Dislocation configuration evolution during extension twinning and its influence on precipitation behavior in AZ80 wrought magnesium alloy

Guo-Liang Shi ,Kui Zhang ,Xing-Gang Li ,Yong-Jun Li ,Ming-Long Ma ,Jia-Wei Yuan ,Hong-Ju Zhang

aState Key Laboratory of Nonferrous Metals and Process,Grinm Group Corporation Limited(GRINM),Beijing 100088,China

b GRIMAT Engineering Institute Co.,Ltd.,Beijing 101407,China

cGuobiao (Beijing) Testing and Certification.Co.,Ltd.,Beijing 101407,China

Abstract Thermomechanical treatment T10 (extension twinning+aging treatment) can largely enhance the precipitation strengthening effect of magnesium alloys.In this study,dislocation structure evolution and precipitation behavior during T10 treatment of an AZ80 extruded bar were analyzed mainly by two-beam diffraction in TEM.At a compressive strain of 1% in the extrusion direction (ED),a typical dislocation configuration,including basal I1 stacking faults (SFs) and dislocations,has been established in extension twins.As the strain reaches 7%,the volume fraction of extension twins increases to more than 90% at which point high dense I1 SFs and dislocations occur.After aging for 2 h at 150 °C for the 7% strained sample,masses of basal I1 SFs and dislocations remain in the extension twins and can act as effective nucleation sites and solute fast-diffusion channels for continuous precipitates.Consequently,the precipitates in extension twins become highly dense.

Keywords: Magnesium alloy;Extension twin;Dislocation;Stacking fault;Precipitation.

1.Introduction

Based on review analysis of over 3000 papers on magnesium and magnesium alloys indexed in SCI in 2020,the control of microstructures of Mg alloys to improve their mechanical properties through developing new alloys and advanced processing technology is still a main research focus[1].Extension twinning ({102}<101>twinning) is an important plastic deformation mechanism in magnesium alloys[2,3],and also beneficial to their recrystallization [4].Recent TEM studies show that a typical dislocation configuration including basalI1SFs anddislocations can be formed in extension twins via the interaction between advancing TBs and basaldislocations in the matrix [5–9],which can considerably amplify the age-hardening effect of Mg–Al alloys during artificial aging [10–12].On the basis of this finding,a new thermomechanical treatment T10 can be potentially developed for wrought Mg alloys[11];however,the evolution of dislocation configuration during extension twinning and its influence on the precipitation behavior of Mg17Al12continuous precipitates remain lacking in-depth TEM analysis,which impedes the precise control of mechanical properties.

When a single crystal of magnesium or its alloys is stretched in the c axis or is compressed perpendicularly to the c axis,extension twinning can readily be activated[3],and the critical resolved shear stress (CRSS) is only about 5–10 MPa[13],only slightly higher than the CRSS of basal slip.The CRSS of extension twinning in polycrystalline magnesium alloys increases;for instance,the CRSS in an AZ31B hot-rolled sheet is about 25–35 MPa[14],or 46.5 MPa in a recent report[15]and about 56 MPa in an extruded ZK60 bar [16].When the total compression strain perpendicular to the c axis or the total tensile strain along the c axis of a pure magnesium single crystal reaches about 5–6%,the single crystal completely transforms to an extension twin[7,13].Extension twinning includes three sequential processes:nucleation,propagation,and growth;moreover,it involves three growth directions: shear(forward) direction,twinning plane normal,and lateral direction=×;when an extension-twin thin crystal nucleates on the grain boundary,it first propagates alonganduntil it is blocked at the opposite grain boundary.It then starts to thicken along.All these processes include the formation and migration of TBs [17],which migrate by the slip of twinning dislocations [8].

Al is an important alloying element for magnesium alloys,because it can bring about higher strength and larger elongation at the same time [18].Thus,Mg–Al series alloys are currently the most widely used magnesium alloys,with Mg17Al12phases as their main aging hardening precipitates.This phase has two kinds of precipitation mechanisms:(1) discontinuous precipitation in which several discontinuous precipitates first nucleate at one grain boundary and then grow in a sheet shape,synchronously drives the grain boundary towards grain interior.A lamellar structure consisting of alternate layers of Mg17Al12andα-Mg is thus obtained.During the growth of precipitates,the grain boundary acts as a fast-diffusion channels of solute Al.However,the Al concentration abruptly changes from one side of the grain boundary to the other,hence the mechanism referred to as discontinuous precipitation.These sheet-shaped discontinuous precipitates cannot effectively hinder basal slip because their habit planes are parallel to the basal plane,and the spaces among them are always wide;(2) continuous precipitation: a large number of lath-shaped particles nucleate in the grain interior;the Al concentration at one point continuously decreases with aging time,so the mechanism involved is referred to as continuous precipitation.Continuous precipitates usually exhibit a high number density,so this mechanism can lead to a desirable hardening effect;moreover,dislocations can act as the preferred nucleation sites and the fast-diffusion channels of solute Al [19–24].Discontinuous precipitation consistently maintains its competitive advantage over continuous precipitation in conventional aging treatment;thus,Mg–Al series alloys often exhibit poor age-hardening effects [23].

Extension twins can largely enhance the age-hardening effect of Mg–Al series alloys,because they can not only impede discontinuous precipitation but also improve the number density of continuous precipitates [10–12,25].Extension twin boundaries can strongly hinder the migration of grain boundaries thereby limiting the growth of discontinuous precipitates.A typical dislocation configuration including basalI1SF anddislocations has been established in extension twins[5–7,11].These high-density dislocations are not only the preferred nucleation sites of continuous precipitates but also the fast-diffusion channels of the solute,resulting in a considerably higher nucleation rate and markedly shortened peak aging time.Meanwhile,extension twins are easily thickened under relatively low stress.Particularly for wrought magnesium alloys with texture,as the strain reaches 6%–8% under suitable loading for extension twinning,the volume fraction of extension twins can exceed 80% [11,14,26],indicating that most regions of the material contain high-density dislocations.Thus,the thermomechanical treatment T10 based on extension twins is a potential strengthening technology for magnesium alloys.

