in different grains,twinning occurs in all grains in the present study;consequently,the residual matrix region has a lath-shaped morphology with a high aspect ratio of 2.04 (Fig.2c).The RD-oriented twin texture has an averageθof 18.7° and its intensity is high as 15.3 m.r.d.Moreover,one prismatic pole of the twin texture is oriented almost parallel to the ND (Fig.2b).These characteristics of the twin texture are consistent with previous results for the {10-12} twin texture in other rolled Mg alloys[40,43,44].The average angleωbetween thec-axis and the ND in the residual matrix region of the precompressed material is 31° (Fig.2c),which is higher than theαvalue of the grains of the as-rolled material (18°).This higherαin the precompressed material is attributable to the fact that grains with a lowerαvalue in the as-rolled material have a higher Schmid factor (SF) for {10-12} twinning under compression along the RD,as a result of which twinning preferentially occurs in these grains during precompression [31,43].

Fig.2.Microstructural characteristics of precompressed material:IPF maps and (0001) and (10-10) pole figure of (a) total region (=twinned region+residual matrix region),(b) twinned region,and (c) residual matrix region. davg denotes the average grain size. ftwin and fmatrix denote the area fractions of the twinned region and residual matrix region,respectively.

3.2.Microstructural variations during subsequent annealing

Fig.3 showsquasi in situEBSD results of the precompressed material before and after annealing at 250°C for 1h.The microstructure changes from the highly twinned structure before annealing to an almost twin-boundary-free structure after annealing,and the average grain size doubles from 19.2μm to 38.5μm (Fig.3a and d).As seen from the boundary maps before and after annealing (Fig.3b and e,respectively),the length of {10-12} twin boundaries decreases significantl from 56.7mm to 8.2mm;therefore,the length ratio of the twin boundaries to all the boundaries decreases greatly from 50.0% to 18.5%.Furthermore,the equiaxed grain structure changes to an irregular grain structure with serrated or bulged high-angle grain boundaries (HAGBs).Fig.3c and f shows the kernel average misorientation (KAM) maps before and after annealing,respectively.In these maps,most regions prior to annealing are green and red in color,which represent high KAM values;in contrast,the regions marked in blue,which represent low KAM values,occupy a considerable portion of the material after annealing.The annealing treatment leads to a decrease in the average KAM value from 1.27 to 0.69,which implies a significan reduction in the stored strain energy in the material upon annealing.

Fig.3. Quasi in situ electron backscatter diffraction (EBSD) results showing microstructural variations during annealing after precompression:(a,d) IPF maps,(b,e) boundary maps,and (c,f) kernel average misorientation (KAM) maps of precompressed material (a-c) before and (d-f) after annealing. davg, ltwin,and KAMavg denote the average grain size,length of {10-12} twin boundaries,and average KAM value,respectively.

