Deformation mechanism of fine grained Mg–7Gd–5Y–1.2Nd–0.5Zr alloy under high tem

Wanru Tang,Zheng Liu,Shimeng Liu,∗,Le Zhou,Pingli Mao,Hui Guo,Xiaofang Sheng

aSchool of Materials Science and Engineering,Shenyang University of Technology,110870,China

b Key Laboratory of Magnesium Alloys and the Processing technology of Liaoning Province,110870,China

c Wanfeng Auto Holding Group,Shaoxing,312500,China

Received 12 July 2019;received in revised form 17 February 2020;accepted 23 February 2020 Available online 26 June 2020

Abstract Fine grained Mg–7Gd–5Y–1.2Nd–0.5Zr alloy was investigated by dynamic compression tests using a Split Hopkinson Pressure Bar under the strain rates in the range 1000–2000 s−1 and the temperature range 293–573K along the normal direction.The microstructure was measured by optical microscopy,electron back-scattering diffraction,transmission electron microscopy and X-ray diffractometry.The results showed that Mg–7Gd–5Y–1.2Nd–0.5Zr alloy had the positive strain rate strengthening effect and thermal softening effect at high temperature.The solid solution of Gd and Y atoms in Mg–7Gd–5Y–1.2Nd–0.5Zr alloy reduced the asymmetry ofα-Mg crystals and changed the critical shear stress of various deformation mechanisms.The main deformation mechanisms were prismatic slip and pyramidalaslip,{102}tension twinning,and dynamic recrystallization caused by local deformation such as particle-stimulated nucleation.

Keywords:Mg–7Gd–5Y–1.2Nd–0.5Zr magnesium alloy;High strain rate deformation;Local deformation mechanism.

1.Introduction

With the continuous improvement of carrying capacity and air-rang requirements,lightweight material has become an important development direction of aircraft,satellites,rockets,and other aircrafts[1–4].Therefore,magnesium alloy with the advantages of low density,high specific strength,good shock absorption performance,and good damping performance has been widely used in the field of aerospace[5–8].In the field of national defense,magnesium alloy with lightweight and high-performance characteristics will effectively reduce the weight of spacecraft and significantly improve its maneuverability[9–12].

The plastic deformation of magnesium alloy has been mainly investigated under quasi-static state[6,13,14]and less studied under high strain rate[15–17],and AZ series alloy has given more attention[18–21].Slip and twinning are the main micromechanism of deformation at high strain rate in AZ31 magnesium alloy[22],and the types of slip and twinning are related to the crystal texture[23],while AZ31 magnesium alloy in rolling or extrusion state usually has strong basal texture.Lou et al.examine the dynamic mechanical properties of the AZ31 alloy parallel to the rolling direction.[10–12]tension twinning is found as the main deformation mechanism in the early deformation stage[24].In addition,the{0001}texture has the most obvious effect on the micromechanism of compression deformation of AZ31 magnesium alloy at high strain rate when it is compressed in the transverse direction[25].There is an inflection point in its true stress–strain curve.The dominant mechanism of plastic deformation before the inflection point is tension twinning,while that after the inflection point is non-basal sliding.Furthermore,when rare-earth elements are added to the magnesium alloy,the texture will become weaker and more random,decreasing the degree of twinning and enhancing the slip deformation mode[26].

Fig.1.The schematic diagram:(a)3-pass MDF;and(b)sample interception.

