Shuxia Ouyang,Guangyu Yang,He Qin,Chunhui Wang,Shifeng Luo,Wanqi Jie

State Key Laboratory of Solidificatio Processing,Northwestern Polytechnical University,Xi’an 710072,China

Abstract Due to the effective precipitation strengthening effect of theβ′phase,Mg-Gd alloys exhibit excellent room temperature mechanical behaviors.However,when served at high temperatures,the metastableβ′phase will transform to other phases,resulting in severe performance degradation.In this study,we investigated the effect of precipitation state achieved by different heat treatments on high temperature tensile and creep behaviors of the Mg-15Gd alloy by comparing the properties of the as-cast,solid-solutioned(T4)and peak-aged(T6)alloys.The results showed that the tensile mechanical properties of the T6 alloy were always highest from room temperature to 300 °C,in spite of an abnormal strength increase with temperature existed in the T4 alloy.For tensile creep properties,the T6 alloy exhibited the lowest steady creep rate below 235 °C while the T4 alloy possessed the best properties above 260 °C.Microstructure characterization revealed that the transition was caused by the stress-promoted precipitation ofβ′phase in the T4 alloy and rapid phase transformation in the T6 alloy at high temperatures.At 260 °C,the calculated stress exponent n was 3.1 and 2.8 for the T4 and T6 alloys,respectively,suggesting the creep deformation mechanism was dislocation slip,which was further confirme by the microstructure after creeping.Our finding can provide new insights into the heat treatment process of Mg-Gd alloys served at high temperatures.? 2021 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:Mg-Gd alloys;Precipitation state;High temperature tensile properties;Tensile creep behaviors;Stress-promoted precipitation.

1.Introduction

Magnesium alloys are known for their low density,high specifi strength and stiffness,and good castability[1].These characteristics make them being important structural materials in the automotive[2]and aerospace[3]industries.As one of the most effective strengthening methods in metallic materials,precipitation strengthening is widely used in Mg alloys.For the most mature commercial Mg-Al alloys,the high strength mainly comes from the hindering effect of Mg17Al12precipitates on basal dislocation movement[4,5].But the Mg17Al12precipitates is an unstable phase at the temperature above 175 °C[6].For the rare earth(RE)element-containing Mg alloys,their excellent high temperature mechanical properties mainly rely on the precipitation strengthening of theβ′phase under peak-aging[7].Since the under-agedβ′′phase has poor thermal stability and the over-agedβphase can be sheared easily by dislocation,the metastableβ′phase,which forms on{2ˉ1ˉ10}αprismatic planes of theα-Mg matrix and is perpendicular to the basal plane,can strengthen the alloy most effectively.However,theβ′phase is also a metastable phase.Generally,it will coarsen and transform to theβ1and further to equilibriumβphase at elevated temperatures,resulting in significan decline of the mechanical properties[8-10].Therefore,the initial precipitation state should have a critical effect on the mechanical properties of Mg-RE alloys served at high temperatures for a long time.

Although there have been many studies about the effect of precipitation state on high temperature mechanical properties of Mg-RE alloys,the conclusions are inconsistent.Wang et al.reported that the peak-aged(T6)sample showed much higher elevated temperature tensile strength than the solid-solutioned(T4)sample in the Mg-10Gd-3Y-0.5Zr alloy[11].Consistent with that,the high temperature creep resistance of the T6 state Mg-18Gd alloy was shown to be higher than the T4 alloy[12].For the hot compressed Mg-2.5Nd alloy,Li et al.also reported that aging can improve the creep resistance[13].However,Mo et al.investigated the Mg-3Gd-2Ca alloy and found that the creep resistance of the T4 alloy was better than the T6 alloy under the condition of 210°C and 100 MPa[14].Furthermore,Fang et al.claimed that the creep resistance of the T6 alloy was better below aging temperature while the T4 alloy was better above aging temperature for the Mg-4Y-2Nd-0.2Zn-0.5Zr alloy[15].These inconsistent conclusions bring difficultie to the selection of heat treatment process for Mg-Gd alloys served at high temperatures.