Fig.4 shows IPF maps,boundary maps,and grain orientation spread (GOS) maps of the white rectangular regions marked in Fig.3a and d,which were drawn to analyze in detail the microstructural changes occurring during annealing.A comparison of the microstructure before and after annealing (Fig.4a and b,respectively) reveals that the microstructural change during annealing is caused by grain growth.This growth behavior of preexisting grains is attributed to the migration of HAGBs (Fig.4c and d).Such grain boundary migration is known to be affected by various factors such as the grain boundary energy,stored strain energy,surface energy,chemical driving force,magnetic field elastic energy,and temperature gradient [45,46].Among these,the grain boundary energy and stored strain energy are considered to be influencin factors in this study.The GOS maps depict the amount of strain in each grain and the strain distribution of grains in a material.The GOS values of the individual grains before annealing are different (Fig.4e).This difference in the stored strain energies in the individual grains of the precompressed material is attributed to the variation in the SF values for slip and twinning under precompression with the crystallographic orientation of the grain and the dependence of activation stresses for slip and twinning on the grain size [42,47].Since the strain hardening caused by deformation twinning is weaker than that caused by dislocation slip,the stored strain energy in a deformed grain varies with the degree of twinning activation in the grain [48-50].The movement directions of the grain boundaries during annealing are marked in the IPF,boundary,and GOS maps before annealing (see the white arrows in Fig.4a,c,and e).From the GOS distribution and the grain boundary movement directions,it can be seen that the grain boundaries move from the region with lower strain energy to the surrounding region with higher strain energy(Fig.4e).This boundary movement caused by the difference between the stored strain energies in adjacent regions is known as strain-induced boundary migration (SIBM),which occurs to reduce the overall strain energy of the material[28,29,51-56].Accordingly,a larger difference between the stored strain energies in two adjacent grains leads to faster grain boundary migration owing to the higher driving force for SIBM [28,45,52-56].In some grains,grain growth hardly occurs during annealing but their GOS values decrease after annealing (see the black arrows in Fig.4e and f).This means that during annealing,the strain energy accumulated in the grains decreases without the occurrence of the migration of HAGBs via SIBM.This dissipation of stored strain energy is a result of a decrease in the dislocation density through a recovery process involving dislocation climb and cross-slip during annealing [39,57,58].Therefore,the stored strain energy of the precompressed material is reduced by both SIBM and static recovery during annealing,and this reduction,in turn,leads to a decrease in both the length of low-angle grain boundaries (from 3.49mm to 1.3mm) and the average GOS value (from 1.41 to 0.91).

Fig.5. Quasi in situ EBSD results showing twin boundary variation during annealing:HAGB and {10-12} twin boundary maps (a) before and (b) after annealing.(c) Superimposed boundary map of those in (a) and (b).(d) KAM map after annealing.

Fig.5 shows grain and twin boundary maps before and after annealing in the same regions as those depicted in Fig.4.The length of {10-12} twin boundaries decreases considerably after annealing (Fig.5a and b).Comparison of the boundary maps before and after annealing (Fig.5c)reveals that the positions of the twin boundaries remaining after annealing hardly change from the initial positions before annealing.This almost no change indicates that the twin boundaries do not move during annealing and that the microstructural change during annealing occurs through only the grain boundary migration.The boundary migration velocity (vb) is proportional to both the boundary mobility(mb) and the driving force for boundary migration (P),i.e.,vb=mb·P[45].The boundary mobility is expressed asmb=mo·exp(-Q/kT),whereQ,m0,k,andTare the activation energy,material constant,Boltzmann constant,and temperature,respectively [59].According to this equation,at a given temperature in a material,a boundary with a lower Q has higher mobility than that with a higher Q.The Q value of HAGBs is generally lower than that of twin boundaries because HAGBs,which are thought to be originally amorphous to some extent,contain considerably more lattice defects than do twin boundaries,which have a coherent or semi-coherent interface with the matrix [59-61];therefore,HAGBs have higher mobility than twin boundaries.In addition,the driving force for boundary migration increases linearly with increasing grain boundary energy[45].Results of previous molecular dynamics simulations show that the energies of HAGBs and{10-12} twin boundaries in a Mg alloy are 1.3-0.4J/m2[62-64] and 0.1J/m2[65,66],respectively.This result indicates that HAGBs have a higher driving force for boundary migration than do {10-12} twin boundaries.As a result,the boundary migration velocity of HAGBs is considerably higher than that of {10-12} twin boundaries,and consequently,the microstructural evolution of the precompressed material during annealing is governed by the migration of HAGBs.

Fig.4. Quasi in situ EBSD results showing high-angle grain boundary(HAGB) migration during annealing in region marked by white rectangle in Fig.3:(a,b) IPF maps,(c,d) grain boundary maps,and (e,f) grain orientation spread (GOS) maps (a,c,e) before and (b,d,f) after annealing.lLAGB and GOSavg denote the length of low-angle grain boundaries (2°-15°)and the average GOS value,respectively.

