A rolled Mg-8Al-0.5Zn-0.8Ce alloy with high strength-ductility synergy via engineering high-density low angle boundaries

Xio M,Min Zh,b,c,?,Siqing Wng,Yi Yng,Hilong Ji,c,?,Dn Go,Cheng Wng,b,

Huiyuan Wanga,b,c

a Key Laboratory of Automobile Materials of Ministry of Education & School of Materials Science and Engineering,Nanling Campus,Jilin University,No.5988 Renmin Street,Changchun 130025,PR China

b State Key Laboratory of Super hard Materials,Jilin University,Changchun,130012,PR China

cInternational Center of Future Science,Jilin University,Changchun 130012,PR China

Abstract Developing low-cost rolled Mg alloys with both high strength and ductility is desirable,while the improved strength is generally accompanied with decreased ductility.Here,by using rotated hard-plate rolling(RHPR)with a total thickness reduction of～85%,we obtained a Mg-8Al-0.5Zn-0.8Ce(wt.%,AZ80-0.8Ce)alloy with a high strength-ductility synergy,i.e.,the yield strength(YS),ultimate tensile strength(UTS)and elongation-to-failure(EF)are～308 MPa,～360 MPa and～13.8%,respectively.It reveals that the high YS is mainly originated from grain boundary strengthening(～212 MPa),followed by dislocation strengthening(～43 MPa)and precipitation hardening(～25 MPa).It is found that a relatively homogeneous fine grain structure containing a large fraction(～62%)of low angle boundaries(LABs)is achieved in the RHPRed alloy,which is benefit for the high tensile EF value.It demonstrates that LABs have important contributions to strengthening and homogenizing tensile deformation process,leading to the simultaneous high strength and high EF.Our work provides a new insight for fabrication of low-cost high performance Mg alloys with an excellent strength-ductility synergy.

Keywords:Mg-Al-Zn alloys;Rolling;Strengthening mechanism;Microstructure;Strength-ductility synergy.

1.Introduction

For growing demand of environment protection and energy saving,magnesium(Mg)alloys have attracted much attention for their low density,high specific strength and ease of recycling[1–3].Among commercial Mg alloys,Mg–Al–Zn(AZ)series,such as AZ80 alloy,are widely used[4–9]due to relatively good mechanical properties and low costs.However,further improving the strength generally accompanied by the loss of ductility.For instance,an AZ80 alloy fabricated via pre-aging and rolling shows a high strength(～401 MPa),while a limited elongation-to-failure(EF,～4.8%)value is presented[10].Therefore,developing low-cost wrought Mg alloys with an excellent strength-ductility synergy is in urgent demand,which is of significance to extend applications of Mg alloys.

Up till now,many attempts have been made to further improve mechanical properties of Mg alloys.Adding rare earth(RE)elements has been considered as an effective way to improve the ductility of Mg alloys[11–13].Among available RE elements,Ce is a promising candidate because of the relatively low cost and strong texture weakening ability[14].It has been demonstrated that the addition of Ce can reduce the relative value of critic resolve shear stress(CRSS)for prismatic and pyramidal〈c+a〉dislocations of pure Mg,leading to a superior EF(～38%)of the Mg-0.2Ce(wt.%)alloy fabricated by equal channel angular processing(ECAP)[15].Moreover,the addition of Ce into Mg–Al–Zn(AZ)series alloys will form large Al-Ce particles,which is benefit for dynamic recrystallization(DRX)through the particlestimulated nucleation(PSN)mechanism.For example,in a previous work[16],the extruded AZ80-0.8Ce alloy exhibited high mechanical properties due to the reduction of grain size via the PSN mechanism.

In addition to RE alloying,microstructure regulation is also an important way for improving mechanical properties[17–19].Recently,bimodal-grain structured Mg alloys have attracted attentions due to the good combination of strength and ductility.In our previous work,hard-plate rolling(HPR)has been developed to improve the formability of Mg alloys,avoiding the formation of cracks that easily appear during traditional rolling.By single-pass HPR,Wang et al.[19]have prepared a bimodal-grained AZ91,which is characterized by DRXed fine grains(FGs)and deformed coarse grains(CGs).The bimodal-grained AZ91 shows a high strength-ductility synergy,i.e.,an ultimate tensile strength(UTS)of～371 MPa and a EF of～26%.This is mainly due to the high work hardening ability originating from the easy activation of basal slip in FGs and the strong dislocation storage ability in CGs[20].However,due to the easy activation of basal〈a〉slips in randomly oriented FGs[20],HPRed AZ series alloys generally show low yield strength(YS＜250 MPa)[19–21].