In previous studies,the T6 state Mg-15Gd alloy with high Gd content has been found to exhibit excellent creep resistance up to 300 °C due to the high density ofβ′phase[16].In this study,in order to clarify the effect of precipitation state on the high temperature mechanical properties,the tensile tests from room temperature to 300 °C and creep tests from 215 °C to 300 °C were performed for the as-cast,T4 and T6 Mg-15Gd alloys.The evolution of precipitates were carefully characterized to explain the difference of mechanical properties.

2.Experimental procedures

2.1.Material preparation

The experimental alloy with the nominal composition of Mg-15Gd(wt.%)was melted at 750 °C using high-purity(≥99.99%)Mg and Mg-28Gd master alloy in a mild steel crucible of an electric resistance furnace with the protection of a shielding gas of 99 vol.% CO2and 1 vol.% SF6.After being refine with Cl6C2and isothermally holding for 15 min,the melt was cast at 720 °C into a steel mold that had been preheated at 250 °C.The size of the ingots was 85 mm in diameter and 110 mm in length.The chemical composition of the prepared alloy ingot was measured by an inductively coupled plasma(ICP)analyzer,which was determined to be Mg-14.66Gd(wt.%).Based on the differential scanning calorimetry(DSC)results conducted on a universal V4.1 DTA DSC apparatus,the solid solution process was studied at 525°C for the different time.For aging treatment,the peak-aging time at 225 °C and 250 °C was determined,respectively.

2.2.Mechanical tests

Vickers hardness test was carried out using a Struers-A300 hardness tester to plot the solution and aging curves.The load was 1 kg and the dwelling time was 30 s.For each sample,at least 9 points was tested.The tensile tests were conducted on an Instron 3382 machine equipped with an environmental chamber that can be heated from room temperature(25 °C)to the pre-set temperature with the error within±2 °C.The strain rate of 1 mm/min was used and three samples were tested for each condition to ensure the reproducibility.The tensile creep tests were conducted on the CSS-3902 creep machine under the creep stresses of 50-90 MPa and the creep temperatures of 215-300 °C.Before creeping,the samples were held at the pre-set temperature for 40 min.

2.3.Microstructural characterization

The microstructure of the samples was characterized using the OLYMPUS-GX71 optical microscopy(OM),VEGA3 TESCAN scanning electron microscopy(SEM)and Themis Z double spherical aberration transmission electron microscope(TEM).For OM and SEM observation,the samples were polished and etched using a solution of 4% nitric acid and 96%ethanol.For TEM observation,thin foils were cut from the longitudinal sections of the creep samples and mechanically grounded to a thickness of 50μm.Then,they were punched into 3 mm diameter disks and further ion milled in a precision ion polishing system(PIPS)operating at 4 kV accelerating voltage and 6° incident angle.

The precipitate size was define using the width measured from TEM.The average precipitate size was calculated according to precipitate size distributions.The volume fraction ofβ′phase was measured by the Image-Pro Plus software.The precipitate number density is the areal density of precipitates.At least three TEM images were used to determine the precipitation size,volume fraction,and number density.

3.Results and discussion

3.1.Heat treatment process and microstructure evolution

3.1.1.Solution treatment and age hardening response

Fig.1(a)shows the DSC heating curve of the as-cast Mg-15Gd alloy.The secondary phase began to dissolve at 544.3°C and ended at 554.6 °C.The position of the endothermic peak indicates that the eutectic temperature of the as-cast alloy is 549.4°C.Accordingly,the optimal solution temperature of 525°C was chosen,which is a little lower than the eutectic temperature.Fig.1(b)presents the variation of micro-hardness of the as-cast Mg-15Gd alloy with solution time at 525 °C.Due to the existence of Mg5Gd phase formed during solidificatio[17],the hardness of the as-cast alloy was highest(? 86.3 HV).The hardness decreased with time and reached stable at?76.8 HV after 8 h.Therefore,the most suitable solid solution condition of the alloy was confirme to be 525 °C for 8 h.