At present,WE54 and WE43 are the most popular highperformance magnesium alloys in the field of aeronautics and astronautics.However,the existing data shows that the newly developed Mg–Gd–Y–Zr alloy,such as JDA2 alloy of China,which has far ahead of that two kinds of alloys in room temperature and high temperature properties[27–29].Moreover,these Mg–Gd–Y alloys can be further strengthened and toughened by severe deformation.Yang Q.et al.[30]apply friction stir processing to Mg–10Gd–3Y–0.5Zr alloy,and Zhang F.et al.[31]apply equal channel angular pressing to Mg–11.90Gd–0.81Y–0.44Zr alloy.Tang L.et al.[32]apply multi-directional forging(MDF)to Mg–10Gd–4.8Y–0.6Zr alloy to further improve the strength and toughness of the alloy through severe deformation induced grain refinement.In order to minimize the increase of alloy density caused by high Gd content,MDF reinforced Mg–7Gd–5Y–1.2Nd–0.5Zr alloy is designed[26].After alloying with various and high contents of rare-earth elements,the average grain diameter of Mg–7Gd–5Y–1.2Nd–0.5Zr alloy can be refined to 5μm,the yield strength can be>250MPa,and the elongation is still as high as 7.5%.Although the alloy has been widely used in rockets,spacecraft and other products,requiring extremely lightweight characteristics,high temperature and dynamic instantaneous performance,there are a few reports on its deformation mechanism at high temperature and high strain rates.Thus,this study aims to investigate the deformation mechanism of MDF reinforced Mg–7Gd–5Y–1.2Nd–0.5Zr alloy under high temperature and high strain rates in order to provide necessary data and design basis for launching rockets,spacecraft,etc.

2.Experimental procedure

2.1.Material

The Mg–7Gd–5Y–1.2Nd–0.5Zr alloy used in the study had a nominal mass fraction of Mg–7Gd–5Y–1.2Nd–0.5Zr,and an actual chemical composition of Mg–6.85Gd–4.52Y–1.15Nd–0.55Zr alloy[26].After 2×3 passes of MDF as shown in Fig.1(a)[26,32],it was extruded,and samples were taken from the extruded parts as shown in Fig.1(b).A pass strain of=0.5 was employed.The MDF was carried out at initial forging temperature of 530°C and a pressing speed of 10mm/s,and the time interval between particular passes was within 15s.The extrusion ratio was 20,and the extrusion temperature was 400 °C.The samples used in this study were processed into the cylinders of8×6 mm2dimension along the normal direction(ND).

Fig.2.The experimental device diagram of Mg–7Gd–5Y–1.2Nd–0.5Zr.

2.2.Dynamic compression experiment

A Split Hopkinson Pressure Bar(SHPB)and a heating furnace were used to characterize high speed deformation behavior of Mg–7Gd–5Y–1.2Nd–0.5Zr alloy in the strain rate range 1000–2000s−1from 293K to 573K,as shown in Fig.2.In the experimental process,the incident waveεiwould be generated first when the bullet impacts the incident bar,and the specimen between the incident bar and the transmission bar would be compressed under the pulse.At the same time,the reflected waveεrwould reflect back to the incident bar,and the transmitted waveεtwill transmit to the transmission bar.The waveform and size of the incident waveεi,the reflected waveεr,and transmitted waveεtwere collected by a data acquisition system.According to one-dimensional stress wave theory,the stressσ(t),strainε(t),and strain rate(t)of the experimental materials could be obtained as follows:

Fig.3.True stress–strain curves of Mg–7Gd–5Y–1.2Nd–0.5Zr at different strain rates under different temperatures.

whereSandS0are the cross-sectional areas of the waveguide bar and sample,respectively,Vis the velocity elastic wave,His the height of the specimen,andεi(t),εr(t),andεt(t)are the signals of the incident wave,the reflected wave,and the transmitted wave,respectively.

2.3.Microstructure observation

Fig.4.True stress–strain curves of Mg–7Gd–5Y–1.2Nd–0.5Zr at the strain rate of 1000 s−1 under different temperatures.

