Effects of Al addition on microstructure and mechanical properties of extruded Mg–3Bi alloy

Sng-Cheol Jin,Je Won Ch,Jun Ho Be,Hui Yu,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

cSchool of Materials Science and Engineering,Hebei University of Technology,Tianjin 300130,China

Abstract Effects of Al addition to a Mg–Bi binary alloy on its microstructural characteristics and tensile properties after extrusion are investigated via extrusion of Mg–3Bi–xAl(x=0,1,and 2 wt%)billets and analysis of the extruded materials.The Al addition negligibly affects the second-phase particles of the extruded alloy;however,an increase in the Al content causes significant decreases in the average grain size and maximum texture intensity of the extruded material owing to an increase in the area fraction of dynamically recrystallized(DRXed)grains.The Al addition improves the strength of the extruded alloy;this improvement is attributed to the enhanced grain-boundary hardening and solid-solution hardening effects induced by grain refinement and Al solute atoms,respectively.As the Al content increases from 0 wt% to 1 wt% and 2 wt%,the tensile elongation increases substantially from 2.8% to 9.4% and 16.9%,respectively.The reduction in the number and size of unDRXed grains with increasing Al content suppresses the formation and coalescence of cracks in the unDRXed grains during tension,which results in a significant improvement in the tensile ductility of the extruded material.During tensile deformation,large undesirable twins that act as crack initiation sites are locally formed in the unDRXed grains of the Mg–3Bi alloy,whereas relatively smaller twins are uniformly formed in both the DRXed and the unDRXed grains of the Mg–3Bi–2Al alloy.Consequently,the extruded Mg–3Bi–2Al alloy has a substantially higher tensile yield strength–elongation product(2924 MPa%)than the extruded Al-free B3 alloy(381 MPa%).

Keywords:Mg–Bi–Al alloy;Extrusion;Microstructure;Strength;Ductility.

1.Introduction

The rapid development of devices aimed at improving the convenience of automobile drivers has been accompanied by an increase in the use of various electronic and electrical components in addition to existing fundamental devices in automobiles.However,such use of additional convenienceenhancing devices inevitably leads to an increase in the total weight of an automobile,which,in turn,causes a decrease its fuel efficiency and an increase in its carbon dioxide emissions[1–3].The intensification of global regulations on fuel efficiency and carbon dioxide emissions of automobiles has resulted in Mg alloys receiving great attention in the automobile industry because of their low density and high specific strength.Cast Mg alloys have mostly been used as automobile components owing to the ease of manufacture and good castability of commercial Mg–Al-based alloys[4–6].In recent years,extensive research has been conducted on extruded Mg alloys—which have much higher mechanical properties than cast Mg alloys—for expanding the application of Mg alloys to automobile body and chassis components,which are required to have high strength and high stiffness.