Due to the high Gd content,the Mg-15Gd alloy can precipitate a large amount ofβ′phase.Fig.2 shows the age hardening curves of the solution treated(525 °C,8 h)T4 alloy at 225 °C and 250 °C,respectively.For each aging temperature,the hardness increased gradually to a peak value,showing a wide peak of hardness plateau(PHP),and then decreased slowly due to the over aging.Compared with 250°C,the hardness increased more slowly when aging at 225°C,and took a longer time(?12 h)to reach the peak value,but the peak-aging hardness was higher(?134.2 HV).The detailed data are summarized in Table 1.Therefore,the most suitable aging condition was determined to be 225 °C for 12 h.Compared with the as-cast and T4 state,the hardness of the peak-aged T6 alloy increases 55.5 % and 44.1 %,respectively,indicating the significan precipitation behavior.

Fig.1.(a)DSC heating curve of the as-cast Mg-15Gd alloy;(b)solid solution curve of the as-cast Mg-15Gd alloy at 525 °C.

Fig.2.Age hardening curves of the T4 state Mg-15Gd alloy at 225 and 250 °C.

Table 1Peak-hardness,PHP time and peak aging time of the T4 alloy aged at 225 and 250 °C.

3.1.2.Microstructure evolution during heat treatment

The microstructures of the as-cast,T4 state and T6 state alloys are shown in Fig.3.As shown in Fig.3(a),the as-cast Mg-15Gd alloy showed dendrite morphology with the secondary phase in the interdendritic regions.The inset shows some eutectics(labeled as A),polygon-shaped(labeled as B)and cuboid-shaped(labeled as C)precipitates coexisted in the Mg matrix.EDS results of different precipitates are given in Table 2.According to the atomic ratio of Mg/Gd,the eutectics consisted of theα-Mg matrix and Mg5Gd.In point B,the atomic ratio of Mg and Gd in the polygon-shaped precipitates was close to 5:1,which can be determined to be Mg5Gd.The cuboid-shaped precipitations(point C)were rich in the Gd element.Combined with previous reports,they were the GdH2phase formed during melting[18].Fig.3(b)shows the microstructure of the T4 alloy after solution treatment.Almost all the secondary phase dissolved into theα-Mg matrix with only a bit cuboid-shaped phase and polygon-shaped precipitates left[19].The polygon-shaped and cuboid-shaped precipitates can still be observed in the T6 state alloy,as shown in Fig.3(c).According to the SAED patterns observed from[20]αzone axis in Fig.3(d),three additional diffraction spots exited at the position of 1/4{010}α,1/2{01ˉ10}α,3/4{01ˉ10}α,indicating the formation ofβ′phase with a basecentred orthorhombic structure(a?2aMg=0.64 nm,b=2.23 nm,c?cMg=0.52 nm).the TEM bright image of the cuboidshaped phase and corresponding SAED patterns are shown in Fig.3(e),indicating that the fcc crystal structure with a lattice parameter of a=0.53 nm.Therefore,the cuboid-shaped precipitations can be determined to be GdH2[18].The XRD patterns are shown in Fig.3(f),both the as-cast,T4 and T6 alloys showed the peak of Mg5Gd phase,confirmin the existence of Mg5Gd in these alloys.Combined with the EDS results of Mg:Gd close to 5:1,the polygon-shaped precipitations can be determined to be Mg5Gd.Therefore,eutectics,Mg5Gd and GdH2can be observed in the as-cast alloy.GdH2besides the undissolved Mg5Gd phase existed in the T4 alloy.The T6 alloy was composed of elliptical shapedβ′phase,cuboid-shaped GdH2phase and a small amount of undissolved Mg5Gd phase.

Table 2EDS results of the different precipitates of T6 state alloy in Fig.3(a)-(c).

