Dislocation characteristics and dynamic recrystallization in hot deformed AM30 and AZ31 alloys

Hyeon-Woo Son,Soong-Keun Hyun

aMetallic Materials Division,Korea Institute of Materials Science,Changwon 51508,Republic of Korea

b Department of Materials Science and Engineering,Inha University,Incheon 22212,Republic of Korea

Abstract Zn addition to Mg alloys activates non-basal slip or twinning with solute softening effects.On the other hand,the effects of the Zn solute on the macroscopic dislocation behavior and dynamic recrystallization are not completely understood.Moreover,it is unclear if＜c+a＞slip can be affected by changes in the c/a ratio of solute atoms.This study was conducted to understand the solute strengthening of Zn addition and its effects on the dislocation characteristics and dynamic recrystallization.A hot torsion test was performed on both AM30 and AZ31 alloys up to a high strain to investigate the Zn solute effect on the dislocation characteristics and dynamic recrystallization.The dislocation components of the hot torsioned alloys were evaluated by X-ray line profil analysis and electron backscatter diffraction.The results showed that the Zn solutes slightly accelerate strain accumulation at the initial stages of hot deformation,which accelerated recrystallization at high strain.The dislocation characteristics were changed dynamically by Zn addition:fortifie＜c+a＞-type slip,dislocation arrangement and strain anisotropy parameters.The most important point was that the dislocation characteristics were changed dramatically at the critical strain for recrystallization and high strain regions.The addition of Zn also acted greatly in these strain areas.This indicates that the rapid formation of＜c+a＞-type slip at the serrated grain boundaries occurs at the initiation of dynamic recrystallization and the increase in the grain triple junction because grain refinemen has a great influenc on the dislocation characteristics at high strain.? 2022 Chongqing University.Publishing services provided by Elsevier B.V.on behalf of KeAi Communications Co.Ltd.This is an open access article under the CC BY-NC-ND license(http://creativecommons.org/licenses/by-nc-nd/4.0/)Peer review under responsibility of Chongqing University

Keywords:Magnesium alloy;X-ray line profile Solid solution;Dislocation density;Dynamic recrystallization.

1.Introduction

Mg alloys are the lightest structural material with the highest specifi strength among metallic materials.On the other hand,they have low formability due to a lack of slip systems.The activation of non-basal slip systems,particularly＜c+a＞-type slip systems,is essential to improving the formability of Mg alloys,while there are controversies that＜c+a＞slips is one of the instabilities which causes inhomogeneous deformation,based on single-crystal experiments[1,2].Solid solution is one of the representative methods for activating＜c+a＞type slip.For example,elements,such as Li and Y,increase the ductility of Mg alloys[3-5].Although fortifie prismatic＜a＞-slip also affects the ductility,the main reason is the activated＜c+a＞-type slip system,which provides c-axis accommodation.The reduction of the c/a ratio by solid solutions is generally considered providing lower critical resolved shear stress(CRSS)of the＜c+a＞-type slip mode.

Zn addition is generally used to enhance the corrosion resistance and improving the mechanical properties in Mg-Albased alloys.In terms of deformation mechanism,the study of Stanford and Barnett represents several important points to Zn addition effects[6].The authors described Zn solid solution mainly decreases CRSS ratio between prismatic and basal slips(τpris/τbasal)rather than CRSS ratio for{10ˉ12}twinning.However,when the range of Zn addition is limited to 0.5wt%,there is a variation in CRSS ratio between{10ˉ12}twinning and basal slip that cannot be ignored as showed in the study.This means that Zn content variation between AM30 and AZ31 alloy can provide large changes in deformation mode because Zn mainly exist as solid solution in AZ31 alloy.Moreover,the importance of Zn content on{10ˉ12}twinning at room temperature imply that Zn content may also affect＜c+a＞-type slip system at high temperature deformation,which accommodates＜c＞-axis deformation like{10ˉ12}twinning.The coincident dependences of non-basal slip and deformation twinning on the grain size support this consideration[7-9].As a result,there is a possibility that the mechanism of dynamic recrystallization(DRX)are changed by Zn addition during hot deformation of Mg alloys.