For bimodal grain structures in Mg alloys,most studies have focused on traditional AZ series alloys without RE elements.For the bimodal-grained AZ-RE alloys,stress concentration easily occurs in front of large Al-RE particles,eventually leading to an early fracture[22].To achieve a high strength-ductility synergy in Mg alloys containing coarse Al-RE particles,it is better to relieve the stress concentration around coarse Al-RE particles.In other words,a relatively homogeneous deformation process is necessary.In the present work,a new rolling strategy named rotated hard-plate rolling(RHPR)has been applied on an extruded AZ80-0.8Ce alloy.The main purpose of this work is to explore the underlying strengthening mechanisms and reasons for enhanced EF of the RHPRed AZ80-0.8Ce alloy based on microstructure characterization.The results will be helpful in fabricating low-cost Mg alloys with a high strength-ductility synergy.

2.Experimental procedures

Commercial pure Mg(99.85 wt%),Al(99.90 wt%),Zn(99.90 wt%)and Mg-28 wt% Ce master alloy were used to prepare AZ80-0.8Ce cast rod.Chemical compositions(Table 1)were measured by an optical spectrum analyzer(ARL 4460 Switzerland).After homogenized at 430 °C for 3 h,the rod with a length of 400 mm was extruded at 430 °C to sheets with a section of 40 mm×5 mm.Rolling samples were cut from as-extruded sheets,which were solid solution treated at 415 °C for 20 h in an air circulation furnace followed by water-quenching to room temperature.For comparison,two rolling procedures with the same total reduction(～85%)were carried out.One is the single-pass HPR along transverse direction(TD)of as-extruded alloys,which features two hard plates covering up and down of rolling samples,like a sandwich structure(Fig.1(a)).The other is RHPR including four-pass HPR process,where the first pass is in parallel to TD of extruded plates while the latter pass is perpendicular to the previous pass.The thickness reduction of the four passes is～55%,～40%,～30% and～15%,respectively.Corresponding schematic illustration is shown in Fig.1(b),where black arrows represent rolling directions.For each rolling pass,samples were pre-heated at 250 °C for 15 min,and rollers were pre-heated to 100 °C.The rolled sheets were annealed at 175 °C for 5 min to relieve the residual inter-stress for subsequent microstructure observations and tensile tests.

Table 1Chemical compositions of the AZ80-0.8Ce alloy.

Table 2Tensile properties of the extruded,HPRed and RHPRed AZ80-0.8Ce alloy.

Microstructure characterization was conducted using an optical microscope(OM,Carl Zeiss-Axio Imager A2m,Germany),a field emission scanning electron microscope(FESEM,Sigma 500)equipped with an energy disperse spectrometer(EDS)and an Oxford instruments Symmetry electron backscatter diffraction(EBSD)detector,and a transmission electron microscope(TEM,JEM-2100F,Japan)operating at an accelerating voltage of 200 kV.Phases were detected by X-ray diffraction(XRD,18KW/D/Max 2500 PC,Rigaku,Japan)with Cu Kαradiation at a voltage of 40 kV and a scanning speed of 4° min-1.The low angle boundary(LAB:2°≤θ≤15°)and high angle boundary(HAB:θ≥15°)are presented as white lines and black lines,respectively.Metallographic samples for OM and SEM observations were ground with sandpapers and mechanically polished with 0.5 μm diamond paste,followed by chemical etching in an acetic picric solution for about 10-30 s.The EBSD samples were mechanically polished and then electro-polished with a solution containing Citric acid and Perchloric acid at 20 V for 60-120 s under-30 °C.The average size of grains and particles were measured using linear intercept method by software Nano Measurer System 1.2.Tensile test samples with a gage size of 10×4×1.2 mm3were cut along rolling direction(RD)of HPRed and RHPRed sheets in the RD-TD section.Uniaxial tensile tests were conducted at room temperature under a strain rate of 1.0×10-3s-1,using a INSTRON 5896 testing machine.At least 3 samples were tested for each condition to ensure the reliability of results.