Fig.3.Microstructure of the Mg-15Gd alloys with different precipitation states.(a),(b)and(c)OM and SEM images of the as-cast,T4 and T6 alloys,respectively.(d)TEM bright fiel image of T6 alloy and corresponding SAED patterns(Z=[2ˉ1ˉ10]α)(theα-Mg was marked by red circles,and theβ′phase was marked by yellow circles).(e)TEM bright fiel image of the cuboid-shaped phase and corresponding SAED patterns(Z=[001])(the cuboid-shaped phase was marked by yellow circles).(f)XRD patterns of the as-cast,T4 and T6 Mg-15Gd alloys,respectively.

3.2.Tensile behaviors

3.2.1.Room and high temperature tensile properties

We firs investigated the effect of precipitation state on the room and high temperature tensile properties of the Mg-15Gd alloy.Fig.4 shows the room temperature tensile mechanical properties of the alloys in three states.Compared with the as-cast specimen,solution treatment resulted in a significan improvement of ductility by eliminating the brittle phase and a slight decrease in both ultimate tensile strength(UTS)and yield strength(YS).Further aging increased the UTS and YS to?183.3 MPa and?186.6 MPa,respectively,while reducing the ductility.

Fig.4.Room temperature tensile mechanical properties of the Mg-15Gd alloy in three states.

The variation in tensile mechanical properties with the temperature of the as-cast,T4 and T6 alloys is shown in Fig.5(a)-(c),respectively.Both the YS and UTS decreased monotonically from room temperature to 300 °C for the ascast and T6 alloy.However,for the T4 alloy,an abnormal increment of strength with the temperature existed between 235 and 285°C.Considering the continuous and slow heating process during high temperature test,the strength increase can be caused by the dynamic precipitation in the T4 alloy.While for the three alloys,the ductility increased with temperature.Despite the abnormal strength increase in T4,it should be noted that the tensile strength of the T6 was highest throughout the test temperature range.Fig.5(d)compares the tensile properties of the three alloys at 260 °C intuitively.Therefore,the T6 alloy is the preferred material for high temperature and short time service.

Fig.5.Variation in the tensile properties(YS,UTS and elongation)with temperature of the as-cast(a),T4(b)and T6(c)alloys.(d)Comparison of the tensile properties of the three alloys at 260 °C.

3.2.2.Tensile fracture mechanisms

Representative fracture features of the three alloys are shown in Fig.6.At room temperature,fracture surfaces of the as-cast alloy were mainly composed of cleavage planes and a small amount of coarse dimples,as shown in Fig.6(a),indicating the primary brittle transgranular fracture.For the T4 alloy,the increased ductility is reflecte by the increased dimples and decreased cleavage planes,as shown in Fig.6(b).While for the T6 alloy,there was a large number of cleavage planes on the fracture surfaces shown in Fig.6(c),exhibiting the typical brittle fracture features.The transition of fracture mode from brittle to ductile with temperature increasing is a common phenomenon in metals.In this study,the fraction of dimples increased for the as-cast and T6 alloys,and the size of dimples became larger in the T4 alloy at 260 °C,accounting for their increased ductility at high temperatures.

3.3.Tensile creep behaviors

3.3.1.Creep properties

Tensile creep curves of the as-cast,T4 and T6 alloys tested at 260 °C under 50 MPa,70 MPa and 90 MPa are shown in Fig.7(a)and(b).For all curves,typical short creep primary stage and long secondary stage existed.During the primary stage,the creep rate declined rapidly from the initial high level within 10 h.Then,the secondary stage of steady creep stage with low creep rate continued until 100 h at 50 and 70 MPa.While at 90 MPa,the tertiary creep stage appeared and the as-cast,T4,T6 alloys ruptured at 75,91,85 h,respectively.Comparing the three states,we found that the T4 alloy showed the best creep resistance and the as-cast alloy showed the worst creep resistance at this temperature,independent of the applied stress.This result is inconsistent with the tensile strength where the T6 alloy showed the best high temperature mechanical properties.