Fig.1.Schematic diagram representing(a)orientation relationship between the hexagonal structure of Mg and shear direction and(b)sectioning area of torsion specimen to observe microstructure.

X-ray line profil analysis(XLPA)has several advantages in investigating the effects of a Zn solid solution on the deformation and recrystallization mechanisms at high temperatures.The strain-anisotropy-based dislocation model was developed to determine the dislocation density,arrangement,and fraction of edge or screw dislocations as well as Burger vector populations[10-12].The dislocation contrast factors were affected by the strain anisotropy and Burgers vector.Therefore,a method to match the experimentally determined dislocation contrast factors from Williamson-Hall and Warren-Averbach plots with the numerically calculated dislocation contrast factors was developed[13,14].Recently,Lorentzian asymmetry in the X-ray line profil from planar-faulted material was found.They analyzed the asymmetry of X-ray line profil using DIFFaX software,which provides the simulated line profile of planar faulted crystals[15-17].The decomposition of the slip system components from the experimentally calculated strain anisotropy parameters was carried out by definin the effective stain anisotropy parameters.Ungár showed that it was possible to calculate the slip system by examining the tolerance conditions for several strain anisotropy parameters[18].Indeed,they carried out dislocation analysis of XLPA and electron backscatter diffraction(EBSD)in the HCP system in pure Ti and emphasized the role of the＜c+a＞dislocations on the plastic deformation of hexagonal-close-packed(HCP)metals[19,20].

Table 1Chemical compositions of the alloys.

This study investigated the dislocation dynamics and DRX behavior of AM30 and AZ31 alloys,particularly in high strain area using the hot torsion test.The information about dislocation characteristics include density,population and arrangement implying the number of dislocations per unit area,fraction of Burger’s vector and order of dislocation configuratio respectively.Also,DRX kinetics in high strain will be helpful to understand the effects of Zn addition on largely deformed Mg alloys.

2.Experimental

2.1.Sample preparation

Hot-extruded bars of AM30 and AZ31 alloys received from the Korea Institute of Industrial Technology(KITECH)were used to examine the effects of Zn addition on the hot deformation and DRX mechanism in Mg-Al based alloys.Table 1 lists the chemical compositions of the used alloys which was obtained by Inductively Coupled Plasma Optical Emission Spectroscopy(ICP-OES).The received alloys were machined to cylindrical torsion specimens with a gage length of 16 mm and a radius of 4 mm.The torsion specimens were annealed at 500 °C for 16 h to obtain full-annealed and equiaxed grain structures.After the full annealing treatment,hot torsion tests were conducted under deformation conditions of 400°C and 0.05/s and using the following orientation conditions:shear direction parallel to the c-axis of the hexagonal crystal structure to activate＜c+a＞slip,as represented in Fig.1a.The specimens were water-quenched immediately after the hot torsion tests in an automatic quenching machine to maintain the dislocation arrangements and recrystallized microstructure.The Von Mises equivalent stressσvmand strainεvmof torsion test was calculated by following equations[21]:

Fig.2.Detailed XLPA procedure of(a)Stokes Fourier approach for instrumental correction,(b)quadratic plot to obtain strain anisotropy parameters,(c)Modifie Warren-Averbach plot and(d)plot of Y/L2 vs.ln(L)to obtain dislocation density.

The specimens were then electrical-discharge-machined to observe hot-torsioned microstructure.The torsion specimens have a strain gradient along the radius of the cylinder.Sun et al.provided that when the torsion specimen is sectioned at 72.4% of radius,strain error can be minimized as shown in Fig.1b[21].The machined samples were then electropolished in a 10% of perchloric acid solution in ethanol solvent at-20 °C.The X-ray line profile were obtained using a PANalytical X’Pert powder diffractometer with a step size of 0.026° and 2θrange of 10-150°.The EBSD patterns of the samples were obtained using a scanning electron microscope(Hitachi 4300SE)with a step size of 1.0μm.Finally,the specimens for transmission electron microscopy(TEM)were prepared by twin-jet polishing using the same conditions for the electropolishing and analyzed using JEM-2100F equipment.