Since a relatively high strain is imposed in a billet and large friction is generated between the billet and the container walls during extrusion,a large amount of heat is produced by deformation and friction[7–9].In commercial highalloyed Mg–Al-and Mg–Zn-based alloys(e.g.,AZ80 and ZK60),thermally unstable phases with low melting temperatures(432°C and 413°C for Mg17Al12and MgZn2,respectively)are formed during hot extrusion[10–12].An increase in the amount of deformation and friction heat with increasing extrusion speed can lead to hot cracking of the abovementioned commercial alloys during their high-speed extrusion owing to the local melting of the thermally unstable phases[4,12–14].Consequently,these commercial alloys have poor extrudability and their extrusion at low speeds translates into a low production rate of extruded products.Mg–Bibased alloys with high mechanical properties and excellent extrudability have been recently developed for overcoming this shortcoming of these commercial alloys.Meng et al.[15]reported that an extruded Mg–8Bi–1Al–1Zn(wt%)alloy has a good strength–elongation balance with a tensile yield strength(TYS)of 291MPa and a fracture elongation of 14.6%,where this balance is attributable to a completely recrystallized grain structure and uniform distribution of numerous Mg3Bi2precipitates.Remennik et al.[16]developed extruded Mg–5Bi–1Ca and Mg–5Bi–1Si(wt%)alloys with high tensile elongations of>40% at room temperature(RT)through a combination of rapid solidification and subsequent hot extrusion processes.This high ductility is a result of an ultrafine microstructure with grain sizes smaller than 2μm,which facilitates the activation of grain-boundary sliding(GBS)during deformation.In addition,Go et al.[17]demonstrated that a Mg–5Bi–3Al(wt%)alloy can be successfully extruded at a very high extrusion speed of 67m/min without any cracking.The extrudability of this alloy is noticeably excellent given that the maximum extrusion speed of the commercial AZ80 alloy is～5m/min[12,17].This extraordinary extrudability is attributed to the very high melting temperature(823°C)of the Mg3Bi2phase formed in Mg–Bi alloys,which is about twice those of the Mg17Al12and MgZn2phases.Somekawa and Singh[18]reported that Mg–Bi binary alloys extruded at low temperatures of 105–210°C have high RT tensile elongations of>50%,which is attributed mainly to the GBS activation by the fine-grained microstructure and equilibrium grain boundary structure.However,when Mg–6Bi and Mg–9Bi(wt%)binary alloys are extruded at 350°C–which belongs to the normal extrusion temperature range(～300–400°C)of Mg alloys,the extruded alloys have relatively low TYSs(129 and 141MPa,respectively)and low elongations(4.1% and 5.1%,respectively)owing to the formation of a partially recrystallized grain structure containing coarse unrecrystallized grains[19].The results of previous studies on Mg–Bi-based ternary alloys signify that the addition of Al,Ca,or Si can improve the mechanical properties of the extruded Mg–Bi binary alloys[16,17,20].However,these previous studies examined high-alloyed Mg–Bi-based alloys with large total amounts of alloying elements(6–10 wt%),whose material cost is high.The present study is aimed at developing low-cost,high-performance Mg–Bi-based alloys;to this end,a Mg–3Bi(wt%)alloy is selected as the base material and 1 wt% or 2 wt% Al is added to the binary alloy as an additive alloying element for improving the mechanical properties of the binary alloy.The variations in the microstructural characteristics and tensile properties of the extruded material caused by the Al addition are systematically investigated with a particular focus on the significant improvement in its tensile ductility.

2.Experimental procedure

Mg–3Bi,Mg–3Bi–1Al,and Mg–3Bi–2Al(wt%)alloys were used in this study;for conciseness and simplicity,they are hereafter referred to as B3,BA31,and BA32 alloys,respectively.Cast billets for extrusion were prepared by the conventional mold casting technique in a CO2–SF6gas mixture according to a previously reported procedure[20].The cast billets were homogenized in an inert gas atmosphere containing an Ar–SF6mixture in an electric furnace at 390°C for 24h.Samples for extrusion were obtained by machining the homogenized billets into a cylindrical shape with a diameter of 68mm and length of 120mm.The resultant cylindrical samples were preheated at 350°C for 1h in a resistance furnace and then subjected to direct extrusion under the following conditions:an extrusion temperature of 350°C,an extrusion ratio of 10,and a ram speed of 1mm/s.Extruded bars with a diameter of 21.5mm were fabricated using a 300-ton horizontal extrusion machine and a flat-faced extrusion die with a central hole.

The microstructural characteristics of the extruded alloys were analyzed on the cross-sectional plane along the extrusion direction(ED)by optical microscopy(OM),field-emission scanning electron microscopy(FE-SEM),X-ray diffraction(XRD)spectroscopy,and electron backscatter diffraction(EBSD).All specimens for the microstructural analyses were mechanically polished with progressively finer grades of emery paper(from #120 to #2000)and then polished using 1μm diamond paste.Specimens for EBSD examination were additionally polished with a colloidal silica solution(0.04μm)for 40min to remove surface strains and obtain reliable crystallographic data.The polished specimens for OM and FESEM examinations were carefully etched in an acetic picral solution(10ml acetic acid+3.0g picric acid+10ml distilled water+100ml ethanol(99.5%)).