To further prove the effect of temperature on creep properties of the three alloys,we performed the creep tests at 215,235 °C(shown in Fig.7(c))and 260,285,300 °C(shown in Fig.7(d))at constant stress of 50 MPa.The creep curves are similar to the results in Fig.7(a)and(b).However,obvious temperature-dependent creep properties of the three alloys can be observed,i.e.,the T6 alloy showed better creep resistance below 235 °C while the T4 alloy showed better creep resistance above 260 °C.The as-cast alloy showed the worst properties at all temperatures,indicating the weakest strengthening effect of the blocky Mg5Gd phase.By calculating the derivative of creep strain with respect to creep time,Fig.7(e)and(f)further compared the creep rate of the three alloys tested at 215 °C and 300 °C,respectively.Despite the minimum creep rate of the T6 alloy was much lower than that of the T4 alloy at 215 °C,the T4 alloy showed lower minimum creep rate at 300 °C.The detail creep data tested at 260 °C with different stresses and at 50 MPa with different temperatures was calculated and summarized in Tables 3 and 4.

Fig.6.SEM images showing room temperature tensile fracture surfaces of the as-cast(a),T4(b),T6(c)alloys,and high temperature(260 °C)tensile fracture surfaces of the as-cast(d),T4(e),T6(f)alloys.

Fig.7.Tensile creep curves of the as-cast,T4 and T6 alloys tested at constant temperature of 260 °C under applied stresses of 50 MPa,70 MPa(a)and 90 MPa(b).Tensile creep curves of the as-cast,T4 and T6 alloys tested at constant stress of 50 MPa under 215 °C,235 °C(c)and 260 °C,285 °C,300 °C(d).Logarithmic scale of creep rate versus time under 50 MPa at 215 °C(e)and 300 °C(f).

Table 3.Creep data of the as-cast,T4 and T6 Mg-15Gd alloy tested at 260 °C and different stresses.

Table 4.Creep data of the as-cast,T4 and T6 Mg-15Gd alloy tested under 50 MPa and different temperatures.

3.3.2.Creep mechanisms

The steady creep rate()can be described by the conventional power-law equation in term of the applied stress(σ)and the temperature(T):[20,21]

where A is a constant,nis the stress exponent,Qis the activation energy for creep.Therefore,the stress exponentncan be obtained from the slope of the log-log plot of the minimum creep rate versus the applied stress,and the activation energyQcan be obtained from the slope of the logarithm of the creep rate versus the reciprocal of temperature[22],as shown in Fig.8(a)and Fig.8(b),respectively.By fittin the points in Fig.8(a),the stress exponent of the as-cast,T4 and T6 alloys was determined to be 4.1,3.1,and 2.8 at 260°C.According to the deformation mechanism map[23,24],the creep mechanisms were presumed to be dislocation climb for the as-cast alloy and viscous gliding of dislocation for both the T4 and T6 alloys.In Fig.8(b),the activation energy Q of the as-cast and T4 alloys changed obviously below and above 260°C.It was calculated to be 50 kJ/mol below 260°C and 210 kJ/mol above 260 °C for the as-cast alloy indicating the transition of creep mechanism from grain boundary slide to cross slip,and 33 kJ/mol below 260 °C and 141 kJ/mol above 260 °C for the T4 alloy,indicating the transition of creep mechanism from grain boundary slide to diffusion of Gd in Mg,This transition can be attributed to the dynamic precipitation during creep at high temperature,as discussed in Section 3.3.3.While for the T6 alloy,the value of Q is 118 kJ/mol at low temperature and 141 kJ/mol at high temperature,which indicates the creep mechanism transition from pipe diffusion to diffusion of Gd in Mg.