3.Results

3.1.Microstructure of the as-extruded AZ80-0.8Ce alloy

Microstructure of the as-extruded AZ80-0.8Ce alloy is shown in Fig.2.It is clear that most grains are equiaxed(Fig.2(a))with an average size of～32 μm,as a result of DRX during hot extrusion[16].At the same time,two typesof particles can be observed in the higher magnification secondary electron(SE)image(Fig.2(b)).Irregular strip shaped secondary phase are located at GBs,while the bulk shaped particles with sizes of 2-5 μm are within grains and at GBs.From corresponding EDS images in Fig.2(c),it can be seen that the Al element is present in both types of secondary phases while Ce element only concentrates in the bulk shaped particles.The XRD analysis in Fig.2(d)shows that there are three kinds of phases in the as-extruded AZ80-0.8Ce alloy,i.e.,primaryα-Mg,β-Mg17Al12and Al4Ce.Therefore,the irregular strip shaped and bulk shaped particles in Fig.2(b)can be confirmed asβ-Mg17Al12and Al4Ce,respectively.In addition,as shown in the(0002)pole figure(Fig.2(f)),the as-extruded sheet exhibits a basal texture character,i.e.,thec-axes of most grains tilt by～45° from the ND to the ED,with a maximum intensity of～19 mr.

Fig.1.Schematic representation of(a)HPR,(b)RHPR process.

Fig.2.(a)and(b)SE images;(c)EDS analysis corresponding to(b);(d)XRD analysis;(e)EBSD IPF map of the as-extruded AZ80-0.8Ce alloy and(f)corresponding(0002)pole figure.

The optical image of AZ80-0.8Ce alloy after solid solution treatment are presented in Fig.3(a).The average grainsize is measured to be～44 μm,which shows a slightly increment compared with that of as-extruded alloy.At the same time,most Mg17Al12particles along GBs have dissolved into theα-Mg matrix.Whereas the number density,sizes and morphology of Al4Ce particles are similar in the as-extruded and as-solutionized alloys(Fig.3(b))due to the high melting point of Al4Ce phase[23].This is consistent with the XRD results presented in Fig.3(c),which shows that only two phases exist in the as-solutionized sample,i.e.,primaryα-Mg and Al4Ce.The EDS analysis(Fig.3(d))corresponding to Fig.3(b)also confirms that the Mg17Al12particles are fully dissolved while the Al4Ce phase is preserved.

Fig.3.(a)optical and(b)SE images;(c)XRD analysis of AZ80-0.8Ce alloy after solid solution treatment;(d)EDS analysis corresponding to(b).

3.2.Microstructure of the RHPRed and HPRed AZ80-0.8Ce alloy

EBSD results of RHPRed and HPRed alloys are presented in Fig.4.It can be seen that both rolled alloys exhibit a basal texture with basal poles nearly aligning to ND.The HPRed alloy shows a higher maximum texture intensity(～22 mr)compared to that of the as-extruded alloy(～19 mr).However,the RHPRed alloy shows a lower maximum texture intensity(～17 mr).As indicated in Fig.4(a)and(d),the HPRed AZ80-0.8Ce alloy exhibits a typical bimodal grain structure[20,21],whereas the RHPRed alloy shows a relatively homogeneous fine grain structure.Note that in this work,grains with sizes smaller than 4 μm are defined as FGs and the others are CGs.The grain size distribution histograms(Fig.4(g)and(h))corresponding to Fig.4(a)and(d)reveal that the HPRed alloy shows a bimodal-grained structure consisting of～42% FGs and～58% CGs.In contrast,a large area fraction(more than 70%)of grains are FGs in the RHPRed alloy.The average grain size is～1.5 μm of FGs and～15 μm of CGs for the HPRed alloy,and～2.3 μm for the RHPRed AZ80-0.8Ce alloy.To further identify the characteristics of rolled microstructure,the selected areas with a higher magnification are illustrated in Fig.4(b)and(e).A distinct feature is that most of the LABs exist mainly in CGs of HPRed alloy,whereas LABs are uniformly distributed in the RHPRed alloy.The LAB length per unit area is calculated to be～0.28 μm-1for FGs and～0.56 μm-1for CGs in HPRed alloy,while it is～1.22 μm-1in RHPRed alloy.The fraction of LABs are～32% for the HPRed alloy and～62% for the RHPRed alloy,respectively.KAM maps corresponding to Fig.4(b)and(e)are illustrated in Fig.4(c)and(f),respectively.These results are in accordance with the distribution of LABs,where the CGs in HPRed alloy and most grains in the RHPRed alloy feature large misorientation gradients.