The area fractions of the dynamically recrystallized(DRXed)grains and unDRXed grains of each extruded alloy were measured over a relatively large area(15.9 mm2)in the optical micrographs via calculation of the manually identified DRXed and unDRXed regions using ImageJ software.For tensile testing,dog-bone-shaped specimens with gauge dimensions of 6mm(diameter)×25mm(length)were machined from the extruded alloys;the tensile loading direction corresponded to the ED.Tensile tests were performed using a Shimadzu AGS-100kNX universal testing machine at RT(23°C)and a strain rate of 1×10-3s-1.For analysis of the fracture mechanisms of the extruded alloys,the microstructures on longitudinal cross-sections of the fractured tensile specimens were observed by OM and EBSD and the fracture surfaces of these specimens were observed by FE-SEM.

Fig.1.Optical micrographs of extruded alloys:(a)B3,(b)BA31,and(c)BA32.fDRX denotes the area fraction of dynamically recrystallized(DRXed)grains.ED and unDRXed denote the extrusion direction and unrecrystallized grains,respectively.

FE-SEM measurements were conducted in secondary electron mode at an accelerating voltage of 15kV.All EBSD measurements were carried out at an accelerating voltage of 15kV with a working distance of 20±3mm using a TSLTMEBSD camera installed in a FE-SEM(Hitachi,SU-70).Automated EBSD scans were performed in stage control mode with step sizes of 1.6μm for the extruded specimens and 1.0μm for the fractured tensile specimens using the TexSEM Laboratories data acquisition software.The resulting EBSD data was analyzed using the TexSEM Laboratories orientation imaging microscopy analysis 7.0 software after one-step cleaning using grain dilation(clean-up parameters:grain tolerance angle=5° and minimum grain size=2μm).Only reliable EBSD data with confidence indexes greater than 0.1 were used to analyze the average grain size,texture,Schmid factor(SF),kernel average misorientation(KAM),and twin boundaries.XRD measurements were performed using Cu Kαradiation at a scan speed of 2°/min in the range of 20°–80°.

3.Results and discussion

3.1.Variation in microstructure of extruded materials with Al addition

Fig.1 shows the optical micrographs of the extruded alloys,which reveals that all three extruded alloys exhibit a bimodal grain structure consisting of equiaxed DRXed grains and elongated unDRXed grains.As the Al content increases from 0 wt% to 1 wt% and 2 wt%,the area fraction of the DRXed grains(i.e.,the DRX fraction)increases from 46%to 68% and 78%,respectively.Moreover,the size of the unDRXed grains also decreases with increasing Al content.These results indicate that Al addition to the Mg–3Bi alloy promotes DRX behavior during hot extrusion,which,in turn,results in a more homogeneous grain structure of the extruded material.The individual characteristics of the DRXed and unDRXed grains are analyzed via generation of the inverse pole figure maps and ED inverse pole figures of the total(DRXed+unDRXed),DRXed,and unDRXed regions of the extruded alloys,as shown in Figs.2 and 3.With an increase in the Al content,the average grain size of the extruded material gradually decreases(153.3,68.9,and 41.7μm for B3,BA31,and BA32,respectively)and the maximum texture intensity also decreases(13.8,9.5,and 4.6 for B3,BA31,and BA32,respectively)(Fig.2).In addition,the average width of the unDRXed grains decreases significantly from 79.8μm for the B3 alloy to 23.5μm for the BA32 alloy,whereas the difference in the average sizes of the DRXed grains of the three alloys is relatively insignificant(24.1–28.6μm)(Fig.3).Therefore,the gradual reduction in the average grain size of the extruded alloy with increasing Al content is attributed to the reduction in both the area fraction and the size of the unDRXed grains,rather than to the refinement of the DRXed grains(Fig.4).In extruded Mg alloys that do not contain any rare-earth elements or Ca,unDRXed grains generally have a highly intense<10-10>texture owing to the lattice rotation caused by the continuous activation of basal and prismatic slips in the grains during hot extrusion[21,22].In the present study,the unDRXed grains of all the extruded alloys have a strong<10-10>texture and their maximum texture intensities(17.0–23.4)are considerably higher than those of the DRXed grains newly formed during extrusion(2.9–3.2)(Fig.3).Given that the texture distribution and intensity of the DRXed grains are similar in the three extruded alloys,the gradual texture weakening of the extruded alloy with increasing Al content results from the reduction in the area fraction of the unDRXed grains,which have a strong texture.