An HAGB is unstable and has high energy because of its disordered crystal structure in the boundary plane.In contrast,a twin boundary has higher thermal stability than an HAGB because a twin has high coherency with its parent matrix and is formed by twin dislocations to small extents [39,67].Therefore,in this study,twin boundaries remain unchanged during annealing at the relatively low temperature of 250°C,because of which they hardly contribute to the variation in the grain structure,in contrast to the significan contribution of grain boundaries.Despite the lack of movement of twin boundaries,their quantity decreases considerably after annealing because grains with lower stored strain energy grow while consuming surrounding twinned grains.The grains that grow through SIBM have low strain energy and a twin-free structure,whereas those containing twin boundaries remaining after annealing still have relatively high strain energy(Fig.5b and d).Therefore,the microstructural variations,including variations in the grain size,boundary fraction,and stored strain energy,during annealing of the precompressed material are governed by the grain growth behavior induced by SIBM,rather than static recovery and twin boundary motion.

3.3.Preferential growth mechanism of residual matrix region during annealing

The twin-free microstructure formed after annealing can be divided into two regions,matrix-originated region (MOR)and twin-originated region (TOR),which are formed by the growth of the residual matrix region and twinned region,respectively [68].Namely,the MOR and TOR of the annealed material are regions with 0° ≤θ≤45° and 45°<θ≤90°,respectively.Fig.6 shows the IPF maps of the residual matrix region and twinned region of the precompressed material and the MOR and TOR of the annealed material.The area fraction of the MOR (32%) is larger than that of the residual matrix region (21%),whereas the area fraction of the TOR (68%)is smaller than that of the twinned region (79%).These area fraction differences indicate that the growth of the residual matrix region during annealing is more pronounced than that of the twinned region,and the residual matrix region partially encroaches on the twinned region during the growth process.Xin et al.[69] reported that during high-temperature annealing (at 450°C) of a wrought AZ31 alloy containing {10-12}twins,narrow twin bands tended to be consumed by the large residual matrix around the twins,which induced a transformation of the twin texture to the matrix texture.In contrast,when the twin bands were much larger than the residual matrix,they preferentially consumed the matrix through thermally activated twin boundary migration (TATBM),and the texture change also co-occurred in reverse.Application of the results of Xin et al.to the twin size effect in the present study indicates that the twinned region may encroach on the residual matrix region during annealing because the area fraction of the twinned region (79%) is much larger than that of the residual matrix region (21%);consequently,the area fraction of the TOR after annealing would be larger than that of the twinned region (79%).However,in this study,the migration of {10-12} twin boundaries does not occur during annealing,because the thermal energy at the relatively low temperature of 250°C is insufficien to cause TATBM (Fig.5),and the fraction of the area corresponding to the twinned region decreases from 79% to 68% after annealing.

Fig.6.IPF maps of (a) residual matrix region and (b) twinned region before annealing and (c) residual-matrix-originated region (MOR) and (d) twinoriginated region(TOR)after annealing.The residual matrix region and MOR have 45° <θ ≤90° and the twinned region and TOR have 0° ≤θ ≤45°,where θ is the angle between the c-axis and the rolling direction(RD).fmatrix,ftwin, fMOR,and fTOR denote the area fractions of the residual matrix region,twinned region,MOR,and TOR,respectively.

Fig.7.(a) IPF map of precompressed material and (b) values of Schmid factor (SF) for basal slip under compression along RD for residual matrix region and twinned region in 20 grains marked in (a).