Recent studies have found that basaldislocations in the matrix can be transformed to sessile Frank partial dislocations ofI1SFs within the extension twins when TBs sweep the matrix during the thickening stage.The constriction of a pair of partial dislocations into a glissiledislocation is energetically beneficial,leading to a typical dislocation configuration in extension twins:a large number of wide basalI1SFs anddislocations [5–9].Another study suggests that these very broad non-equilibrium basal SFs within the extension twins of hexagonal close-packed (HCP) metals are associated with the migration of incoherent TBs [27].The precise application of T10 treatment to regulate precipitation behavior and mechanical properties requires an in-depth understanding of the evolution of dislocation configuration in extension twins with strain and how it influences the precipitation behavior;however,this subject is rarely reported.In the present study,EBSD and two-beam diffraction techniques in TEM have been used to reveal the evolution of dislocation configuration in extension twins during cold compression of an extruded AZ80 bar and its influence on the precipitation behavior of Mg17Al12continuous precipitates.This study lays a theoretical foundation for a new strengthening technology of magnesium alloys based on extension twinning.

2.Materials and methods

Specimens in the current study were cut from an extruded AZ80 bar with a cross-section of 230 mm × 140 mm and the following chemical composition:Mg-8.87Al–0.62Zn–0.15 Mn–0.27Ce–0.0034Fe–0.02Si–0.0012Cu (wt.%).This bar has three characteristic directions,i.e.ED (extrusion direction),TD (parallel to 230 mm sides of the rectangular cross-section),and ND (parallel to 140 mm sides of the rectangular cross-section).A semi-continuous casting ingot was homogenized at 395°C for 16 h,and the bar was produced by“multi-direction forging (MDF)+extrusion+online cooling.” MDF was conducted at temperatures exceeding 350 °C,and a graphite–oil mixture was used as a lubricant.After MDF,the forged ingot was directly extruded at 350 °C with an extrusion ratio of 10 and an extrusion speed of about 0.35 mm/s.To maintain a high degree of supersaturation and weaken grain coarsening,the bar was quickly quenched using an on-line water quenching device as soon as it escaped from the mold.The specimens with dimensions 14 mm(ND) × 16 mm (TD) × 30 mm (ED) were compressed on a universal testing machine at RT,with the compression parallel to ED;this loading path facilitates the activation of extension twinning;graphite flakes were used between the specimen and anvils as the lubricant.

The twinning characteristics in samples with 1 and 7%strains were analyzed by EBSD.Dislocation configuration and precipitation behavior in extension twins were observed by two-beam diffraction technique in TEM.EBSD was conducted on JSM-7001F field-emission SEM equipped with the EBSD system produced by TSL-EDAX.Ultimately,EBSD samples were electrolytically polished in 20% (volume fraction) nitric acid solution in alcohol at a voltage of about 15–20 V,with the temperature set below 0 °C.The Burgers vectorof different dislocations can be distinguished using the·=0 invisibility criterion,and<110>and<113>zone axes are normally chosen for HCP metals.Meanwhile,[0002],[10],[100],[101],[10],and [011]are the commonly usedvectors;under the diffraction condition of=[0002],dislocations are invisible but dislocations with acomponent can be seen [28].A strong fiber texture has been formed in an extruded AZ80 bar;thus,the (0001)basal planes of most grains are parallel to the ED [29,30];when the as-extruded samples are compressed in the ED,extension twinning can be activated,and the original lattice is rotated through 86.3 ° around<110>,rendering the (0001)basal planes in extension twins almost perpendicular to the ED [26,31].On the basis of this orientation transformation,the TEM foils were cut parallel to the ED,given the high probability that the electron beam is parallel to the<110>zone axes of the extension twins,which facilitates the analysis of dislocation configuration in extension twins via two-beam diffraction.TEM analysis was conducted on JEM-2100F and FEI-Tecnai-F20 with a 200 KV operating voltage,and TEM foils were first ground to a thickness of 50 μm.Disks measuring 3 mm in diameter were punched,and ion milling was ultimately conducted on Gatan PIPS691 for further thinning and perforation.

3.Results and discussions

3.1.Twinning characteristics

The engineering stress–strain curves of the as-extruded sample with a rectangular section during compression in the ED at RT have a typical S shape,as presented in Fig.1.Extension twinning is the dominant plastic mechanism in the early stages of compression;thus,the yield strength is relatively low at 160 MPa;when the strain reaches about 6%,non-basal slip ofdislocations and contraction twinning dominate the subsequent plastic strain [26],leading to rapid strain-hardening in the late stages,resulting in considerably high ultimate compressive strength (460 MPa).In the current study,plastic strains at 1 and 7% were chosen for microstructure observation in the present work (Fig.1).

Fig.1.Typical engineering stress–strain curve of as-extruded 14 mm(ND) × 16 mm (TD) × 30 mm (ED) sample during compression in the ED at RT (Graphite flakes were used between specimen and anvils as the lubricant;strain rate=0.0005 s-1).