2.2.X-ray line profil analysis

The X-ray line profile of the(100),(002),(101),(102),and(103)reflection were used for XLPA.The instrumental broadening of the experimentally obtained X-ray line profile was corrected using the Stokes method[22].Fig.2a presents an example of the correction procedure.Although instrumental-corrected line profil by Fourier transform represents serration in the baseline profile it is sufficien to obtain the full-width at half-maximum(FWHM).Considering that X-ray line broadening occurs by the size and strain parameters,double-Voigt fittin was used to obtain the FWHM of the instrumental-corrected line profiles The strain anisotropy parameters can be represented in hexagonal systems[13]:

Fig.3.Initial microstructures of(a)AM30 and(b)AZ31 and initial texture distributions of(c)AM30 and(d)AZ31.

whereandare the measured strain anisotropy parameters associated with the elastic constants;irepresents the Burgers vectors;is the average contrast factor at each Burgers vector;hiis the fraction of each slip system.

The quadratic equation converted from themodifieWilliamson-Hall(WH)equation,as represented in Fig.2b,was used to determine theand:

whereK=2 sinθ/λ;K=2 cosθ(Δθ)/λ;α=(0.9/D)2;;x=(2/3)(l/ga)2.Dis the crystallite size;Mis a constant;ρis the dislocation density;ais the lattice constant in the basal plane;lis the last Miller index of the(hkl)set;gis the diffraction vector.Theandvalue can be determined via regression analysis in the quadratic fitting The fractions of the slip system,ha,hc,andhc+a,were calculated using effective strain anisotropy parameters,and,under the tolerance conditions suggested by Ungár et al.[18].The dislocation contrast factors ofandcan be determined by substitutinghiinto Eq.(6)[13]:

AmodifieWarren-Averbach(WA)plot,as shown in Fig.2c,was used[23]:

whereA(L)is the real part of Fourier coefficientsLis the Fourier variable;Reis the effective outer cut-off radius of dislocations.A(L)can be given by Arechabaleta et al.[23]:

whereFLandFGare the Lorentzian and Gaussian parts of FWHM.Using themodifieWA plot of lnLvs.,Y=,can be given by:

Fig.4.Flow curves of the AM30 and AZ31 alloys.

Fig.5.Plots of the dislocation density and dislocation arrangement parameter of AM30 and AZ31 alloys.

Fig.6.XLPA results representing fraction of(a)＜a＞-type dislocations,(b)＜c+a＞-type dislocations and(c)screw dislocations.

As shown in the Fig.2d,the dislocation density can be calculated using a linear plot betweenand lnL.

3.Results

Fig.3 represents the initial microstructures of AM30 and AZ31 alloys.The long-time heat treatment generated equiaxed structures with similar grain sizes and minimized internal strain.The average values of initial grain sizes were measured as 29.7μm in AM30 and 27.2μm in AZ31.The initial texture distributions of both AM30 and AZ31 have strong basal components perpendicular to the shear plane normal(SPN),even though AZ31 has a slightly dissipated distribution.As a result,＜c+a＞slip can be strongly activated in this orientation relationship as represented in Figs.1 and 3.

Fig.4 shows the fl w curves of AM30 and AZ31 alloys.The hot-torsion tests were conducted by three times to obtain experimental reliability.When the experiment was repeated,the fl w curve showed no significan deviation which would be originated from the following considerations:Since the torsion specimens was machined from hot-extruded rods,there were almost no casting defects.In addition,the high temperature heat treatment for a long time provided full-annealed and equiaxed structure,so that the fl w curve showed no significant deviation.As shown in other studies,the fl w stresses of the AM30 and AZ31 alloys did not show a significan difference.On the other hand,there was a deviation in the slope of the fl w curve,which resulted in two intersection points in the fl w curve.The intersection of the fl w curves occurred mainly in the critical strain for DRX and high strain region where DRX was sufficientl progressed.Hence,different strain hardening as well as softening mechanisms were expected between AM30 and AZ31 alloys.