As described in Section 3.2,the microstructural evolution during annealing is caused by grain growth through SIBM,and not TATBM.Since the grain boundary migration via SIBM occurs from a region with lower stored strain energy toward an adjacent area with higher stored strain energy,the pronounced growth behavior of the residual matrix region can be attributed to the lower stored strain energy in this region.As basal slip is dominantly activated during RT deformation,the strain energy accumulated in the residual matrix region and twinned region is proportional to the SF for basal slip under precompression.In other words,in a region with a higher SF,dislocation slip occurs more vigorously during precompression at RT,as a result of which this region has higher stored strain energy.For comparison of the SF for basal slip between the residual matrix region and the twinned region,20 partially twinned grains of the precompressed material were randomly selected and the SF values of the residual matrix region and twinned region in each grain were calculated(Fig.7).The calculation results reveal that in 16 of the 20 examined grains,the twinned region has a higher SF than the residual matrix region;that is,the residual matrix region has a higher SF in only 4 of the 20 examined grains.The average value of SFs of the twinned regions in the 20 grains (0.33) is higher than that of the corresponding matrix regions (0.29),which means that the twinned region is more favorably oriented for the activation of basal slip during precompression.Accordingly,{10-12} twins are formed in the grains at the early stage of precompression owing to the ND-oriented initial texture and low CRSS for {10-12} twinning,and subsequently,many dislocations are formed in the twins during further precompression deformation owing to the high SF for basal slip of the twins.

Fig.8.Variations in SF for basal slip of a grain and a {10-12} twin formed in it as a function of angle θ between c-axis of grain and loading axis.The curves of the maximum and minimum SF values of the grain and twin were generated using the angle α between the a-axis of the grain and the loading axis under the condition that the loading axis was projected onto the basal plane (refer to the schematic illustrations on the right side of the figure) The marked value of 79.7° is the average value of angle θ between the compressive loading axis (i.e.,RD) and the c-axis of grains of the as-rolled material (refer to Fig.1c).

The reasons for the higher average SF of the twinned region of the precompressed material than that of the residual matrix region but a higher SF of the residual matrix region in some grains (Grains 4,11,12,and 13 in Fig.7) were analyzed via calculation of the variations in the SF for basal slip with the crystallographic orientations,i.e.,θandα(whereαis the angle between thea-axis of a grain and the loading direction),of the initial grain.Fig.8 shows the variations in the SF for basal slip of two grains,one havingαof 0° and the other havingαof 30°,with a variation in theirθf(wàn)rom 50° to 90° This figur also shows the variations in the SF of{10-12} twins formed in these two grains at a givenθunder the assumption that {10-12} twinning occurs in one twin variant with the maximum SF for {10-12} twinning under compression along the loading direction.In this study,under compression along the RD,the grains of the as-rolled material haveθin the range of 55°-90° (Fig.1c) andαin the range of 0°-30° because the material has an ND basal texture with thec-axes preferentially oriented along the ND and a uniform prismatic texture with thea-axes randomly oriented along the RD-TD plane (Fig.1b).When a grain has the highestαof 30°,the SFs of both the grain and the twin formed in it are the lowest at a givenθ(see the SF curves with solid squares in Fig.8).In contrast,when a grain has the lowestαof 0°,the SFs of both the grain and the twin formed in it are the highest at a givenθ(see the SF curves with open circles in Fig.8).Comparison of the SFs of the grain and twin reveals that atα=30°,the SF of a grain is slightly higher than that of the twin formed in it over theθrange of 50°-88°However,atα=0°,the SF of the twin formed in the grain is considerably higher than that of the grain over the entire consideredθrange;moreover,the difference in the SF between the twin and the grain increases with increasingθowing to a gradual decrease in the SF of the grain.Consequently,the average SF of grains (gray-shaded region in Fig.8) over the complete ranges of orientations considered for calculations in Fig.8 (i.e.,0° ≤α≤30° and 50° ≤θ≤90°) is considerably lower than that of the twins formed in the grains(light-red-shaded region in Fig.8).