The microstructure and microtexture were characterized by X-ray diffractometry(XRD),optical microscopy(OM)using a zeissoptical microscope,electron backscatter diffractometry(EBSD),and transmission electron microscopy(TEM).The metallographic sample was cut along the axis direction of the cylindrical sample as shown in Fig.1(b),then inlaid,ground,and polished.A mixture of 16mL ethanol,2mL glacial acetic acid,0.96g picric acid,and 2mL distilled water was used to corrode the sample for 20s,and then the samples were washed with ethanol and dried using a fan.For the EBSD samples,electrolysis was required after grinding and polishing.The electrolyte was a mixture of 10%perchloric acid and ethanol solution.The parameters of electrolysis experiments are as follows:electrolysis voltage,8V;electrolysis current,0.04 A;electrolysis temperature,−30 °C;electrolysis time,60s.For the preparation of TEM samples,the samples were ground to 50μm size and punched into a3mm disk using a hole punch,and then the TEM observation was carried out after the disk samples were thinned using an ion shearing instrument.In this way,the materials could be analyzed comprehensively from different angles.

Fig.5.The yield strength–strain rate of Mg–7Gd–5Y–1.2Nd–0.5Zr under different temperatures.

Fig.6.Diffraction patterns of Mg–7Gd–5Y–1.2Nd–0.5Zr alloy at 293K.

Fig.8.Distribution of SF in grains under different slip systems:(a)basal slip;(b)prismatic slip;(c)pyramidalslip;and(d)pyramidalslip.

3.Results

3.1.High strain rates deformation under different temperatures

Fig.7.The initial microstructure and microtexture of Mg–7Gd–5Y–1.2Nd–0.5Zr:(a)OM map;(b)EBSD map;and(c)pole figure.

The compressed deformation of Mg–7Gd–5Y–1.2Nd–0.5Zr alloy along ND in the strain rate range 1000–2000s−1under the temperature range 293–573K is displayed in Fig.3,clearly indicating that the flow stress increases with increasing strain rate at the same temperature,showing the positive strain rate strengthening effect.The curves begin to change continuously without obvious yield platform,showing the deformation characteristics of continuous yield.It is noteworthy that all the curves are convex upward rather than concave first and then convex upward,as in the case of ultra-fine grain AZ31 alloy specimens[33].The concave upward curve is the result of tension twinning with lower critical shear stress and its promotion to the basal slip[34].

Fig.9.Numerical statistics of SF in different slip systems:(a)basal slip;(b)prismatic slip;(c)pyramidalslip;and(d)pyramidalslip.

Fig.4 illustrates the comparison of true stress–strain curves at a strain rate of 1000 s−1under different temperatures.The true stress–strain curves keep moving down at the same strain rate with increasing temperature,i.e.,the dynamic compressive strength decreases continuously.

Fig.5 depicts the yield strength–strain rate of Mg–7Gd–5Y–1.2Nd–0.5Zr at different strain rates under different temperatures,indicating increasing yield strain with increasing strain rate and decreasing yield strain with increasing temperature.The mechanical behavior of the alloy was found to be much more sensitive to temperature than to strain rate.

3.2.Effect of rare-earth solid solution on lattice constants

Diffraction patterns of Mg–7Gd–5Y–1.2Nd–0.5Zr alloy at 293K were measured by XRD,as shown in Fig.6,and the lattice constantsa,c,andc/awere calculated.The test results are as follows:the lattice constantsa1=0.3226nm andc1=0.5209nm were obtained by fitting the test data.Thus,the axial ratioc1/a1of Mg–7Gd–5Y–1.2Nd–0.5Zr alloy was determined to be 1.6147.At 293K,the lattice constanta0of the pure magnesium is 0.3202nm,c0is 0.5199nm,and the axial ratioc0/a0is 1.6237[35].Obviously,when rare-earth elements are added to magnesium,the change ofaandcis shown in the following equations:

whereis the change ofa,andis the change ofc.When rare-earth atoms enter the magnesium matrix as solid solution atoms,the atomic spacing increases.The lattice constantsaandcof Mg–7Gd–5Y–1.2Nd–0.5Zr alloy are larger than those of magnesium.Because the effect of solid solution atoms on the a-axis is greater than that on the c-axis,is larger than.Soc1/a1is smaller thanc0/a0,indicating that the solid solution of Gd and Y intoα-Mg will decrease the asymmetry of the close-packed hexagonal crystals.