Fig.2.Inverse pole figure maps and ED inverse pole figures of extruded(a)B3,(b)BA31,and(c)BA32 alloys.davg denotes the average grain size.

Fig.3.Inverse pole figure maps and ED inverse pole figures of unDRXed and DRXed regions of extruded(a)B3,(b)BA31,and(c)BA32 alloys.wunDRX and dDRX denote the width of unDRXed grains and the average size of DRXed grains,respectively.

Fig.4.Variations in average grain size of extruded alloy,average size of DRXed grains,and average width and area fraction of unDRXed grains with Al content.

SEM micrographs of the extruded alloys are shown in Fig.5a–c,which reveal that numerous fine(0.5–3.0μm in size)second-phase particles are uniformly distributed in all the extruded alloys.The XRD patterns in Fig.5d demonstrate that all the particles formed in the extruded alloys are of the Mg3Bi2phase and that no Mg17Al12particles are present even in the Al-containing BA31 and BA32 alloys.Fig.5e shows the equilibrium phase diagram for Mg–3Bi–xAl(x=0–4 wt%),as calculated using FactSage software.From this diagram,it can be seen that at the extrusion temperature of 350°C,the B3,BA31,and BA32 alloys all lie in a two-phase region consisting ofα-Mg and Mg3Bi2.According to the phase diagram,the Mg17Al12phase can be formed below extrusion temperatures of 108°C and 164°C in the BA31 and BA32 alloys,respectively.After the material exits the extrusion die,it is naturally air-cooled to RT,but Mg17Al12precipitates cannot be formed in the extruded BA31 and BA32 alloys during air-cooling because the time taken for the diffusion of Al solute atoms is insufficient for the formation of the Mg17Al12phase owing to a relatively fast cooling rate.From the SEM micrographs in Fig.5a–c,it appears that the number density of the Mg3Bi2particles increases slightly with increasing Al content,but the difference among these number densities in the three alloys is insignificant.Indeed,the fractions of the Mg3Bi2phase at 350°C calculated using the FactSage software are 3.06,3.16,and 3.25 wt% for the B3,BA31,and BA32 alloys,respectively.Although the equilibrium state in the phase fraction calculation is quite different from the dynamic state during the extrusion process,the thermodynamic calculation results along with the SEM observations suggest that the addition of Al(at least up to 2 wt%)to the Mg–3Bi alloy has little effect on the amount of Mg3Bi2particles formed in the extruded alloy.Since the extrusion temperature(350°C)is lower than the homogenization temperature(390°C),dynamic precipitation of Mg3Bi2phase can occur during extrusion.However,given that the temperature in the deformation zone near the die exit increases owing to the heat generated by deformation and friction that occurs during extrusion[7,12],the actual deformation temperature during extrusion is likely to be similar to the homogenization temperature.Consequently,dynamic precipitation of the Mg3Bi2phase will not or rarely occur under the extrusion condition conducted in this study.It is confirmed from additional SEM observations of the homogenized billets that numerous fine Mg3Bi2particles are present in all the homogenized alloys and their size and number density are similar to those of Mg3Bi2particles in the extruded alloys.This suggests that Mg3Bi2particles formed in the solidification process are partially remained after homogenization treatment because the homogenization temperature lies in a two-phase region comprisingα-Mg and Mg3Bi2in the equilibrium phase diagram(Fig.5e).These remained fine particles are rearranged along the metal flow direction during extrusion.