Additional analyses of variations of the SFs of grains and twins at variousαvalues reveal that atα=～23°,a grain and a twin formed in it have similar SF values at a givenθ.These calculation results of the SF suggest that in this study,in low-αgrains (0° ≤α≤23°) of the as-rolled material,the twinned region may have a higher SF than the residual matrix region,whereas in grains with rather highα(23°≤α≤30°),the residual matrix region may have a higher SF.To verify this suggestion,the angular relations of the crystallographic orientation and the precompression direction(i.e.,θandα) with the SFs of the twinned and residual matrix regions were obtained from the EBSD results (Fig.9)for 4 of the 20 grains (Grains 19,3,4,and 11) marked in Fig.7.Indeed,in Grain 19,which has a lowαof 3°,the SF of the twinned region (0.35) is higher than that of the residual matrix region (0.29) (Fig.9a).In addition,although Grain 3 has a higherα(11°) than Grain 19,the former has a very highθof 89°,and as a result,the SF of its twinned region (0.28) is significantl higher than that of its residual matrix region (0.01) (Fig.9b);this result is consistent with the calculated result that the SF difference increases with increasingθ(Fig.8).In contrast,in Grain 4—having the largestαof 30°—the SF of the twinned region (0.27) is lower than that of the residual matrix region (0.35) (Fig.9c).Similar to Grain 4,the other three grains in which the residual matrix region has a higher SF than the twinned region as shown in Fig.7 (Grains 11,12,and 13) have relatively higherαexceeding 22°,which coincides with the calculated SF curves.Given that the as-rolled material has a uniform distribution ofαand a high averageθof 79.7° (Fig.1),it is obvious that the SF for basal slip under precompression is higher for the twinned region than for the residual matrix region;according to the calculated SF variations,atθof 79.7°,the average SF of twins (0.26) is 64% higher than that of grains (0.16) (Fig.8).Therefore,during precompression,more dislocations accumulate in the twinned region than in the residual matrix region.In addition to reporting the accumulation of more numerousbasal dislocations because of the higher SF of the twinned region,Kadiri and Oppedal [70] proposed the dislocation transmutation theory wherein the dislocation density inside twins increases rapidly through twinning-induced dislocation transmutation,which results in a higher hardening rate of the twins than of the residual matrices.Recently,Wang and Agnew [71] performed transmission electron microscopy observations of a rolled AZ31 alloy subjected to compression along the RD at RT in order to observe the dislocation transmutation reaction at the boundaries of {10-12} twins.They reported that 〈a〉 matrix dislocations transmute into 〈c+a〉 dislocations in the {10-12} twins and that these 〈c+a〉 dislocations cause hardening within the twin itself during subsequent deformation because these dislocations are most probably sessile.Given that the material that they used and the deformation conditions that they adopted in their study are identical to those in the present study,it is reasonable to conjecture that sessile〈c+a〉dislocations are probably formed in the {10-12} twins during precompression in this study and that their formation may be a contributing factor to the higher stored strain energy in the twinned region than in the residual matrix region.Therefore,during subsequent annealing,the grain boundaries preferentially move toward the twinned region,which has higher stored strain energy;consequently,the area fraction of the twinned region decreases after annealing.

Fig.9.IPF maps of (a) Grain 19,(b) Grain 3,(c) Grain 4,and (d) Grain 11,all marked in Fig.7a,along with crystal orientations and SF for basal slip of residual matrix region (M) and twinned region (T) of these grains. ORgrain denotes the angular relationship between a grain and the compressive loading direction (i.e.,RD). SFmatrix and SFtwin denote the values of the SF for basal slip under loading along the RD for the residual matrix region and twinned region,respectively.

The growth of the residual matrix region into the surrounding twinned region can be confirme from the results of thequasi in situEBSD measurements.It can be seen in the IPF map before annealing (Fig.10a) that one partially twinned grain comprising a residual matrix (M1) and twins with the same twinning variant (T1) is surrounded by {10-12} twins formed in four adjacent grains (T2,T3,T4,and T5).During annealing,the matrix grows substantially while consuming the surrounding twins,and this growth consequently results in the formation of a coarse MOR (Fig.10b).However,even after annealing,the twin present in the grain (T1) remains within the grown matrix (i.e.,MOR) because the twin boundaries do not move during annealing.In a certain grain of the precompressed material,the twinned region generally has higher stored strain energy than the residual matrix region (Fig.8).However,preferential growth of the twinned region can occur in some grains;this scenario is depicted in Fig.10c and d.Specificall,Fig.10c shows the presence of a twinned grain comprising a residual matrix (M2) and twins (T6),and in this grain,the area fraction of the twins is much larger than that of the matrix,unlike in the grain comprising M1and T1as depicted in Fig.10a.As a result,most of the boundary of the grain comprising M2and T6becomes the twin-twin interface between the twins in the grain (T6) and the twins formed in the surrounding grains (T7,T8,T9,and T10).In general,the strain energy accumulated in twins varies with the orientation of the parent grain and the extent of slip activation;therefore,a twin with lower stored strain energy can grow into adjacent twins with higher stored strain energy through the movement of grain boundaries constituting the twin-twin interface.During annealing,twins T6in Fig.10c grow significantl toward T8and T10but not toward T7(Fig.10d),which suggests that the stored strain energy in T6is lower than those in T8and T10but similar to or higher than that in T7.Consequently,after annealing,a coarse TOR is formed by the growth of T6,and the residual matrix (M2) remains unchanged because its stored strain energy is lower than that in T6and the interfaces between M2and T6are twin boundaries with low mobility.