Fig.9(a)presents the TEM bright fiel image of T4 alloy crept at 260°C and 50 MPa for 100 h,with the incident beam direction parallel to[011]α.The characters of dislocations were determined according to the contrast.All of dislocations were visible in Fig.9(a),while the dislocations marked by red arrows were out of contrast under the two-beam condition ing1=[101]αandg2=[012]α,as shown in Fig.9(b)and(c).According to theg·b=0 invisibility criterion,the burgers vectorbof the invisible dislocations was determined to be(1/3)[ˉ12ˉ13][25].Therefore,the long and straight dislocations were distinguished as＜a+c＞dislocations,indicating the creep mechanism of pyramidal slip.

3.3.3.Microstructure evolution during creep

To reveal the abnormal better creep properties of the T4 alloy than the T6 alloy at high temperature,the microstructure evolution during creep of the T4 alloy was characterized.Fig.10 shows the TEM images of the T4 alloy crept at 260°C under 50 MPa at 10 and 20 h.TEM dark fiel image of the sample crept for 10 h is shown in Fig.10(a).An amount of thin platelet precipitates with 60-100 nm in length and 10-15 nm in thickness(aspect ratio about 6:1)were found.With the electron beam parallel to[010]α,the corresponding SAED patterns are shown in Fig.10(b).From the diffraction patterns and morphology,the precipitations were identifie as theβ′′phase giving rise to the faint spots at 1/2{110}α.Theβ′′phase,which is metastable and fully coherent with the matrix,has an ordered DO19type structure(a?2aMg=0.64 nm,c?cMg=0.52 nm).The orientation relationship of theβ′′precipitates and theα-Mg matrix was identifie to be[010]β′′//[010]α,(20)β′′//(20)α,and the habit plane was parallel to＜0001＞direction[9,26].Since the structure ofβ′′phase is similar to that ofα-Mg,the nucleation process ofβ′′phase can be completed by replacing the Mg atoms periodically with rare-earth atoms without structural changes.TEM bright fiel image of the sample crept for 20 h is shown in Fig.10(c).The lens-shaped precipitations with the size of 100-120 nm in length and 20-30 nm in thickness appeared.According to the SAED patterns observed from[11]αzone axis in Fig.10(d),three additional diffraction spots exited at the position of 1/4{010}α,1/2{010}α,3/4{010}α,indicating that those precipitates wereβ′phase[27].Theβ′phase has a base-centred orthorhombic structure(a?2aMg=0.64 nm,b=2.23 nm,c?cMg=0.52 nm).The orientation relationship between theβ′precipitates and theα-Mg matrix was found to be[001]β′//[0001]α,(010)β′//(010)α,and the habit plane ofβ′with three variants were parallel to the＜010＞αdirection[28,29].Since the structure ofβ′phase with less rare-earth element content is similar to that ofβ′′phase,the transformation fromβ′′toβ′can be achieved by decreasing rare-earth atoms.This process is accompanied by the increase of precipitation volume fraction.

Fig.8.(a)Double logarithmic plot of the minimum creep rate versus applied stress at 260 °C.(b)Logarithmic of the minimum creep rate vs the reciprocal absolute creep temperature under 50 MPa.

Fig.9.(a)TEM bright fiel image of the T4 alloy crept at 260 °C and 50 MPa for 100 h,with the incident beam direction z=[011]α(theα-Mg was marked by red circles).(b)and(c)Dislocations marked by red arrows in(a)are invisible under two-beam condition with g1=[101]αand g2=[012]α,respectively.