Fig.5 presents the changes in dislocation density and dislocation arrangement parameters with varying strains for AM30 and AZ31 alloys.The variation of the dislocation density showed a similar pattern in both AM30 and AZ31 alloys.The dislocation density in both alloys decreased sharply at a strain of 0.3,and then showed a slow but gradual decrease.It indicates that typical DRX occurred at both AM30 and AZ31 alloys.On the other hand,the dislocation density of the two alloys showed different behavior based on a strain of 0.7.Although the dislocation density of AZ31 was high near the critical strain for the DRX region,it showed a more decreasing pattern than AM30 as recrystallization proceeded.This suggests that the continuous accumulation of strain energy due to Zn addition affected the recrystallization kinetics at high strain.In the case of the dislocation arrangement parameter,the two alloys showed a similar pattern.Unlike the dislocation density,however,it increased in proportion to the strain and then decreased in the high strain region.The physical meaning of the dislocation arrangement parameter can be described as following:when the dislocation arrangement parameter is smaller than 1.0,it shows strong screening behavior indicating dislocations arranged in an order,such as in a dipole configuratio which can reduce strain fiel of dislocation.While larger value than 1 shows a random arrangement of dislocations[24].Both AM30 and AZ31 alloys have smaller values than 1.0 at the low strain area and reached one at about a strain of 0.6 to 0.7.The strain of 0.7 is the intersection point of the two fl w curves in Fig.4 and the reverse point of the dislocation densities of AM30 and AZ31 alloys.This also suggests that there is a change in the deformation and recrystallization mechanism after a strain of 0.7,where the DRX is occurring.The dislocation arrangement parameter of the AZ31 alloy showed higher values than AM30 at strains lower than 0.9,but rather lower values at strains larger than 0.9.This means that the turning point of the deformation and recrystallization mechanism,which was previously define as a strain near 0.7,occurred at a lower strain in AZ31.

Fig.7.Semi-quantitative strain mapping using GOS value representing strain distributions of(a,c,e)AM30 and(b,d,f)AZ31.Red-fille grains indicate recrystallized grains determined by GOS value lower than 2°.

The slip system variation with strain showed a similar tendency with the changes in dislocation density and dislocation arrangement parameters.Fig.6 shows the variation of the slip system and dislocation characteristics with strain changes.Fig.6a and b present the fraction of＜a＞-type and＜c+a＞-type slip variations.The change in the slip system also had a large transition point in the critical strain for DRX and high strain regions,as previously emphasized.The＜c+a＞-type slip system change shown in Fig.6b represent that both AM30 and AZ31 alloys had a strong increase in the slip system fraction in the critical strain for DRX and high strain region.On the other hand,the change in the fraction of this slip system was changed significantl in the AZ31.The＜c+a＞-type slip fraction of AZ31 tended to increase significantl compared to AM30 in the critical strain for DRX and high strain area.When Zn was added,the change in screw dislocation was also significantl changed near strain of 0.3 and 0.7,similar to the＜c+a＞slip behavior.Although the results of XLPA alone are inconclusive,these results suggest that＜c+a＞-type slip with screw components can be activated in the critical strain for DRX and high strain region.

Fig.8.(10ˉ11)pole figure of(a)AM30 and(b)AZ31.Red-circled area and internal numbers indicate pole intensity of the normal direction of(10ˉ11)plane parallel to SPN.

Fig.7 presents the stored energy distribution of the hotdeformed microstructures.As predicted from the XLPA results,strain hardening is much higher in the AZ31 alloy by observing high grain orientation spread(GOS)values at the strain of 0.2.Similarly,a slightly accelerated DRX rate was observed in the AZ31 alloy after critical strain for DRX.The smaller average grain size of AZ31 as represented in Fig.7 also show that AZ31 alloy has a higher DRX rate during hot deformation.

The evidence for enhanced＜c+a＞-type slip in AZ31 alloy was provided by comparing texture evolution of AM30 and AZ31 alloys.Fig.8 shows the(10ˉ11)pole figure of the AM30 and AZ31 alloys.If＜c+a＞slip occurs strongly in the torsion test,the slip surface of＜c+a＞slip will be strongly oriented along the direction of SPN.Based on this assumption,a strong(10ˉ11)pole intensity was observed in AZ31 when analyzing the(10ˉ11)pole intensity in the SPN direction on the pole figure Hence,the high＜c+a＞-slip fraction of the AZ31 alloy orients the＜c+a＞-type slip plane in the SPN direction,as analyzed in XLPA.