Fig.10. Quasi in situ EBSD results showing (a,b) growth of residual matrix region during annealing and resultant formation of MOR and (c,d) growth of twinned region during annealing and resultant formation of TOR.The basal poles of the residual matrices (M1 and M2) and {10-12} twins (T1-T10) are depicted in the (0001) pole figure on the right side of the figure

Fig.11.(a) Variations in area fractions of grains before and after annealing as a function of angle between c-axis and RD (θ).(b,c) (0001) Pole figure before and after annealing for (b) region with 45° <θ ≤90° and (c) region with 0° ≤θ ≤45° The region with 0° ≤θ ≤45° corresponds to the twinned region before annealing and the TOR after annealing,and that with 45° <θ ≤90° corresponds to the residual matrix region before annealing and the MOR after annealing.

The preferential growth of the residual matrix region during annealing leads to a variation in the orientation distribution.Fig.11a shows the area fractions of grains before and after annealing as a function ofθ;the regions with 0° ≤θ≤45° correspond to the twinned region before annealing and the TOR after annealing,and those with 45°<θ≤90° correspond to the residual matrix region before annealing and the MOR after annealing.Comparison of the area fractions before and after annealing reveals that the fraction of the region with 0° ≤θ≤45° decreases whereas that of the region with 45°<θ≤90° increases upon annealing because the residual matrix region preferentially grows while consuming the twinned region during annealing.In theθrange corresponding to the twinned region and TOR (0° ≤θ≤45°),the fraction of the region with 15° ≤θ≤45° decreases after annealing whereas that of the region with 0° ≤θ <15° increases slightly after annealing.As the SF for basal slip of a grain gradually increases with an increase inθf(wàn)rom 0° to 45°[31],the twinned region with a higherθ(15° ≤θ≤45°) has higher stored strain energy than that with a lowerθ (0° ≤θ<15°).Therefore,the twinned region with 15° ≤θ≤45° is likely to be consumed by the surrounding twinned or residual matrix region with lower stored strain energy,whereas the twinned region with 0° ≤θ <15° is likely to grow during annealing owing to its low stored strain energy.This dependence of the growth behavior of the twinned region on its orientation causes a texture change after annealing.For the region with 45°<θ≤90°,the change in the maximum intensity of the (0001) pole figur after annealing is not large(from 9.4 m.r.d.to 10.3 m.r.d.),because of the increase in the overall area fraction in thisθrange (Fig.11b).However,in the twinned region,the RD-oriented region with 0° ≤θ<15° tends to grow whereas the off-RD region with 15° ≤θ≤45° tends to disappear during annealing;therefore,the basal texture of the region with 0° ≤θ≤45° becomes more concentrated along the RD and its maximum texture intensity increases significantl from 15.3 m.r.d.to 23.9 m.r.d.after annealing (Fig.11c).

Fig.12.Variation in overall texture of material with application of precompression and subsequent annealing treatment.