Fig.5.(a–c)SEM micrographs of extruded(a)B3,(b)BA31,and(c)BA32 alloys.(d)XRD patterns of extruded alloys.(e)Equilibrium phase diagram for Mg–3Bi–xAl(x=0–4 wt%),calculated using FactSage software;here,Text denotes the extrusion temperature(350°C).

Fig.6.Tensile properties of extruded alloys.TYS,UTS,and EL denote the tensile yield strength,ultimate tensile strength,and elongation,respectively.

3.2.Variation in tensile strength of extruded materials with Al addition

Fig.7.(a,b)Schmid factor(SF)maps and(c)distributions of SF for basal slip under deformation along ED of extruded alloys:(a)BA31,(b)BA32,and(c)BA31 and BA32.SFavg denotes the average SF value.

Fig.6 shows the tensile properties of the extruded alloys.The extruded BA31 and BA32 alloys have significantly higher tensile strengths than the extruded B3 alloy.In particular,the addition of just 1 wt% Al causes a large improvement in both the TYS(by 39MPa,from 136MPa to 175MPa)and the ultimate tensile strength(UTS)(by 58MPa,from 179MPa to 237MPa).This strength improvement caused by the addition of 1 wt% Al can be attributed mainly to the enhanced grain-boundary hardening effect caused by the grain refinement(from 153.3μm to 68.9μm).The Al solute atoms in the extruded BA31 alloy can also contribute to the strength increment through the solid-solution hardening effect.The extruded BA32 alloy has a finer grain structure and more numerous Al solute atoms than the extruded BA31 alloy;however,the TYS and UTS of both the extruded alloys are almost the same.Fig.7 shows the maps and distributions of the SF for(0001)<11-20>basal slip under deformation along the ED of the extruded BA31 and BA32 alloys.Since the unDRXed grains have a strong<10-10>texture,they are unfavorably oriented for basal slip during tension along the ED.Hence,the unDRXed grains have a lower SF than the DRXed grains,which is evident from blue colored areas representing the unDRXed grains in the SF maps(Fig.7a and b).As the extruded BA31 alloy has a larger amount of un-DRXed grains,the number fraction of measured points with very low SF values—lower than 0.025—is much higher in this alloy(19.9%)than in the extruded BA32 alloy(8.1%)(Fig.7c).As a result,the average SF value of the extruded BA31 alloy(0.15)is lower than that of the extruded BA32 alloy(0.18),which indicates that the texture hardening effect during tension is more pronounced in the former than in the latter.The extents of the strain hardening effect are also different in the two alloys.Fig.8a and b shows the KAM maps of the extruded BA31 and BA32 alloys,respectively.The KAM quantifies the average misorientation around a measurement point with respect to a defined set of nearest or nearest plus second-nearest neighboring points[23,24].The average KAM value of the extruded BA31 alloy(0.86)is higher than that of the extruded BA32 alloy(0.51),which means that the internal strain energy stored in the former material is higher.On the basis of the strain gradient model developed by Gao et al.[25],Kubin and Mortensen[26]proposed a simple equation for calculating the density of geometrically necessary dislocations(GNDs)using the KAM value extracted directly from EBSD data.The simplest estimate of the density of GNDs,ρGND,from the EBSD data can be expressed asρGND=2ϑ/ub,whereϑis the average misorientation angle,uis the distance over which the misorientation is measured,andbis the magnitude of the Burgers vector.From this equation,the GND density of the extruded BA31 alloy is calculated as 7.15×1013m-2,which is 68% higher than that of the extruded BA32 alloy(4.25×1013m-2)(Fig.8c).Since the flow stress increment caused by strain hardening is roughly proportional to the square root of the dislocation density[27],the extruded BA31 alloy has a stronger strain hardening effect during tensile deformation than does the extruded BA32 alloy.Therefore,although the grain-boundary hardening and solid-solution hardening effects are more pronounced in the extruded BA32 alloy,the texture hardening and strain hardening effects are more pronounced in the extruded BA31 alloy.As a consequence,the extruded BA31 and BA32 alloys have similar tensile strengths probably because of the offset between the stronger and weaker hardening mechanisms in both the alloys.