Fig.11(a)shows the TEM bright fiel image of the T4 alloy crept at 260 °C under 50 MPa for 50 h.Both small and large plate-shaped precipitates coexisted in the matrix.According to the morphology,the large plate-shaped precipitate was confirme as theβ1precipitates,as indicated by Nie and Muddle[30].Theβ1precipitate was reported to have a fcc structure with a lattice parameter of a=0.74 nm,and the orientation relationship betweenβ1phase with theα-Mg matrix was[110]β1//[0001]α,(112)β1//(1100)α,and the habit plane was parallel to the＜110＞direcltion.Fig.11(b)presents the TEM bright fiel image containing small plate precipitates with two directions.According to the SAED patterns observed from[43]αin Fig.11(c),they were theβ′phase.Theβ1phase formed between twoβ′phases.Along with the formation of theβ1phase,theβ′phase diminished gradually.However,it is still controversial for the formation mechanism of theβ1phase.[31,32].Therefore,after creep for 50 h,part ofβ′phase disappeared while others only coarsened.Fig.11(d)shows the TEM bright fiel image of the T4 alloy crept at 260 °C under 50 MPa for 100 h,which shows the similar morphology with Fig 11(a).The large plate-shaped precipitates were stillβ1phase and the small plates wereβ′phase.High-magnificatio image in Fig.11(e)and SAED patterns in Fig.11(f)further confirme the morphology and crystal structure of theβ′phase.However,the size of bothβ1phase andβ′phase coarsened and the fraction of theβ1phase increased,indicating that moreβ′disappeared withβ1phase formed.Based on the above analysis,we found a slower phase transformation in the T4 alloy than the T6 alloy under the same creep condition.For the T6 alloy,theβ′phase disappeared completely within 100 h[16].Since it is the mainly strengthening phase in this alloy system,the rapid transformation resulted in the worse creep resistance of the T6 alloy.On the other hand,for the T4 alloy,theβ′phase formed during the primary creep stage and provided effective strengthening effect during the entire secondary creep stage.The key strengthening phase isβ′in the Mg-15Gd alloy,which plays an important role on the creep resistance.Since the coherentβ′phase forms on the prismatic planes{20}αof the matrix in a dense triangular arrangement,the three-dimensional shape of the precipitates resembled a convex lens,which is perpendicular to the basal plane.When the alloy is plastically deformed,a complete lattice translation in theβ′crystal necessitates two＜a＞dislocations,generating great anti-phase boundary energy[33].So theβ′phase can act as effective barrier to dislocation motion on the basal plane.As for theβ1phase,it is much larger in size is incoherent with the Mg matrix,making dislocations tend to bypass rather than cut through it[34].Consequently,the strengthening effect ofβ1phase is much weaker thanβ′.

Fig.10.TEM images showing the morphology and distribution of precipitates in the T4 state Mg-15Gd alloy crept at 260 °C under 50 MPa for 10 h(a)and 20 h(c).(b)and(d)show corresponding SAED patterns taken from Z=[01ˉ10]αand Z=[11]α,respectively,(theα-Mg was marked by red circles,and the precipitate phase was marked by yellow circles).

Fig.11.TEM images showing the morphology and distribution of precipitates in the T4 state Mg-15Gd alloy crept at 260 °C under 50 MPa for 50 h(a),(b)and 100 h(d),(e).(c)and(f)show corresponding SAED patterns taken from Z=[ˉ24ˉ23]αand Z=[11]α,respectively(theα-Mg was marked by red circles,and theβ′phase was marked by yellow circles).