Fig.10.The microstructure of Mg–7Gd–5Y–1.2Nd–0.5Zr alloy at a strain rate of 1000 s−1:(a)OM map at 293K;(b)EBSD orientation map at 293K;(c)boundary misorientation map at 293K;(d)OM map at 373K;(e)EBSD orientation map at 373K;(f)boundary misorientation map at 373K;(g)OM map at 473K;(h)EBSD orientation map at 473K;(i)boundary misorientation map at 473K;(j)OM map at 573K;(k)EBSD orientation map at 573K;and(l)boundary misorientation map at 573K.

3.3.Microstructure under different temperatures

The initial microstructure and microtexture along the ND are displayed in Fig.7 with an average grain size of 5μm.The results show that{0002}basal texture formed during the extrusion process,and the c-axes of most grains are parallel to transverse direction with a maximum intensity of 4.85.The green color in Fig.7(c)shows the characteristic dispersion distribution.Compared to AZ31 alloy[25],which does not contain rare-earth elements and has a grain size of approximately 10μm,the{0002}base texture strength of fine grained Mg–7Gd–5Y–1.2Nd–0.5Zr alloy is obviously weakened and has more random distribution.

Fig.8 displays the Schmid Factor(SF)of the slip systems along the ND.Clearly,the color of most of the grains is red and dark yellow in Fig.8(b)and 8(c),respectively,showing that the SF of the prismatic slip and pyramidalaslip is larger.Fig.9 displays numerical statistics of SF in different slip systems.

The microstructure of Mg–7Gd–5Y–1.2Nd–0.5Zr alloy at a strain rate of 1000s−1was characterized by OM at different temperatures as shown in Fig.10(a),(d),(g),and(j).The microstructure of the compressed specimens shows uneven grain size.It can be clearly seen that dynamic recrystallization(DRX)and some deformation bands appear in the microstructure at a strain rate of 1000 s−1under different temperatures as characterized by EBSD shown in Fig.10(b),(e),(h),and(k).With increasing temperature,DRX grains grow gradually,and new DRX grains appear continuously.A small amount of{10-12}tension twinning is shown in Fig.10(c),(f),(i),and(l)at a strain rate of 1000 s−1under different temperatures.Almost all of these{10-12}tension twinning occurs in large grains.

Fig.11 displays the grain size distribution of Mg–7Gd–5Y–1.2Nd–0.5Zr alloy at a strain rate of 1000 s−1under different temperatures from EBSD.The average grain diameter calculated from the EBSD data is 3.43,2.67,2.70,and 3.00μm.The average grain diameter first decreases with the temperature,and then increases with the temperature,but it is always smaller than the initial state,which is due to the generation and growth of DRX by high strain rate deformation.

Fig.11.The grain size distribution of Mg–7Gd–5Y–1.2Nd–0.5Zr alloy at a strain rate of 1000 s−1 under different temperatures:(a)293K;(b)373K;(c)473K;and(d)573K.

At present,some studies[23,36]have revealed the effect of grain refinement on twinning and slip.Meyers et al.[37]has shown that twinning stress increases with decreasing grain size more rapidly than the stress required to activate slip,i.e.,Kin the Hall-Petch relationship for twinning-dominated flow is frequently greater than that for the slip-dominated flow.

Not only the grain size,but also the precipitated second phase has a great effect on the microstructure evolution and mechanical properties of the alloy.Fig.12 shows the TEM images of Mg–7Gd–5Y–1.2Nd–0.5Zr alloy at a strain rate of 1000 s−1under different temperatures.The stacking of dislocations occurs after dynamic compression,as shown by the arrow in Fig.12(a)and 12(b).Fig.12(c)and 12(d)show white and bright dislocation-free areas in the high dislocation areas along the front of the second phase particles,indicating that DRX may occur in these areas Fig.13.