Fig.8.Kernel average misorientation(KAM)maps of extruded(a)BA31 and(b)BA32 alloys.(c)Dislocation densities of extruded BA31 and BA32 alloys.KAMavg denotes the average KAM value.

3.3.Significant improvement in ductility of extruded materials with Al addition

As the Al content increases from 0 wt% to 1 wt%and 2 wt%,the tensile elongation of the extruded material increases significantly from 2.8% to 9.4% and 16.9%,respectively;namely,the addition of 2 wt% Al results in a six-fold improvement in ductility.When extruded Mg alloys contain large second-phase particles or wide particlerich bands,the particles and bands can act as crack initiation sites during tensile deformation because of the stress concentrated around them,and such cracking eventually causes deterioration in the ductility of the extruded material[28,29].However,in the present study,in all the extruded alloys,the fine Mg3Bi2particles are uniformly distributed throughout the material(Fig.5a–c).Moreover,in extruded Mg alloys with a partially DRXed grain structure consisting of fine DRXed grains and coarse DRXed grains,{10–11}contraction twins and{10–11}-{10–12}double twins are easily formed in the coarse unDRXed grains during tension along the ED because the critical stress required to activate twinning(i.e.,twinning stress)decreases with increasing grain size[30–34].Once the contraction and double twins are formed in grains during tension,deformation is highly localized in the twinned region because the crystallographic orientations of the twins are favorable for basal slip;this consequently leads to the formation of microcracks along the twins[31,35].In the present study,optical micrographs of the longitudinal cross-section of the fractured tensile specimen of the extruded B3 alloy clearly show that several large-sized microcracks(50–300μm in length)are formed in the coarse unDRXed grains(see the red dotted circles in Fig.9a and b).Because of the large width of the unDRXed grains of the extruded B3 alloy,large twins are formed in these grains and cracking subsequently occurs along the twins during tension.In addition,since the unDRXed grains are closely spaced in the extruded B3 alloy,the twin-induced cracks formed in these grains easily coalesce with those formed in adjacent unDRXed grains,and this consequently causes the formation of sharp facture lines along the twins(see the green arrows in Fig.9b).In other words,the coalescence of adjacent cracks formed in the unDRXed grains causes the formation of a relatively large macrocrack;this formation is confirmed from the fractograph,which depicts the connection of three cleavage planes to form a large cleavage plane 710μm in length(Fig.9c).Therefore,the presence of numerous coarse unDRXed grains facilitates the formation and coalescence of cracks during tension,which eventually results in the premature fracture and resultant poor ductility of the extruded B3 alloy.

Microcracks formed in the unDRXed grains of the extruded BA31 alloy are also observed in the optical micrographs of its fractured tensile specimen(Fig.9d and e),similar to the observation in the case of the extruded B3 alloy.However,the cracks formed in the extruded BA31 alloy(30–150μm in length)are smaller than those in the extruded B3 alloy(50–300μm in length).Since the twinning stress increases with decreasing grain size[30–34],the smaller width of the unDRXed grains of the extruded BA31 alloy makes twin formation during tension more difficult.In addition,in the extruded BA31 alloy,each unDRXed grain is surrounded by many fine DRXed grains owing to the fairly high DRX fraction,and therefore,the distance between the unDRXed grains in the direction perpendicular to the tensile loading direction is quite large.Hence,as shown in Fig.9d,the distances between the cracks formed in the unDRXed grains are also fairly large—greater than 200μm;this wide distribution of cracks is contrary to the narrow distribution of cracks in the extruded B3 alloy.Therefore,although several cleavage planes are observed on the fracture surface of the extruded BA31 alloy,each cleavage plane tends to be surrounded by the fine dimples formed in the DRXed grains because of an infrequent coalescence of these planes(Fig.9f).These results reveal that the decreases in the area fraction and size of the unDRXed grains by the addition of 1 wt% Al suppress both the formation and the coalescence of twin-induced cracks in these grains during tensile deformation,and this suppression,in turn,causes an improvement in the ductility of the extruded material.