We further investigated the effect of creep on the content of precipitations.Since theβ′phase is the strengthening phase,we compared the volume fraction ofβ′in the T4 and T6 alloys.As shown in Fig.12(a),the phase fraction ofβ′was 18.5% in the T6 alloy before creeping,representing for the peak aging effect without stress.For the T6 alloy crept at 260 °C under 50 MPa for 5,25,50 h,where over-aging proceeded under stress,theβ′phase fraction was 19.6,18.4,14.4%,respectively[16].However,for the T4 alloy crept at 260 °C under 50 MPa for different times,where aging proceeded under stress,theβ′phase fraction was 30.6,27.5,23.4% at 20,50,100 h.The increased fraction ofβ′phase indicated a stress-promoted precipitation in the T4 alloy during creeping,which is also conducive to the improvement of high temperature creep performance.Fig.12(b)shows the number density of theβ′phase with creep time in the T4 and T6 alloys.For the T4 alloy,the number density ofβ′was highest at 20 h and decreased gradually with creep time,which is due to the phase transformation toβ1.Similarly,for the T6 alloy,the number density ofβ′was highest before creeping and decreased with creep time.It is worth mentioning that the number density of the T4 alloy was always higher than that of T6.The average size variation of theβ′phase in the T4 and T6 alloys is shown in Fig.12(c).The precipitates in the T6 alloy coarsened faster and their average size was larger than that in T4 at all time.Fig.12(d)and(f)further show the size distribution ofβ′with creep time in the T4 and T6 alloys,respectively,which also revealed the larger precipitation size in the T6 alloy.The dynamic precipitation of T6 alloy is completed.Therefore,it is reasonable to infer that the aging during creep induced the formation of more and fine precipitations.The Vickers hardness of the ascast,T4 and T6 alloys before and after creeping at 260 °C under 50 MPa is shown in Fig.12(f).It shows that the hardness of the as-cast sample increased 12.3 HV after creeping,which is related to the slight precipitation.The hardness of the T4 state sample increased 33.7 HV after creeping.Since the creep strain of the sample is less than 0.1%,the effect of dislocation on hardness change can be ignored.Therefore,the hardness increment can be attributed to the higher density of precipitates formed during creeping.However,for the T6 alloy,the hardness decreased slightly from 134.2 to 104.1 HV.This is caused by the disappearance of theβ′phase.Theoretically,the T6 alloy which was peak aged at 225 °C should have more precipitations than the T4 alloy creeping at a higher temperature of 260 °C from the perspective of thermodynamics.However,as shown in Fig.12,the T4 alloy after creeping for 10 h showed the higher density,higher fraction and fine precipitations than the T6 alloy without creeping.Since the temperature and applied stress during creeping is the only two variables,we can attribute the higher density and fine precipitations in the T4 alloy to be stress-promoted.

Fig.12.Evolution of(a)the volume fraction,(b)the number density,(c)the average size in the T4 and T6 Mg-15Gd alloys,(d)size distribution in the T4 alloy,(e)size distribution in the T6 alloy,of theβ′phase with creep time in the Mg-15Gd alloys.(f)Vickers hardness of the as-cast,T4 and T6 alloys before and after creeping at 260 °C under 50 MPa for 100 h.

To conclude,moreβ′phase formed in the T4 alloy during creeping at high temperature and existed in the whole creep test,leading to the higher hardness and better creep resistance than the T6 alloy.As for T4 alloy creeping at low temperature,the precipitation kinetics is so slow that theβ′phase cannot precipitate to provide effective strengthening.Therefore,the T6 alloy showed better creep properties at low temperature.This result can provide guidance for the selection of heat treatment of the Mg-15Gd alloy served at high temperature for a long time.The T6 treatment is recommended for serving below 235 °C,while the T4 treatment is recommended for serving above 260 °C.

4.Conclusions

In this study,the effect of precipitation state on high temperature tensile and creep behaviors of the Mg-15Gd alloy was investigated by comparing the properties of the as-cast,solid-solutioned(T4)and peak-aged(T6)alloys.The main finding are summarized below:

(1)The suitable solid solution condition of the alloy was confirme as 525 °C for 8 h.The optimum aging heat treatment was confirme as 225 °C for 12 h.

(2)The tensile properties of the T6 alloy were always highest from room temperature to 300 °C,in spite of an abnormal strength increase with temperature existed in the T4 alloy.The fracture mode transitioned from brittle to ductile with temperature increasing in the Mg-15Gd alloy.

(3)For creep properties,the T6 alloy exhibited the lowest steady creep rate below 235 °C while the T4 alloy possessed the best properties above 260 °C.The stresspromoted precipitation of theβ′phase contributed to the better creep resistance of the T4 alloy at high temperature.

(4)The stress exponent n of the T6 and T4 alloys crept at 260 °C was calculated to be 2.8 and 3.1,indicating the creep deformation of dislocation glide.This was further confirme by TEM analysis

Acknowledgments

This work was supported by the National Natural Science Foundation of China(Grant No.51771152),the National Key Research and Development Program of China(Grant No.2018YFB1106800).