4.Discussion

Fig.3 shows that the true stress–strain curves of Mg–7Gd–5Y–1.2Nd–0.5Zr alloy at different strain rates and temperatures are convex upwards,indicating that the tension twinning–basal slip does not play a dominant role in the early deformation stage,and the deformation mechanism is mainly controlled by the non-basal slip,which is similar to those reported in literature[7].In other words,a small amount of tension twinning and basal slip do not play a dominant role in the macro-deformation of the alloy,whereas prismatic or pyramidal slip plays a dominant role in the whole deformation process.

4.1.Plastic deformation dominated by non-basal slip mechanism

For close-packed hexagonal crystals,based on the rigid sphere model,when the axial ratioc/ais equal to 1.633,the atoms on the(0002)plane are most closely aligned.When the axis ratioc/ais less than 1.633,the difference in the atomic arrangement density among(0002),(10-10),and(11-22)planes decreases with decreasingc/avalue.When thec/avalue is small enough,the(10-10)prismatic plane or(11-22)pyramidal plane may be also the slip planes[38,39].

Fig.12.TEM images of Mg–7Gd–5Y–1.2Nd–0.5Zr alloy at a strain rate of 1000 s−1 under different temperatures:(a)293K;(b)373K;(c)473K;and(d)573K.

Fig.13.(a)Important crystal planes involved in tension twinning;and(b.)atomic displacements in(110)plane during tension twinning.

The XRD results shows that the lattice constants of Mg changed by the addition of the solid solution of rare-earth elements Y and Gd inα-Mg,decreasing thec/aaxis from 1.6237 of pure Mg to 1.6147 of Mg–7Gd–5Y–1.2Nd–0.5Zr alloy.Thus,the difference in the critical shear stress between the basal slip and non-basal slip decreased,and the non-basal slip such as prismatic plane or pyramidal plane slip can occur easily.

In addition,SF is also widely used to measure the feasibility of the plastic deformation mechanism of alloys[40,41].Fig.9 shows that the number of high SF grains in pyramidalaslip system is the largest,followed by the number of high SF grains in the prismatic slip system,and the number of high SF grains in the basal and pyramidalc+aslip system is less.Therefore,prismatic slip and pyramidalaslip should be the main slip mechanism of Mg–7Gd–5Y–1.2Nd–0.5Zr along the ND combined with the axial ratio.

4.2.Tension twinning-basal slip mechanism controlled by c/a axial ratio

Fig.10 shows that the tension twinning in Mg–7Gd–5Y–1.2Nd–0.5Zr alloy is basically in larger grains,and the number of tension twinning is very small compared to AZ31 magnesium alloy[42–44].The grain refinement can enable the suppression of{10-12}tension twinning[45].The stress required for twinning has a higher grain-size dependence than that required for dislocation slip,suggesting that the activation of{10-12}tension twinning is inhibited in the ultrafine grained microstructure[46].In fact,it is not only related to the grain size but also to thec/aaxis ratio of the crystal.Among all the twinning types of magnesium alloys,the shear variable of tension twinning is the smallest;therefore,it is also the easiest to produce.However,tension twinning usually produces when the alloy is compressed perpendicular to the c-axis or stretched parallel to the c-axis[47,48].

Fig.13(a)shows that the(10-12)plane is a tension twinning plane,and the atoms on both sides of the plane will be symmetrically arranged after the tension twining occurs at(1-210)plane.Fig.13(b)shows the arrangement of atoms on the(10-12)plane.Among them,the lattice Q is the mirror image of the lattice N.When it is compressed perpendicular to the c-axis or stretched parallel to the c-axis,the lattice P moves to lattice Q and tension twinning occurs.