Unlike in the cases of the extruded B3 and BA31 alloys,apparent microcracks are not observed in the optical micrographs of the fractured tensile specimen of the extruded BA32 alloy(Fig.9g and h).As explained earlier,twins are mostly formed in the unDRXed grains of the extruded B3 and BA31 alloys.In contrast,twins are uniformly formed in both the DRXed grains and the unDRXed grains of the extruded BA32 alloy because the width of the unDRXed grains(23.5μm on average)is almost identical to the size of the DRXed grains(24.1μm on average).As a result,sharp fracture lines are not formed along the cracks in the coarse unDRXed grains of the extruded BA32 alloy,unlike in the cases of the extruded B3 and BA31 alloys(Fig.9h).The extruded BA32 alloy also shows a quasi-cleavage fracture surface containing several cleavage planes;however,the cleavage planes in this alloy are much smaller than those in the other two alloys and distributed uniformly on the facture surface(Fig.9i).These results indicate that because of their small size,the unDRXed grains of the extruded BA32 alloy do not cause degradation of the ductility through the preferential formation of twin-induced cracks in them.Consequently,deformation is homogeneously imposed throughout the specimen of this alloy during tension,without excessive localization of strain in the twins formed in the unDRXed grains;this homogeneous deformation results in the substantially better ductility of the extruded BA32 alloy.

The difference in the characteristics of deformation twins—which act as crack initiation sites—formed in the three extruded alloys during tension is analyzed through EBSD measurements of the longitudinal cross-sections of the fractured tensile specimens(Fig.10).In the extruded B3 alloy,several large twins are formed in a very coarse unDRXed grain with a width of 380μm,whereas few or no twins are formed in the DRXed grains(whose average size is 28.6μm)and in a relatively smaller unDRXed grain having a width of 40–98μm(Fig.10a).This twin formation in only a large unDRXed grain suggests that among the numerous unDRXed grains present in the alloy,twins are preferentially formed in large unDRXed grains owing to the lower twinning stress in them.Subsequently,cracks are initiated along the twins,which eventually causes premature fracture before the occurrence of twinning in the relatively smaller unDRXed grains.

Fig.9.Comparison of tensile fracture behavior of extruded(a–c)B3,(d–f)BA31,and(g–i)BA32 alloys:(a,b,d,e,g,h)optical micrographs and(c,f,i)SEM fractographs of fractured tensile specimens.

In the extruded BA31 alloy,many twins are formed in a large unDRXed grain having a width of 245μm,and their number density is higher than that of twins formed in the large unDRXed grain of the extruded B3 alloy(Fig.10b).Furthermore,in the case of the DRXed grains,twins are observed in some DRXed grains that are relatively larger than others(see the red arrows in Fig.10b).This observation indicates that the reduction in the size of the unDRXed grains by the addition of 1 wt% Al decreases the size of the twins formed in these grains and thereby suppresses the premature fracture caused by(i)the formation of excessively large twins and(ii)subsequent cracking at the twins.In the extruded BA32 alloy,numerous twins are formed in both the unDRXed grains and the DRXed grains(Fig.10c).Moreover,the sizes of these twins formed in the unDRXed and DRXed grains are comparable,unlike the much larger sizes of twins formed in the unDRXed grains than of those formed in the DRXed grains of the extruded BA31 alloy.This size comparability indicates that the strain imposed during tension in the extruded BA32 alloy is homogeneously accommodated via twinning uniformly throughout the material,which is contrary to the strain localization in the large unDRXed grains of the extruded B3 alloy.