Fig.12 shows the variation in the overall texture of the material with the precompression and subsequent annealing treatment.The as-rolled material has a strong basal texture with the basal poles aligned almost parallel to the ND.When this material is compressed along the RD,the matrix region,havingθ >>45°,disappears almost completely,whereas the twinned region,having 0° ≤θ≤45°,develops owing to the formation of {10-12} twins.The twinned region governs the overall texture and the maximum texture intensity of the precompressed material owing to the small area fraction of the residual matrix region.Accordingly,the precompressed material has an RD-oriented twin texture with a maximum texture intensity of 12.1 m.r.d.After subsequent annealing at 250°C,although the texture of the region withθ45° intensifie slightly,the overall texture is still governed by the region with 0° ≤θ <45° owing to the large area fraction of the TOR.The basal poles of the TOR are more closely aligned with the RD than are those of the twinned region before annealing;therefore,the subsequently annealed material has a higher maximum texture intensity than the precompressed material.Consequently,the precompression and subsequent annealing treatment leads to a significan texture change of the material from an ND-oriented texture to an RD-oriented texture.

The mechanical properties of wrought Mg alloys at RT vary greatly with texture and loading direction because of the limited number of active slip systems at RT and the polar nature of deformation twinning [72-74].For example,in a rolled Mg alloy with a strong ND-oriented basal texture,accommodation of plastic deformation along the out-of-plane direction (i.e.,thickness direction) is difficul because both the plane and the direction of basal slip are parallel to the in-plane direction [75].Since {10-12} twinning induces significan lattice reorientation,many studies have attempted to modify texture using {10-12} twins for the purpose of improving the various physical properties of wrought Mg alloys,e.g.,strength [43,75-78],rollability [79,80],bendability [68],stretch formability [81],fatigue resistance[82,83],and damping capacity [84,85].However,the application of pre-deformation to induce the formation of {10-12}twins can cause a decrease in the ductility of the material owing to the accumulation of the imposed strain [68],and the stability of the microstructure can decrease owing to the high mobility of the {10-12} twin boundaries [86,87].Then,the process of subsequent annealing of {10-12}-twinned materials is vital for overcoming these drawbacks.Therefore,the present study is expected to expand the understanding of microstructural evolution mechanisms in pre-twinned Mg alloys during heat treatment and effectively contribute to an improvement in the various mechanical properties of wrought Mg products via texture control using {10-12} twinning.

4.Conclusions

This study investigates the microstructural evolution of a{10-12}-twin-containing precompressed AZ31 alloy during annealing at 250°C for 1h throughquasi in situEBSD analysis.The analysis results reveal that the partially twinned structure consisting of the twinned region and the residual matrix region changes to an almost twin-free structure comprising grown grains with serrated boundaries in the TOR and MOR after annealing.The annealing treatment leads to doubling of the average grain size from 19.2μm to 38.5μm and a decrease in the average KAM value—which represents the internal strain energy—from 1.27 to 0.69.Such microstructural variations are attributed mainly to grain growth by the migration of HAGBs through SIBM,and not by the movement of{10-12} twin boundaries.The twinned region formed at the early stage of precompression is more favorably oriented for basal slip than is the residual matrix region,and therefore,more dislocations are generated in the twinned region during precompression.The slip-favorable orientation of the twins and the dislocation transmutation reaction occurring in them lead to the accumulation of greater stored strain energy in the twinned region than in the residual matrix region.Therefore,to reduce the total strain energy accumulated in the material,the residual matrix region in a grain preferentially grows while consuming the twinned regions formed within adjacent grains during annealing.As a result,the area fraction of grains with matrix orientations increases whereas that of grains with twin orientations decreases after annealing.The variation in the matrix texture during annealing is insignificant but the twin texture intensity increases significantl and this texture becomes more concentrated along the RD because the RDoriented twins tend to grow during annealing owing to their fairly low stored strain energy.

Declaration of Competing Interest

The authors declare that they have no conflic of interest.

Acknowledgements

This work was supported by a grant from the National Research Foundation of Korea (NRF) funded by the Korean government (MSIP,South Korea) (No.2019R1A2C1085272).