The EBSD images in Fig.2 present twinning behaviors in three typical grains in a sample compressed to 1%strain,with the horizontal directions of all images parallel to compression in ED and the vertical directions parallel to ND.The twinning characteristics in MA grain are exhibited in Fig.2(a) and (b),and the inverse pole figure (IPF) in Fig.2(a) demonstrates the orientation of MA—that is,compression is applied parallel to its c axis,which hinders the activation of extension twinning but facilitates the activation of contraction twinning.Several parallel needle-shaped contraction twins are thus observed in Fig.2(b).The c axis of grain MB is approximately normal to compression.Thus,two extension twin variants are activated,as labeled in Fig.2(c).Different twin variants can form an“apparent crossing” twin structure,which results from twin–twin boundary formation [32];meanwhile,the adjacent twins belonging to the same variant can merge together,as shown in Fig.2(c) and (d).Fig.2(e) presents the kernel average misorientation (KAM) distribution map of MB grain,and a large number of low-angle boundaries with a misorientation angle of about 2° are found near TBs and around compound particles,suggesting a considerably high dislocation density in these regions.As shown in Fig.2(f),the c axis of grain MC is strictly perpendicular to ED,which significantly contributes to the activation of extension twinning,and only one variant is activated.Individuals of the same variant can also form an“apparent crossing”twin structure,and extension twin appears to bypass compounds and continue propagating,as shown in Fig.2(g).The KAM distribution map in Fig.2(h) also indicates a considerably high dislocation density near TBs and around compound particles.

The twinning behaviors in the 7% strained sample are shown in Fig.3.In the earlier study,the area fraction of the extension twins rises to 90% when the compression strain in the ED reaches 6%;meanwhile,the extrusion fiber texture is completely transformed[9].The IPF image in Fig.3(a)shows that most grains are completely transformed into extension twins,but several original structures remain in some grains,labelled by arrows in Fig.3(a).The original extrusion fiber texture is (0001)‖ED (c axes⊥ED).After the strain increases to 7%,the texture is fully transformed into (0001)⊥ED (c axes‖ED) by extension twinning,as shown in Fig.3(d).The KAM distribution map in Fig.3(c) shows that most areas present a misorientation angle about 2 °,particularly near the grain boundaries and TBs and around compounds,indicating the high density of dislocations in most regions in the material.

Fig.3.Twinning behaviors in the 7% strained sample,horizontal directions of all images ‖ ED,and vertical directions ‖ ND.(a): IPF image;(b): IQ image;(c): KAM image (blue: KAM=0 °,red: KAM=5 °);(d): (0001) pole figure (For interpretation of the references to color in this figure legend,the reader is referred to the web version of this article.).

3.2.Dislocation configurations in extension twins

The dislocation configurations in an twin T1 of the 1%strained sample are analyzed by two-beam diffraction in TEM,as shown in Fig.4,with the electron beam along the〈20〉 zone axes of the T1.Fig.4(a) and (b) present the bright-and dark-field images,respectively,under a normal diffraction condition.Tangled dislocations exist in the matrix MD and basalI1SFs in the T1 can stretch across the entire twin,and a larger version of SFs is shown in the inset of Fig.4(b).With reduced strain (1%) and small thin thickness(～1 μm),this extension twin is most likely in the early stages of growth;TBs migrate and continuously react with thedislocations in the matrix to produce sessile Frank partial dislocations ofI1SFs in the extension twin,and these SFs can continue stretching with the migration of TBs [6].With the orientation of the T1 adjusted to the two-beam diffraction condition=[0002],the bright-and dark-field images of the T1 are presented in Fig.4(c) and (d).BasalI1SFs anddislocations with low number density are observed in the T1,indicating that reactions of TBs withdislocations in the matrix are insufficient in the early stages of growth.Bothdislocations labeled in Fig.4(c) attach one of their ends to TB and the other end to SF.The morphology that the line directions ofdislocations closely aligned with the basal plane traces in the matrix has been observed in [6];thus,T1 is most likely an extension twin because the angle between the basal planes of T1 and MD is near 86 °,as shown in Fig.4(c).TBs in Fig.4(c) and (d) present as equal inclination fringes,indicating the larger inclination of TBs relative to the incident electron beam.The locally magnified image of the TB in Fig.4(d) shows that numerous parallel black lines cut off the equal inclination fringes and are most likely twinning dislocations (interfacial step) on the twin boundary.Bright-and dark-field images in Fig.4(e)–4(h)are under the=[010]diffraction condition,and the dark field-image can more clearly demonstrate the dislocation configuration than the bright-field image.Under this condition,basalI1SFs present them as fringes of equal inclination,and some SFs stretch across the entire twin,whereas other SFs are enclosed inside the twin.Thedislocation and SFs labeled in Fig.4(f) are the same as those labeled in Fig.4(d),which help present their three-dimensional structure and the possible reaction mechanism.The two SFs marked in Fig.4(f)may be formed during propagation of the extension twin and continue to become enlarged with the migration of the left TB.Subsequently,both SFs are enclosed inside the twin as soon as the second basal dislocations in the matrix arrive at the left TB;moreover,thedislocation with one end at the enclosed SF and the other end at the left TB is created by the constriction of a pair of partial dislocations.Similar configurations among TB,SFs,anddislocations can also be observed in this twin,as labeled in Fig.4(g) and (h).

Fig.4.Dislocation configuration in the 1% strained sample,with an electron beam along 〈20〉 zone axes of T1.(a) and (b): normal diffraction condition;(c) and (d): two-beam diffraction condition of the T1 with =[0002];(e)–(h): =[010]for the T1.