The SE images of rolled alloys are presented in Fig.5.As shown in Fig.5(a)and(d),the Al4Ce particles are similar in size,morphology and distribution with those in the as-extruded alloy.The number density of the Al4Ce particles of both rolled alloys are estimated to be～1.6×104m-2.As indicated,area A and area B represent grains larger than 4 μm and smaller than 4 μm,respectively.The Mg17Al12particles within area A in RHPRed alloy(Fig.5(b))exhibits an ellipse shape with length and width in the range of～90-300 nm and～40-110 nm,respectively.However,nearly no Mg17Al12particles can be observed within area A in HPRed alloy(Fig.5(e)),which is different to the formation of ellipse shaped Mg17Al12particles in the CG area of HPRed AZ91 alloy[19].This may due to the formation of Al4Ce phase,which consumes most of the Al atoms.In contrast,the RHPRed featuring four passes rolling leads to a uniform precipitation of Mg17Al12particles in the whole matrix.The Mg17Al12particles in area B of RHPRed alloy are spherical with larger sizes and a lower number density(Fig.5(c))compared to that in HPRed alloy(Fig.5(f)).According to the distribution histograms illustrated in the upper right corners of Fig.5(c)and(f),the average size and number density of Mg17Al12particles are～230 nm and～2.2×1012m-2for the RHPRed alloy and～120 nm and～7.2×1012m-2for the HPRed alloy.Due to the formation of Al4Ce particles,the size of Mg17Al12precipitates in present study is smaller compared with that in a HPRed AZ91 alloy(～600 nm)[24].The larger size and lower number density of Mg17Al12particles in RHPRed alloy is mainly due to the thermodynamic coupling between multiple rolling passes and intermediate annealing during RHPR process,where dynamic precipitation and redissolution of Mg17Al12particles are both present.The particles with irregular shape,which dynamically precipitate during rolling deformation tended to dissolve into matrix during intermediate annealing because of their relatively poor thermal stability[25].As a result,spherical shape Mg17Al12precipitates with larger size are driven by a decreased interface free energy between the particles and matrix[26,27].The larger size and lower number density of Mg17Al12particles in the RHPRed alloy lead to the weaker pinning effect on GBs compared to the HPRed alloy.Thus,the average grain size of RHPRed alloy(～2.3 μm)is larger than that of FGs in HPRed alloy(～1.5 μm).It should be mentioned that there are alsoa few sub-micron Al4Ce particles in the matrix(Fig.6(a)),which could act as obstacles to dislocation motion.However,the number of the sub-micron Al4Ce particles is quite limited,from which the strengthening effect can be ignored.

Fig.4.EBSD results of(a)-(c)HPRed,(d)-(f)RHPRed AZ80-0.8Ce alloy;(c)and(f)are corresponding KAM maps of(b)and(e),respectively;(g)and(h)are corresponding grain size distribution histograms of(a)and(d),respectively.

Fig.5.SE images of(a)-(c)RHPRed,(d)-(f)HPRed AZ80-0.8Ce alloy.

Fig.6.(a)BF-TEM micrograph of sub-micron Al4Ce phases;(b)corresponding EDS analysis.

Typical SE images of fracture surfaces of RHPRed and HPRed alloys after tensile test are shown in Fig.7.By comparison,it can be found that the microscopic failure mechanism is different in both cases.Obvious necking can be observed in the RHPRed alloy(Fig.7(a)),which cannot be seen in the HPRed alloy(Fig.7(d)).Fig.7(b)and(c)exhibit a large number of uniformly distributed dimples(as indicated by yellow arrows),suggesting a ductile fracture.On the contrary,cleavage planes with large surfaces(red arrows in Fig.7(e))and few heterogeneously distributed dimples are presented in the HPRed alloy,suggesting that cleavage fracture is the main fracture mechanism.