The Schmid factor is used to quantify the＜c+a＞-type slip surface oriented in the SPN direction.Fig.9 shows the Schmid factors for shear directions of(10ˉ11)and(11ˉ22)＜c+a＞slip into the shear direction.The(10ˉ11)slip plane oriented to the direction of SPN showed a slight decrease in AM30 while it increased in AZ31.In the case of the(11ˉ22)slip plane,both alloys represented an increasing tendency in the Schmid factor.In the case of the(11ˉ22)slip plane,the Schmid factor was generally lower than that of the(10ˉ11)slip plane,but it increased significantl with increasing strain.Although the average Schmid factor does not have a linear relationship with the fraction of slip system,＜c+a＞-type slip occurs strongly in AZ31 where both Schmid factors for slip planes of(10ˉ11)and(11ˉ22)increase continuously with strain.

Finally,strain accumulation in the Schmid factors for＜c+a＞-type slip planes was measured semi-quantitatively,which is helpful for identifying strain hardening of＜c+a＞-type dislocations.Fig.10a shows the kernel-average misorientation(KAM)values in terms of the Schmid factors of AM30 and AZ31 at a strain of 0.2 before DRX initiates.In this strain,higher KAM values are distributed in AZ31 due to the strain hardening effect by Zn,as described above.The notable point is that the deviation between the minimum and maximum values in the KAM distribution for each Schmid factor of AM30 and AZ31 is significantl different.In the case of AM30,the deviation of the KAM value at low and high Schmid factors was approximately 0.10,but in AZ31,it was approximately 0.18.Similar behavior is also observed in the high strain area,as shown in Fig.10b.A low KAM value was noted in AZ31,which was found to have a higher DRX rate because it is a strain region after DRX.In addition,the deviation of the calculated KAM value in Fig.10b was approximately 0.15 at AM30,whereas it was approximately 0.31 at AZ31.The KAM deviation between the grains with a low Schmid factor and the grains with a high Schmid factor was higher in AZ31,which was similar to the case of low strain.This indicates that the rate of orientation transition into the condition of SPN//＜c+a＞-type slip plane was much faster in the AZ31 alloy because the high deviation of the KAM value indicates a high driving force to change their slip-induced grain rotation.

Fig.9.Average Schmid factors of(10ˉ11)and(11ˉ22)into the shear direction in AM30 and AZ31 alloys.

Fig.10.Schmid factor vs.KAM of AM30 and AZ31 alloys in low(0.2)and high(0.8)strain.

4.Discussion

According to the recent atomistic simulation results,it was reported that an increase in the Zn content activates the＜c+a＞-type slip by reducing the critical resolved shear stress required for＜c+a＞-type slip[25].On the other hand,the simulated strain range was limited to low strain(0.1),unlike this study;thus,a novel interpretation in the results of this study is necessary.Considering that the Zn content within 1%does not have a significan effect on the c/a ratio of the Mg crystal,the change in＜c+a＞-type slip caused by a change in the Zn content in the high strain area mainly depends more on the strain hardening effect of Zn solutes than the variation in the lattice constant[26].The strain hardening effect in the initial strain of hot deformation due to an increase in the Zn content is represented by the increase in the dislocation arrangement parameter in AZ31 alloy.The high dislocation arrangement parameter at the initial strain in the AZ31 means that the addition of Zn can increase the dislocation density by increasing the interactions between the dislocations and solute atoms.On the other hand,in the high strain area of AZ31,the softening rate increases due to the enhanced dislocation forest formation and the high dislocation density,suggesting that the dislocation screening is rapidly dissociated accordingly.