Fig.5 presents the dislocation configuration in extension twins T2 and T3 in the 1% strained sample with an electron beam along the<20>zone axes of the extension twin T3.The bright-field image under normal diffraction condition in Fig.5(a) shows that high-density dislocations exist in both the matrix and the extension twins.Moreover,the orientation relationship between the extension twin and the matrix can be determined by their diffraction patterns in the insets in Fig.5(a) and is labeled in Fig.5(a) by basal plane traces.The angle between these two traces is about 86 °,consisting of the orientation transformation of extension twinning.Dislocation projections in the matrix are mainly arrayed along the basal-plane traces.Linear defects along basal planes also appear in the extension twin and are most likely basalI1SFs.With the orientation of the T3 adjusted to the=[010]diffraction condition,bright-and dark-field images are shown in Fig.5(b)–(f) with (b)–(e) for the T2 and (f) for the T3.Dislocation contrasts between the T2 and the T3 completely vary,indicating a large misorientation between them.In addition,T2 and T3 may belong to different variants and connect,similar to the case in Fig.2(d).A long basalI1SF with a length of 750 nm is found in the T2,which is connected to the right TB;however,many short SFs with the contrast of equal inclination fringes are connected to the left TB,as marked in Fig.5(b)–(e),implying that the migration in the right TB is longer than that in the left one.The long SF is speculated to have only a small part remaining in the TEM film sample,which may be the reason it does not present equal inclination fringes.Numerous dislocations are present in the T2,most of which are connected to TB and are most likelydislocations,as marked in Fig.5(b)–(d).Under this diffraction condition,the dislocation contrast in the T3 is indistinct,only a fringe feature from SF can be distinguished,as marked in Fig.5(f).Orienting the T3 to the=[0002]diffraction condition helps observe thedislocations in the T3 becausecomponents are invisible,as shown in Fig.5(g)–(j).However,the dislocation contrast in the T2 is blurred,as can be seen in Fig.5(k)and(l).A typical dislocation configuration in the extension twin including long basalI1SFs and non-basaldislocations is exhibited in the bright-field image of the T3 in Fig.5(g).In most cases,eachdislocation is connected to SF on one end and to TB on the other end.An SF is first formed by a reaction between migrating TBs and onedislocation in the matrix,and continues to lengthen with the TB migration until anotherdislocation in the matrix is incorporated into the TB and then transform to a sessile Frank partial dislocation,which can enclose the SF.Adislocation linking TB and SF is then formed by the constriction of partial dislocations [6].Typical SFs anddislocations are labeled in Fig.5(g)–(j).Two styles of configurations between SF and thedislocation are often observed—that is,theZandYstyles,as labeled in Fig.5(g)–(j).A probable origin of the Z-style configuration is that another SF connected to thedislocation dipole is generated as the segment of the third basaldislocation in the matrix strike the TB on the line between the two nodes of the dislocation dipole.The most likely cause of the Y-style configuration is that two near-screwdislocations of the dipole attached to the same SF glide in different directions.As shown in Fig.5(h)–(j),TBs are marked by dotted lines,and the residual interfacial disconnections on TBs,which are connected to SFs ordislocations,can be straight lines (a,b,d,and e) or curves (c and f).Moreover,the straight lines can be parallel to basal-plane trace of the matrix (a and d) or the twin (b and e),implying that complex dislocation reactions occur at TBs.The connecting interface between twin variants T2 and T3 is shown in Fig.5(i),which present as equal inclination fringes.Thecomponents in the T2 in Fig.5(c)–(f) are invisible under the=[0002]diffraction condition of the T3.However,many dislocations with a partly invisible contrast can be observed in the T2,which is parallel to the basal-plane trace of the T2,and these dislocations may becomponents,as labeled in Fig.5(k).The long SF marked in Fig.5(b)–(d)is also observed in Fig.5(k),whereas short SFs connected to the left TB in Fig.5(e) no longer present as equal inclination fringes but as single lines,as labeled in Fig.5(l).

Fig.5.Dislocation configuration inside two connecting extension twin variants T2 and T3 in the 1% strained sample,with an electron beam along the <20>zone axes of the T3.(a): normal diffraction condition;(b)–(f): =[010]for the T3;(e)– (h): =[0002]for the T3.