3.3.Tensile properties

Typical tensile engineering stress-strain curves of asextruded,HPRed,and RHPRed alloys are shown in Fig.8(a),with the corresponding tensile properties summarized in Table 2.It can be seen that both rolled alloys show an enhanced strength compared to the as-extruded alloy.A significantly enhanced YS from 147.5 MPa for the as-extruded alloy to 308.2 MPa for RHPRed alloy is achieved,which is more than doubled.The ultimate tensile strength(UTS)of HPRed(357.5 MPa)and RHPRed(359.9 MPa)alloys is almost identical,but the YS of the latter is 37.9 MPa higher.The EF of RHPRed AZ80-0.8Ce alloy(13.8%)exhibits barely no sacrifice compared with the as-extruded alloy(16.6%).More importantly,it shows an advantage to the HPRed alloy(7%).Corresponding curves of calculated workhardening rate(Θ=?σtrue/?εtrue)are shown in Fig.8(b).During the whole tensile process,the HPRed alloy exhibits a higher work hardening rate compared to that of the RHPRed alloy.This is mainly due to the easy activation of basal slip in random orientated FGs,the smaller sizes of Mg17Al12particles and the high storage ability of dislocations in large size CGs.Comparisons of tensile properties between the present RHPRed AZ80-0.8Ce alloy and the wrought AZ80 series alloys reported in literature are illustrated in Fig.8(c).The present AZ80-0.8Ce alloy processed by RHPR shows improved mechanical properties in terms of a high strengthductility synergy.The YS of RHPRed alloy exhibits an advantage as compared to the wrought AZ80 alloys previously reported.

4.Discussion

4.1.Strengthening mechanisms for RHPRed and HPRed AZ80-0.8Ce alloys

As shown in Fig.4,LABs are uniformly distributed in the RHPRed alloy.Wang et al.[19]reported that randomly oriented FGs formed in HPRed AZ91 are associated with the continuous dynamic recrystallization(CDRX)process,which characterized a transformation from dislocations to LABs and then HABs.In the present study,FGs in the HPRed alloy are fully dynamically recrystallized due to the large thickness reduction during single pass(～85%).In contrast,four rolling passes with a relatively small thickness reduction in each RHPR process(Fig.1b)lead to an incomplete CDRX,where most LABs cannot transform into HABs.Moreover,during intermediate annealing,both rearrangement and annihilation of dislocations occur.The formation of LABs and the transformation of LABs into HABs are accompanied by a decreased dislocation density.In summary,due to the incomplete CDRX during RHPR process and the rearrangement of dislocations during intermediate annealing,large fraction of LABs(～62%)are preserved in the RHPRed alloy.The schematic representation of microstructure evolution from initial as-solutionized state to the fourth pass of RHPR process is illustrated in Fig.9.

In order to clarify the strengthening effect from LABs,the bright-filed(BF)TEM image of the RHPRed alloy(Fig.10(a))shows the distribution of LABs in the matrix.Higher magnification and corresponding diffraction patterns of three selected areas marked by A,B and C are shown in Fig.10(b),(c)and(d),respectively.Evidences of LABs could be found by comparing the selected area diffraction patterns(SADPs)between the lattices within one grain.The misorientation angle of LABs in area A,B and C are～6.4°,～5.1° and～5.4°,respectively.A high-resolution TEM(HRTEM)image of area C is shown in Fig.10(e).The inverse fast-Fourier-transform(IFFT)image of D area in Fig.10(e)is presented in Fig.10(f).Plenty of dislocations(marked by red T shape)are distributed in the matrix adjacent to the LABs,indicating that dislocations are inhibited by LABs.Similar results are also reported by Chen’s work[34],where the evidence of hindering effect of LABs on dislocations was provided by in-situ TEM experiments.Inconsistentwith the common thought that the LABs have poor resistance to dislocations,it is reasonable to consider that LABs in the RHPRed alloy contribute to a certain amount of the YS.

Fig.7.SE images of fracture surface of(a)-(c)RHPRed,(d)-(f)HPRed AZ80-0.8Ce alloy.

Fig.8.(a)Tensile engineering stress-strain curves of extruded,RHPRed and HPRed AZ80-0.8Ce alloys;(b)corresponding work-hardening rate Θ vs.the true strain of(a);(c)the YS(UTS)vs.engineering strain for wrought AZ80(Ag)alloys(extruded[28,29];MDF[30];ECAP[31];rolling[32,33]).

To quantify the room-temperature strengthening mechanisms of two rolled alloys,the contribution of grain boundary strengthening,dislocation strengthening,precipitation hardening and solid solution strengthening were calculated.Hansen[35]implied that GBs and dislocations act as independent and linearly additive contributors to the total hardening in metallic alloys.Accordingly,σyscan be expressed as:

Fig.9.Schematic representation for the microstructure evolution of RHPRed AZ80-0.8Ce alloy from initial solution treatment sample to fourth rolling pass:(a)solid solution,(b)the first pass rolling,(c)intermediate annealing,(d)the fourth pass rolling.