Dislocation characteristics have two significan turning points at the critical strain for DRX and strain of 0.8.Zn content affects the turning point of the dislocation characteristics.An increase in＜c+a＞-type slip and screw dislocation component in the critical strain for DRX is likely associated with the stress concentration at the serrated grain boundary.Grain boundary serration is one of the results of dynamic recovery,which has the maximum degree at the initiation of DRX[27,28].Since the grain boundary serration can induce stress concentration by grain boundary sliding(GBS)during hot deformation,they can form a＜c+a＞-type slip,and their geometrical properties can be related to the formation of screw dislocations[29].Fig.11a and b show nucleation of the dislocation array at the serrated grain boundary and the transition procedure into DRX grain.The dislocation arrays nucleated from the serrated grain boundary are being converted to the low-angle grain boundary(LAGB)from the region close to the serrated grain boundary.The dislocations enclosing the serrated grain boundary have the same contrast and line vector in the given g-vector,so they are all the same component.On the other hand,their moving direction is very different from the gliding direction.This behavior of dislocations means that a strong stress concentration occurs in this region.Hence,the climb or cross slip of dislocations can occur relatively freely.This means that stress concentration occurs near the serrated grain boundary,and＜c+a＞-type slip can be generated strongly at the serrated grain boundary.Fig.11c also shows evidence for the screw component of the dislocation arrays at the serrated grain boundary.Although thermal recovery of dislocations makes difficul to analyze the Burgers vector accurately,a thermally recovering dislocation array was formed from the vertices of serrated grain boundary and found traces that changing paths by cross-slip(red circle).The screw component and＜c+a＞-type slip can be interpreted because the accommodation for the GBS at the serrated grain boundary causes stress concentrations in the direction parallel to the grain boundary[29].This suggests that the Zn solute increases the stress concentration at the serrated grain boundary by grain boundary segregation or strengthening the interaction between dislocations.

Fig.11.TEM image representing hot deformed AZ31 at 400 °C,0.05/s and strain of 0.7.Each figure show(a)formation of a dislocation array and its transition procedure into LAGB at the serrated grain boundary,(b)same figur obtained from a different g-vector and(c)another example representing LAGB formation at the serrated grain boundary,which is a more recovered state than(a)and(b).

In the case of the high strain area,grain refinemen caused by DRX would dominate＜c+a＞-type slip generation.TEM study of the grain size effect on slip system showed that the＜c+a＞-type slip was activated when the grain size was 5.5μm in pure Mg[30].This means that the accelerated grain refinemen by the Zn content in the high strain area has advantages in activating＜c+a＞-type slip.In particular,when grain refinemen occurs by DRX,the number of grain triple junctions increase,which has a similar role to the serrated grain boundaries on the stress concentration[31].When grain refinemen occurs by DRX,grain rotation deviation between the parent and DRX grains may occur significantl in the early stages of DRX because the small DRX grains cause an independent grain rotation system[32].This causes stress concentration in the grain triple junction and may help increase the fraction of＜c+a＞-type slip at high strain during high-temperature deformation.

5.Conclusions

This study examined the effects of Zn solutes on the dislocation characteristics and slip system variation in Mg-Albased alloys by comparing commercial AM30 and AZ31 alloys.The XLPA,EBSD and TEM study on the hot-torsioned AM30 and AZ31 revealed the following results:

(a)Zn addition accelerates strain hardening at the initial stage of hot deformation and overall DRX rate.

(b)At the strain hardening stage,strong interactions between dislocations and weak screening behavior of the dislocations were expected by increased dislocation arrangement parameters in the AZ31 alloy.In contrast,the dislocation arrangement parameter decreased when DRX was sufficientl progressed in AZ31 alloy.This indicates that higher DRX rate causes rapid dislocation screening at high strain area.

(c)The dislocation characteristics,such as the slip system and dislocation components,were changed rapidly at a critical strain for DRX and specifi high strain areas where DRX sufficientl occurs.The Zn content accelerates the rapid changes in physical parameters,particularly＜c+a＞-type slip and screw dislocation component.

(d)The reason for the higher fraction of＜c+a＞-type dislocations and screw components were the maximized grain boundary serration at critical strain for DRX and the increased number of grain triple junctions caused by grain refinemen of DRX at the high strain area.

Acknowledgment

This work was supported by the Inha University Research Grants.