Our earlier study indicates that the same AZ80 extruded bar reaches its limit of compression deformation in ED at RT when the strain increases to 7%,and the volume fraction of the extension twins exceeds 90%,as shown in Fig.3.Under this condition,the dislocation configuration in the extension twins is analyzed in Fig.6.In the bright-field image in Fig.6(a) under normal diffraction condition,five boundaries divide the field of view into six zones,and four zones with larger thicknesses are most likely extension twins because the same basal-plane trace can be determined in these zones,and they are defined as T4,T5,T6,and T7.Meanwhile,the basalplane trace in the narrow zone in the middle varies from that of the extension twins and thus is most likely the matrix,defined as ME.By deduction,these extension twins are originally four entities belonging to the same variant,similar to the scenario in Fig.2(f);moreover,they continue to thicken with the strain,and when the strain approaches the maximum T4,T5,and T6 are completely merged,whereas,T4 and T7,separated by a narrow ME,are about to merge.Three new extension twins are activated at the narrowest site of ME,as denoted by arrows.As indicated in Fig.6(b),these three twins nucleate at the TB between ME and T7;they traverse ME and T4 and even the merging interface between T4 and T5,indicating a markedly higher local stress.Under the=[0002]diffraction condition for the T4 twin,bright-and dark-field images are presented in Fig.6(b)–(h).Fig.6(b) is a low-magnification image with a field of view roughly identical to that in Fig.6(a).Meanwhile,Fig.6(c)–(h) are high-magnification images revealing three fields of view in Fig.6(a) and (b),as marked by rectangular frames in Fig.6(a) and (b).Contrary to those in Fig.6(a),most dislocations are invisible in Fig.6(b),except for retained dislocation contrast in the vicinity of the narrowest site of ME where T4 and T7 are about to merge;and a much higher dislocation density can be predicted.Consequently,a strong local stress concentration is obtained,which activates three new twins.Moreover,dislocation contrast and secondarily activated twins also appear near the left TB of T7.Compared with the invisible dislocation contrast in the bright-field image,darkfield images can clearly show the basalI1SFs anddislocations,and their number density grows exponentially,compared with that in the 1% strained sample as shown in Figs.4 and 5.Fig.6(c) and (d) present the bright-and darkfield images,respectively,of the same field of view located immediately above the white rectangle frame in (e) and (f)in Fig.6(b),a small portion of which is marked in Fig.6(a).In this field of view,three twins—T4,T5,and T6—merge.The bright-field image can only show high-density tangles of basalI1SFs anddislocations near the TBs,whereas the dark-field image can clearly exhibit dislocation configurations in almost all areas.In the dark-field image,T4,T5 and T6 show completely different configurations of basalI1SFs anddislocations;T4 contains a considerably wide,densely tangled area near the TB between T4 and ME.By deduction,this TB transforms to the merged TB between T4 and T7,and the other region of T4 is dominated by densely arrayed basalI1SFs;the dislocation density near the merged TBs (denoted by dotted lines in Fig.6(c)) is much higher than that in the interiors of T5 and T6.The top half of T5 is a wide,densely tangled area,whereas the lower half is dominated bydislocations;by contrast,the T6 interior is dominated by densely arrayed basalI1SFs,among which somedislocations exist.The two fields of view corresponding to Fig.6(e)–(h) are labeled by white rectangular frames in Fig.6(b),reflecting the dislocation configurations in the neighborhood of the narrowest sites of ME.In this area,T4 and T7 are about to merge,and high stress concentration induces secondary twinning.The dislocation density in these new twins is largely reduced relative to those in their surrounding areas.The dislocation density in T5 is lower than that in T4,and a marked difference in dislocation configuration in different areas of T5 is observed—that is,dislocations dominate the upper-left area,as shown in Fig.6(f),whereas basalI1SFs dominate the lower-right area,as shown in Fig.6(h).Dislocations in T7 are invisible in the brightfield images in Fig.6(e) and (g);by contrast,densely arrayed SFs anddislocations are visible in the dark-field images in Fig.6(f) and (h).The bright-and dark-field images in Fig.6(i)–(n) are under the=[010]diffraction condition for the T4 twin,exhibiting dislocation configurations in three fields of view,denoted by rectangular frames in Fig.6(b).In Fig.6(i) and (j),the same fields of view as those in Fig.6(g)and (h) are presented.In Fig.6(m) and (n),the field of view is denoted by a black rectangular frame in Fig.6(b).Under this diffraction condition,the dislocation-contrast distribution of ME is similar to that under the=[0002],as shown in Fig.6(b)—that is,the contrast in high-density dislocation tangles exists at the narrowest sites of ME,as shown in Fig.6(m) and (n);however,the contrast disappears in other sites of ME,as shown in Fig.6(i)–(l).T4 contains a considerably broad dislocation tangle zone;however,dislocation contrast is invisible around these secondary twins,as shown in Fig.6(i) and (m).This observation indicates that secondary twinning can lower the dislocation density,whereas highdensity dislocation contrast occurs in these secondary twins,compared with that under the=[0002].By deduction,these dislocations belong todislocations.The fringe contrast of basalI1SFs is observed in T4 and T5,as shown in Fig.6(i)–(l).

Fig.6.Dislocation configuration in the 7%strained sample,with the electron beam along the <20> zone axes of the T4 twin.(a): normal diffraction condition;(b)–(h): =[0002]for T4 twin;(i)–(n): =[010]for T4 twin.

The dislocation configurations in another extension twinned area of the 7% strained sample are shown in Fig.7 with the electron beam along the〈20〉zone axes of twinned area.The two-beam diffraction condition in Fig.7(a)–(d) is=[010]for the twinned area.Fig.7(a) demonstrates the macrograph of this twinned area,which consists of three merged twins labeled as T8,T9,and T10.The dislocation density in T8 is not uniform–that is,characterized by a lowdensity area in the upper part and a high-density area in the lower part.A magnified image of the top–left corner of Fig.7(a) is shown in Fig.7(b),and the merged TBs and the inhomogeneity of the dislocation density in T8 are clearly presented.The fields of view of the two dark-field images in Fig.7(c) and (d) are denoted by black rectangular frames in Fig.7(a).High density dislocation exists in both fields of view,and the dislocation density in the lower part of T8(Fig.7(d)) is higher than that in the upper part of T9(Fig.7(c)).Moreover,Fig.7(c) and (f) are two dark-field images of almost the same field of view under differentconditions,exhibiting completely different dislocation contrasts.By deduction,the dislocation contrast in Fig.7(a)–(d) originates partly from thecomponents.Fig.7(e)–(h) are under the=[0002]condition of the twinned area.Fig.7(e) clearly shows that the twinned area is consists of three merged twins,and the T10 area is enlarged upward,which may be attributed to a slight change in orientation on the edge of the ion thinning hole.The fields of view corresponding to Fig.7(f)–(h)are labeled with black rectangular frames as that in Fig.7(e).Fig.7(f) shows the dislocation configuration in the upper part of T9,with almost the same field of view as that in Fig.7(c),which includes densely arrayed basalI1SFs anddislocations.The higher the dislocation density,the closer it is toward the TB,and the TB between T9 and the matrix MF present as fringes of equal inclination.Fig.7(g) and(h) present the dislocation configuration in T8.Fig.7(h) is a magnified image of part of Fig.7(g).The TB between T8 and the matrix MF does not present as fringes of equal inclination but as a wide dislocation tangle zone,implying a considerably high intensity of reaction between TB anddislocations in matrix MF,relative to the reaction intensity at the TB between T9 and the matrix MF in Fig.7(f).Moreover,the merged TB between T8 and T9 is also composed of high-density dislocation tangles,and the low-density area in the upper part of T8 in Fig.7(g) and (h) also consists of densely arrayed basalI1SFs anddislocations,similar to those in T9,as shown in Fig.7(f).