Fig.10.(a)BF-TEM micrograph of the RHPRd AZ80-0.8Ce alloy;(b)-(d)selected area of A,B and C in(a)with higher magnification;(e)HRTEM image of area C in(a);(f)inverse fast-Fourier-transform(IFFT)image of the area D in(e).

whereσysis the yield stress;σGB,σdislo,σpptandσssrepresent the strengthening contributions from GBs,dislocations,precipitates and solute atoms,respectively.

The contribution of GBs to hardening can be obtained as follows:

wherekyis the Hall-Petch constant and various for different grain size.The value of Taylor factorMis～2.5 for simplicity[36];αis a constant(=0.2);Gis the shear modulus(～17 GPa).TheSVis the total area of grain boundaries per unit volume,which can be estimated from the total boundary length per unit area when applying the relationSV=.Here,BAis the total boundary length per unit area derived from the 2-D orientation micrographs.is the average misorientation angle of LABs,fHABis fraction of HABs anddHABis size of HAB grains.The values of,fHABanddHABcan be calculated from IPF maps,which are illustrated in Table.3.In Mg alloys,the Hall-Petch parameters are texture and grain size dependent[37,38].Note that both rolled alloys show a basal texture,and the discrepancy of texture intensity between RHPRed alloy and the FGs of HPRed alloy can be ignored.Accordingly,thekyis selected as 209 MPa·μm-1/2for the RHPRed alloy and FGs in HPRed alloy[39],while 281 MPa·μm-1/2for CGs[40]in HPRed alloy.The total strengthening contribution of grain boundaries is estimated to be～212 MPa and～165 MPa for RHPRed alloy and HPRed alloy,respectively.

Table 3Structural parameters estimated from EBSD orientation maps of RHPRed and HPRed alloy.

The YS contribution from dislocations can be evaluated as follows[41]:

whereρis the dislocation density.It can be estimated by the following approach[42,43]:

whereuis the unit length of a circuit around a specific point of interest(500 nm,which is equal to the scanning step length);bis magnitude of Burgers vector(|b|=0.32 nm).To ensure the accuracy,Θis evaluated from local misorientation profile in the KAM map with size of 200×200 μm,which is～1.14° for the RHPRed and～0.98° for the HPRed alloys,respectively.The corresponding dislocation density is～2.47×1014m-2and～2.14×1014m-2for RHPRed and HPRed alloys,respectively.Accordingly,the calculated dislocation strengthening value is～43 MPa and～40 MPa for RHPRed and HPRed alloys,respectively.

The strengthening effect of Mg17Al12particles can be estimated by the Orowan mechanism,which can be calculated as follows[44]:

whereνis Poisson ratio with the value of 0.35 in Mg[45];dpandfrepresent the average size and volume fraction of Mg17Al12particles,respectively.In the present work,dpandfare～230 nm and～5.8% in RHPRed alloy while they are～120 nm and～6% in the FGs of HPRed alloy.Thus,σpptis calculated to be～25 MPa and～46 MPa for RHPRed and HPRed alloys,respectively.

Fig.11.SE images of RHPRed AZ80-0.8Ce alloy for deformed at a strain of(a)～5% and(b)～10%;(c)HPRed AZ80-0.8Ce alloy for deformed at a strain of～5%.

For the AZ80-0.8 Ce alloy studied in the present study,Ce has the lowest solubility limit in the Mg matrix[46]among Al,Zn and Ce atoms.Wang et al.[16]reported that the Al4Ce eutectic phase appears even only 0.2 wt% Ce addition into AZ80.Therefore,the main solute atoms in Mg matrix are Al and Zn.Jin et al.[24]reported that for the AZ91 processed by HPR with a 80% thickness reduction,the solid solution strengthening effect from Al and Zn atoms is estimated to be～40 MPa.As a result,it is reasonable to estimate that the solute strengthening from Al and Zn atoms in the two rolled samples is lower than 40 MPa due to the formation of Al4Ce phase consuming partial Al atoms.