The distance between P and Q depends on thec/aaxial ratio.Whenc/a=1.732,POMN is a square,and P and Q are symmetrical;therefore,the distance between P and Q is zero.Whenc/aincreases,the distance between P and Q increases,i.e.,higher critical shear stress is required for tension twinning.When the axial ratio decreases from 1.6237 of pure magnesium to 1.6147 of Mg–7Gd–5Y–1.2Nd–0.5Zr alloy,the possibility of tension twinning decreases,which may be the reason why the tension twinning-basal slip does not become the dominant deformation mechanism.Thus,the curve in Fig.3 does not show a concave–convex shape.

4.3.Localized deformation and particle-stimulated nucleation mechanism

The localized deformation is more likely to occur when the alloys with uneven grain size are under compressive load with high strength and low thermal conductivity[49,50].Local deformation occurs preferentially in the fine grained zones with high SF on the prismatic slip or pyramidalaslip,and a large number of dislocations accumulate at the grain boundaries and phase boundaries(Fig.12).With increasing stacking dislocations,the macroscopic behavior is local deformation hardening.Dislocation accumulation will increase local deformation energy and thermal energy.When the stored deformation energy and heat energy reach a certain critical value,recrystallization will occur,and it is characterized by local grain refinement in the structure and local thermal softening in performance.

Microstructural inhomogeneity also contributes to the local accumulation of dislocations and even leads to dynamic recrystallization.Particle-stimulated nucleation(PSN)has been discovered in some magnesium alloys[51–53].PSN mechanism was also observed in Mg–7Gd–5Y–1.2Nd–0.5Zr alloy in this experimental investigation.As shown in Fig.12,highdensity dislocation entanglements gathered around the secondphase particles with a size of approximately 0.8μm or more.These intense deformation zones are mainly composed of several dislocation cells,which can provide nucleation sites for DRX grains as shown by the arrow in Figs.12(c)and 12(d).The higher the temperature,the greater the effect of PSN mechanism.With the increase of temperature,more dislocations are activated,which helps to improve the nucleation rate of DRX under the PNS mechanism,and further refines the grains.On the other hand,with the increase of temperature,the mobility of dislocation increases,and the tendency of DRX grains growth increases,which inevitably leads to the increase of grain size.Obviously,as the DRX proceeds,the grain size is further refined and the twinning will be more difficult.

5.Conclusions

The conclusions of this study are as follows.

(1)The flow stress increases with increasing strain rate along the ND during dynamic compression progress,exhibiting the positive strain rate strengthening effect.The true stress–strain curves keep moving down at a strain rate of 1000s−1with increasing temperature,exhibiting the thermal softening effect of high temperature.

(2)The lattice constants of Mg changed by the addition of the solid solution of rare-earth elements Y and Gd inα-Mg,resulting in the difference in the critical shear stress among the basal slip and non-basal slip,decreasing the tension twinning.The decrease in thec/aaxial ratio and the increase in the high SF grain number make the non-basal slip become the dominant mechanism of plastic deformation of Mg–7Gd–5Y–1.2Nd–0.5Zr alloy.

(3)Microstructural inhomogeneity also contributes to the local accumulation of dislocations and even leads to dynamic recrystallization.As the dynamic recrystallization proceeds,the grain size is further refined and the twinning will be more difficult.Therefore,the prismatic slip and pyramidalslip become the dominant deformation mechanism of Mg–7Gd–5Y–1.2Nd–0.5Zr magnesium alloy.

Acknowledgment

Thank Professor Zhang Kui,General Research Institute for Nonferrous Metal,for providing Mg–7Gd–5Y–1.2Nd–0.5Zr magnesium alloy Sheet for this article.Furthermore,this work was supported by National Natural Science Foundation of China(Nos.51571145,51404137),City of Ningbo"science and technology innovation 2025major special project(new energy vehicle lightweight magnesium alloy material precision forming technology)(No.2018B10045).