Fig.7.Dislocation configuration in extension twins in the 7% strained sample,with the electron beam along the 〈20〉 zone axes of the T9 twin.(a)–(d): =[010]for the T9 twin;(e)–(h): =[0002]for the T9 twin.

The following rules can be determined from the aforementioned analysis: (a) under strain conditions,several extension twins belonging to the same variant usually form simultaneously in one grain and ultimately merge,generating merged TBs that consist of high-density dislocation tangles;(b) during thickening of an extension twin,TBs migrate towards the matrix and react withdislocations in the matrix simultaneously,resulting in a typical dislocation configuration consisting of basalI1SFs anddislocations.The spaces among basalI1SFs increasingly decrease,and the density ofdislocations among SFs progressively increases,leading to severe dislocation tangles.Thus,individual dislocations or SF cannot be distinguished by two-beam diffraction in TEM.How does this dislocation configuration in extension twins influence the precipitation behavior? In the following sections,pre-compressed samples with 6–7% strain undergo artificial aging at 150 °C for 2 and 50 h.The dislocation configuration and precipitation behavior are analyzed by two-beam diffraction in TEM.

3.3.Precipitation behavior in extension twins

Fig.8 presents the dislocation configuration and precipitation behavior in the extension twins after treatment consisting of pre-compression with 7% strain+artificial aging at 150 °C for 2 h,with the electron beam along [20]of the T12 twin.The age-hardening curves at 150 °C for 50 h of the present AZ80 extruded bar with and without pre-deformation have been reported in the literature [11],from which the following conclusions can be drawn: (a) Cold compression with 6% strain in the ED increases the hardness from 73 HB (asextruded) to 90 HB by work hardening (increasing the dislocation density).After the sample undergoes aging for 2.5 h,the hardness is reduced to 86 HB by recovery (decreasing the dislocation density).The hardness increases rapidly to the peak value of 93 HB at 20 h and then remains stable until 94 HB at 50 h.(b) The age-hardening curve without predeformation has an incubation period of 15 h with hardness lower than 75 HB.The hardness then rises slowly but does not peak until the duration of 50 h(86 HB)is reached.Consequently,sufficient extension twinning can completely change the precipitation behavior,not only increases the peak hardness value,but also markedly reduces the time to peak.Moreover,after the sample undergoes aging for 2 h,the recovery stage is basically completed,and the dislocation density is reduced.Simultaneously,the nucleation stage of the precipitates basically ends,whereas the stage of fast growth begins;meanwhile the hardness rises sharply.Three extension twins with the same orientation are shown in Fig.8(a)—T11,T12,and T13.Fig.8(b)–(k) present the bright-and dark-field images from five fields of view.The two-beam diffraction condition in Fig.8(a)–(i) is=[0002]for the T12 twin.As shown in the figure,I1SFs anddislocations,still have a high number density,despite undergoing recovery during aging at 150 °C for 2 h.Moreover,continuous precipitates nucleated on dislocations are speculated to hinder the dislocation annihilation.Meanwhile,continuous precipitates in extension twins have a very high number density.By contrast,precipitation in the matrix is difficult to visualize,implying that the typical dislocation configuration can largely enhance the nucleation and growth of continuous precipitates.Along [20],most continuous precipitates in extension twins present themselves as lath-shaped or rectangular,with the pair of surfaces parallel to the (0001) basal plane as the length,and the other pair of surfaces as the thickness.The length,thickness,and lengthto-thickness ratio are three important features for precipitation strengthening.Determined by statistical analysis,the average values of length,thickness,and length-to-thickness ratio of continuous precipitates in these three twins are 49 nm,23 nm and 3,respectively.In areas dominated by SFs,such as zones near the right lower TB of T12 and left upper TB of T13,precipitates referred to as long-thin precipitates have large length-to-thickness ratios,as shown in Fig.8(b),(c),(f),and(g).However,in areas dominated bydislocations,the length-to-thickness ratio is reduced,and such precipitates are referred to as short-thick precipitates.Some particles are even thicker and present themselves as giant blocks,as shown in Fig.8(d) and (h).Moreover,large irregular blocks tend to precipitate on TBs mainly connected todislocations,such as those on the right lower TB of T11 and left upper TB of T12,as shown in Fig.8(e) and (i).However,no large irregular blocks have precipitated on TBs mainly connected toI1SFs,such as those on the right lower TB of T12 and left upper TB of T13,as shown in Fig.8(g).The two-beam diffraction condition=[010]of the T12 twin in Fig.8(j)and (k) can reveal the fringes of equal inclinationfromI1SFs but cannot clearly show the precipitates.