For comparison of the twins formed in the DRXed and unDRXed grains of the extruded alloys,the twin boundary length per unit area for two types of twins that cause cracking(i.e.,{10–11}contraction twins and{10–11}-{10–12}double twins)is calculated from the EBSD results in Fig.10a–c;the corresponding calculation results are shown in Fig.10d.The twin boundary length in the unDRXed grains increases linearly with increasing Al content,whereas that in the DRXed grains increases slightly upon the addition of 1 wt% Al and then increases rapidly with the addition of 2 wt% Al.When 2 wt% Al is added,the twin boundary lengths in the DRXed and unDRXed grains are almost identical.On the basis of the results in Figs.9–10,it can be concluded that in the extruded Al-free B3 alloy,a few large twins are formed in some coarse unDRXed grains in the early stage of tensile deformation,which eventually results in the poor ductility of the material owing to the high plastic instability caused by localized deformation in the twins.With an increase in the Al content,both the area fraction and the size of the unDRXed grains decrease gradually,and the resultant enhancement of microstructural homogeneity leads to more uniform deformation during tension.Consequently,the tensile ductility of the extruded Mg–3Bi alloy improves significantly upon Al addition.Because of the simultaneous increases in the strength and ductility of the extruded Mg–3Bi alloy by the addition of 2 wt% Al,the comprehensive tensile property(TYS–elongation product)of the extruded material improves drastically from 381 MPa%to 2924 MPa%.

Fig.10.EBSD measurement results of fractured tensile specimens of extruded alloys:(a–c)inverse pole figure maps of(a)B3,(b)BA31,and(c)BA32;and(d)variations in twin boundary length per unit area for{10–11}contraction twins and{10–11}-{10–12}double twins in DRXed and unDRXed grains with Al content.

4.Conclusions

This study demonstrates that the addition of Al to the Mg–3Bi alloy significantly affects the microstructure,texture,and mechanical properties of the extruded material.As the Al content increases from 0 wt% to 2 wt%,the amount of second-phase particles in the extruded alloy remains almost unchanged but the average grain size decreases considerably from 153.3μm to 41.7μm and the maximum texture intensity decreases from 13.8 to 4.6;both these decreases are attributed to the decreases in the area fraction and size of the unDRXed grains with Al addition.The addition of 1 wt%Al improves the TYS of the extruded alloy from 136MPa to 175MPa because of the enhanced grain-boundary hardening and solid-solution hardening effects induced by the grain refinement and Al solute atoms,respectively.However,further addition of 1 wt% Al to the Mg–3Bi–1Al alloy does not improve its strength,which can be attributed to the weakened texture hardening and strain hardening effects.The extruded Mg–3Bi binary alloy has an extremely low tensile elongation because the formation and coalescence of relatively large cracks in the coarse unDRXed grains cause premature fracture.However,the decreases in the area fraction and size of unDRXed grains with increasing Al content suppress both the formation and the coalescence of large cracks during tension,which results in a remarkable improvement in the elongation from 2.8% to 16.9%.In the extruded BA32 alloy,contraction and double twins—which act as crack initiation sites—are uniformly formed throughout the material during tension,without preferential twinning and subsequent cracking in the unDRXed grains.Consequently,the addition of 2 wt% Al to the Mg–3Bi alloy improves the TYS–elongation product of the extruded alloy tremendously from 381 MPa%to 2924 MPa%.

Declaration of Competing Interest

The authors declare that they have no conflict of interest.

Acknowledgments

This research was supported by a National Research Foundation of Korea(NRF)grant funded by the Ministry of Science,ICT and Future Planning(MSIP,South Korea)(No.2019R1A2C1085272)and by the Materials and Components Technology Development Program of the Ministry of Trade,Industry and Energy(MOTIE,South Korea)(No.20011091).