Finally,by accounting for the contributions fromσGB,σdislo,σpptandσss,the calculated YS values are in the range of～280-320 MPa and～251-291 MPa for the RHPRed and the HPRed alloys,respectively.They are in good agreement with the experimental YS values of both rolled alloys.In summary,it is reasonable to state that the discrepancy of YS between the two rolled samples is mainly come from GB strengthening,while dislocation strengthening,precipitate hardening and solid solution strengthening provide relatively smaller contributions to the total YS for both rolled alloys.

4.2.Reasons for high EF of the RHPRed AZ80-0.8Ce alloy

The usual outcome of rolled Mg alloys is a significant enhancement in strength with great ductility loss[47,48].In the present study,the RHPRed alloy exhibits an improved tensile strength with little sacrifice in EF,as compared to the HPRed alloy.Specifically,EF of the RHPRed alloy is about twice that of the HPRed alloy.

To investigate the reasons for different performance in EF of the two rolled alloys,SE images of RHPRed and HPRed alloys at different tensile strains are presented in Fig.11.Note that micro-cracks are rarely seen in the RHPRed alloy with a tensile strain of～5%(Fig.11(a)).While some micro-cracks(red arrows)appear along GBs in the RHPRed alloy with a tensile strain of～10%(Fig.11(b)),no obvious micro-cracks around the large size Al4Ce particles can be observed.In the present work,a high density of Al4Ce(～1.6×104m-2)par-ticles with sizes in the range of 2-5 μm are present in the matrix.Generally,the large size particles are considered as the preferential position of crack initiation than grain boundaries[49].As shown in Fig.4,the RHPRed alloy shows a relatively homogeneous fine grain structure,which seems to be beneficial for the strain coordination among grains during tensile deformation and effective in avoiding stress concentrations around the large size particles.In contrast,the micro-cracks around Al4Ce particles(yellow arrows)show an obviously larger size than that in GBs in the HPRed alloy at a tensile strain of～5%(Fig.11(c)).This likely due to the heterogeneous deformation of the HPRed alloy[20],where the basal slip is easy to be activated in the randomly orientated FGs at the early stage of tensile deformation and the pyramidal slip systems may activate in the basal orientated CGs at the later stage of tensile deformation.The heterogeneous deformation of HPRed alloy is conducive to the EF enhancement,but also leads to an inhomogeneous distribution of tensile strain.As a result,for traditional HPRed AZ series alloys without large size particles,an enhanced EF is accompanied by the improved strength,i.e.,～19.6% for HPRed AZ91 alloy[21]and～17% for HPRed ATZ821 alloy[20].However,for the HPRed AZ80-0.8Ce alloy in the present study,the heterogeneous deformation is considered to accelerate micro-cracks around Al4Ce particles,resulting in the premature fracture.

Furthermore,a large fraction of LABs(～62%)is preserved in the RHPRed alloy.In addition to hindering the movement of dislocations,LABs are reported to be served as dislocation nucleation sites under stress as revealed by Chen et al.[34],which play an important role in homogenizing the deformation process[50].It has been reported that LABs have better ability to tune the balance of strength and ductility than HABs[34].As a result,it is reasonable to claim that for alloys containing large size brittle particles,the microstructure containing relatively homogeneous fine grainswith uniformly distributed LABs is beneficial in alleviating stress concentrations,thereby delaying premature fracture and enhancing the EF.

5.Conclusion

We have successfully fabricated a AZ80-0.8Ce alloy with a high strength-ductility synergy by the RHPR process,i.e.,the YS,UTS and EF are?308 MPa,?360 MPa and?13.8%,respectively.The predominant mechanism for the high YS is grain boundary strengthening.The high EF is attributed to the relatively homogenous fine grain structure with a large fraction of LABs(～62%)preserved in the matrix,which is conducive to the alleviation of stress concentrations around coarse secondary phase particles.The LABs originated from the combined effect of incomplete DRX during rolling and the dislocation rearrangement during intermediate annealing.Accordingly,this kind of homogeneous fine grain structure with a high density of LABs is effectively in delaying fracture and enhance the EF of alloys containing large particles during tensile deformation.

Acknowledgements

This work was primarily supported by The Natural Science Foundation of China under Grant Nos.51922048,51871108 and 52001133.Partial financial support came from the Fundamental Research Funds for the Central Universities,JLU,Program for JLU Science and Technology Innovative Research Team(JLUSTIRT,2017TD-09)and The Science and Technology Development Program of Jilin Province(Nos.20200201193JC and 20210201115GX).