Fig.9 presents the dislocation configuration and precipitation behavior in one grain MG when the aging time is extended to 50 h at 150 °C (pre-deformation strain≈6%),under the two-beam diffraction condition=[0002]for the T14 twin(electron beam‖[20]of the T14 twin).This grain MG has five extension twins—T14,T15,T16,T17,and T18—with the same orientation.The number density of the lathshaped precipitates is markedly higher in T14,T17 and T18 with large thickness because basalI1SFs anddislocations can fully develop during the long-distance migration of TBs;by contrast,few lath-shaped precipitates are formed in thin T15 and T16 because the typical dislocation configuration cannot be fully developed;however,coarse short-thick precipitates and giant irregular blocks are formed on TBs.Fig.9(c) and (d),are respectively the bright-and dark-field images of the same field of view in T18 in which basalI1SFs,dislocations,and short–thick and long–thin precipitates coexist.As determined by statistical analysis,the average length,thickness,and length-to-thickness ratio of the continuous precipitates in T18 are 117 nm,42 nm,and 4,respectively;moreover,both length and thickness are twice those of the precipitates in the sample aged for 2 h,whereas the length-to-thickness ratio only slightly increases.Meanwhile,a large number of long–thin continuous precipitates precipitate in the matrix MG.

Fig.10 presents the dislocation configuration and precipitation behavior of two extension twins—T19 and T20—in the same sample treated by pre-compression with 6% strain+artificial aging at 150 °C for 50 h.The two-beam diffraction condition of=[0002]for extension twins is used under the incident way of the electron beam ‖ [20]of the extension twins.Fig.10(a) and (b) correspond to the T19 twin,and (c)–(f) correspond to the T20 twin.The number density of continuous precipitates in the T20 is markedly higher than that in the T19 because the T20 has a higher dislocation density.In the T20,most giant short–thick laths are connected to severaldislocations simultaneously,indicating that thedislocation has an intense strain field and exhibits high solute transport,facilitating the nucleation and growth of precipitates.By statistics analysis,The average length,thickness,and length-to-thickness ratio of the continuous precipitates in the T20 are statistically calculated to be 125 nm,40 nm and 4,respectively,which are close to the statistical results in the T18 twin in Fig.9.

For the engineering applications of the AZ80 wrought magnesium alloy with a larger cross-section,the preferred cold pre-deformation strain is about 6% in the ED or rolling direction which can induce a high volume fraction of extension twins.The average length,thickness,and length-tothickness ratio of the lath-shaped continuous precipitates are about 120 nm,40 nm and 4,respectively,after peak aging treatment at 150 °C for 20～50 h.

Scanning-transmission electron microscopy (STEM) images of continuous precipitates in the sample treated by precompression with 6% strain+artificial aging at 150 °C for 50 h are shown in Fig.11.The incident beam in Fig.11(a)–(c) is nearly parallel to the [0001]direction of the extension twins and shows largely diverse precipitation.Most continuous precipitates in Fig.11(a) present themselves as laths with six variants because of their Burgers orientation relationship with the matrix [21].In addition to laths,precipitates with equiaxial and needle-like morphologies are observed in Fig.11(b) and (c),dominated by only one variant.The incident beam in Fig.11(d) is nearly perpendicular to the [0001]of the extension twin,and the field of view is dominated by long–thin precipitates,in addition to few rod-shaped precipitates,which are perpendicular or nearly perpendicular to the basal plane.Such diversity in precipitation in extension twins is potentially caused by the diverse and complex dislocation configurations.

Fig.11.Scanning–transmission electron microscopy images of continuous precipitates in the sample treated by pre-compression with 6% strain+artificial aging at 150 °C for 50 h.

4.Discussions

In recent years,numerous studies have confirmed that extension twinning significantly enhances the age-hardening effect of wrought Mg–Al alloys [10–12,25],given that nanoscale particles continuously precipitated in extension twins gain a considerably higher number density than those in matrices and that TBs can hinder the growth of coarse discontinuous precipitates.On the basis of these findings,a new thermomechanical treatment referred to as T10 can be potentially developed for wrought Mg alloys.The characteristics,evolution,and formation mode of the dislocation substructure in extension twins—in addition to its influence on precipitation—should be elucidated to precisely control the strengthening effect of T10 treatment,which is the objective of the present study.

The dislocation substructure in {102} twins is formed by the interactions between matrix dislocations and TB,and studies on this subject can be traced back to 1950 s.One of the most representative early studies was conducted by M.Yoo and C.Wei concerning the growth of the {102} deformation twin in Zn.They proposed one important interfacial reaction on the {102} TB by correspondence matrices—that is,two [a1,3]dislocations in the matrix can be transformed to one [c±a2]dislocation on the prismatic plane of the twin and one double twinning dislocation on the TB [8].A large volume of TEM studies have recently revealed that a typical dislocation configuration is formed in {102} twins of HCP metals consisting of basalI1SFs anddislocations[5–7,11,27].F.L.Wang et.al have conducted continuous and in-depth research on the subject,and a physical mechanism underlying the formation and evolution of the dislocation substructure in the {102} twin has been proposed [5–7].In the present study,TEM images of the dislocation structure in the 1% strained sample are consistent with the physical mechanism proposed by F.L.Wang et.al.This study also analyzes the formation of Z-style and Y-style configurations,which commonly occur in the {10–12} twin.This study further reveals the dislocation structure in the 7% strained sample and its influence on precipitation behavior during aging at 150°C.

On the basis of findings reported in previous researches and the current study,a schematic of the complete microstructure evolution during the T10 treatment of wrought Mg alloys is presented in Fig.12.Fig.12(a) shows an original grain in the wrought material,which contains a large number of pre-existing basaldislocations (blue short lines in the matrix).High-density dislocations can be retained within the matrices by on-line quenching after hot deformation.The c axis of this grain is perpendicular to the compression direction(indicated by arrows in Fig.12(a)),which contributes to the activation of {102} extension twinning.After a nucleus of the extension twin is formed at the grain boundary,it rapidly propagates the whole grain and then starts to grow by the glide of twinning dislocations on TB.During twin growth,a typical dislocation configuration is continuously formed via the interaction between the migrating TB and basaldislocations in the matrix,as shown in Fig.12(b).

Fig.12.Dislocation structure evolution and precipitation behavior during the thermomechanical treatment T10: blue short lines in the matrix represent dislocations on basal planes;yellow lines denote TBs;short green lines in the twin indicate basal I1 SFs;red curves in the twin are dislocations;black rectangular blocks in (f) represent precipitates (For interpretation of the references to color in this figure legend,the reader is referred to the web version of this article.).

The fundamental physical mechanism [6]can be described as follows: (1)Formation of basal I1 SF.As one basaldislocation segment in the matrix is incorporated by the extension twin,it is first transformed into a 1/6<203>-type sessile Frank partial dislocation on the TB (yellow line);anI1SF (denoted by short green lines) is then emitted on the basal plane of the twin between the sessile Frank partial dislocation and the migrating TB,the boundaries of which include sessile Frank partials on the basal plane and retained twinning dislocation on the TB.The width of the basalI1SF increases with the TB migration.(2)Enclosed basal I1 SF.The emitted basalI1SF cannot be enlarged infinitely and can be enclosed when the second basaldislocation segment in the matrix strikes the TB only on the retained twinning dislocation of the SF.(3)Formation of adislocation dipole(red lines).After the basalI1SF is enclosed,two interfacial dislocation segments from the first and seconddislocations can be constricted,resulting in a pair ofdislocations between the sessile SF and the advancing TB,which have the same Burgers vector but opposite line directions—that is,they constitute a dislocation dipole.Since near-screwdislocation dipoles are glissile,so self-annihilation can occur in some scenarios.(4)Z-style configuration.When the third basaldislocation segment in the matrix strikes the TB just on the line between the two nodes of thedislocation dipole,another SF connected to the dislocation dipole is generated;thus a Z-style configuration appears.(5)Y-style configuration.The two near-screwdislocations of the dislocation dipole attached to the same SF glide in different directions.(6)Separation between the dislocation dipole and the SF.After thedislocation dipole is largely extended by the advancing TB,it can be separated from the SF because of the attractive force between the twodislocations;the final dislocation configuration consists of ahalf-loop attached to the TB and an independent double-looped basalI1SF.

The volume fraction of the extension twin increases with an increase in the local strain,as shown in Fig.12(c).As the local strain reaches about 6%,the extension twin can occupy the whole grain,as shown in Fig.12(d).Currently,extension twinning ceases,and the whole grain gains a completely different dislocation configuration.When local strain further increases,the dominant plastic mechanism changes to the non-basal slip ofdislocations and contraction twinning.The number density ofdislocations rapidly increases to form dislocation tangles,as shown in Fig.12(e).Under higher local strain,two neighbouring twins in the same grain can merge only when they belong to the same variant,and dislocation tangles exist in the vicinity of merged TBs.During subsequent aging,these high-density dislocations can act as effective nucleation sites and fast diffusion channels for solutes,thereby obtaining nanoscale precipitates with a high number density.Thus,{102} twinning can significantly enhance the age-hardening effect of Mg alloys.

5.Conclusions

In this study,the microstructural evolution of an AZ80 extruded bar during T10 treatment (extension twinning+aging treatment) was analyzed.Extension twinning was induced by cold compression in the ED and characterized by EBSD.Dislocation configurations and precipitation behaviors were analyzed by two-beam diffraction in TEM.The conclusions are thus summarized:

(1) When as-extruded polycrystal samples are coldcompressed in the ED,extension twinning can be readily activated.In the majority of cases,several extension twins belonging to the same variant simultaneously nucleate in one grain and can merge with an increase in strain.The volume fraction of extension twins exceeds 90% when compression strain reaches its limit value of～7%.

(2) Under 1% compression strain,a typical dislocation configuration,including basalI1SFs anddislocations,is established in extension twins.In the majority of cases,adislocation connects SF on one end and TB on the other end;Z-and Y-style configurations can also be observed.The Z-style configuration may be attributed to the generation of another SF connected to thedislocation dipole as the segment of the third basaldislocation in the matrix strikes the TB only along the line between the two nodes of the dislocation dipole;meanwhile,the Y-style configuration may be attributed to the two near-screwdislocations of the dislocation dipole attached to the same SF gliding in different directions.By the time the strain reaches its limit value of 7%,extension twins are fully thickened and merged in most areas.BasalI1SFs anddislocations are intensively tangled in extension twins,and dislocation tangles appear near merged TBs.

(3) After the 7% strained samples undergo aging at 150 °C for 2 h,masses of basalI1SFs anddislocations remain in extension twins and can serve as effective nucleation sites and solute fast-diffusion channels for continuous precipitates,resulting in a high number density of precipitates in extension twins.These continuous precipitates mostly appear as long-thin and shortthick laths.Most short-thick laths are associated withdislocations,implying that thedislocation has an intense strain field and exhibits high solute transport.In the[20]direction of extension twins,the average length,thickness,and length-to-thickness ratio of lath-shaped continuous precipitates are about 49 nm,23 nm and 3,respectively.After aging for 50 h,abundant basalI1SFs anddislocations remain pinned by lath-shaped precipitates,and the average length,thickness,and length-to-thickness ratio of the lath-shaped precipitates increase to about 120 nm,40 nm and 4,respectively.

Declaration of competing interest

There is no conflict of interest.

Acknowledgments

This study is supported by Beijing Natural Science Foundation (No.2194090).We thank Ting Li and Hai-Yan Yu for their careful operation of the TEM instrument during the twobeam diffraction analysis of the dislocation configuration in extension twins and Xiao-Lei Han for his expertise in EBSD operation.