Xi Zho,Shuhng Li,Zhimin Zhng,Pnghng Go,Shuiling Kn,Ff Yn

a School of Mechanical and Electrical Engineering,North University of China,Taiyuan 030051,China

b Engineering Technology Research Center for Integrated Precision Forming of Shanxi Province,North University of China,Taiyuan 030051,China

c China Aerodynamics Research and Development Center,Mianyang 621000,China

d The People’s Liberation Army NO.5702,Xianyang 712000,China

e Ningbo Branch of Chinese Academy of Ordnance Science,Ningbo 315103,China

Received 11 January 2020;received in revised form 24 April 2020;accepted 18 May 2020 Available online 25 June 2020

Abstract This paper provided an eff cient single pass severe plastic deformation(SPD)method,annular channel angular extrusion(ACAE),for fabricating AZ80 magnesium alloy shell part.The effect of ACAE process on the microstructure homogeneity,texture,and mechanical properties of extruded part was experimentally investigated.For comparison,conventional backward extrusion(BE)was also conducted on processing AZ80 part with same specification The results showed that ACAE process has a better capacity to refin the microstructure and dramatic improve the deformation homogeneity of the extruded part than BE process.Due to two strong shear deformations were implemented,ACAE process could also concurrently modify the basal texture more notably than BE process.In particular,a bimodal texture was found in ACAE extruded part,which was greatly related to the enhanced synergetic action of basal slip and secondary＜c+a＞slip caused by the effective shear stress.More uniform and superior hardness along the thickness and height of part were achieved via ACAE process.Further surveying of tensile tests also showed that the part fabricated by ACAE process exhibited significantl higher and far more homogeneous tensile properties with an excellent balance of strength and ductility.The remarkable enhanced tensile properties of ACAE extruded part could be primarily attributed to the significan grain refinement which provided a powerful grain boundary strengthening effect and meaningfully suppressed the development of twin-sized cracks during tensile deformation.It was established that ACAE process seemed to be a very promising single pass SPD method for manufacturing Mg-based alloy shell parts with more homogeneous microstructure and superior performance.? 2020 Published by Elsevier B.V.on behalf of Chongqing University.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;Annular channel angular extrusion;Microstructure;Texture;Mechanical properties.

1.Introduction

As the lightest commercially available metal,magnesium(Mg)alloys fin widespread interest in the fiel of transportation and aeronautical industries due to their high strength-to-weight ratio,satisfie castability and machinability[1-4].Nonetheless,the hexagonal close-packed(HCP)crystallographic structure of Mg,with limited available slip systems,brings about poor mechanical properties and deformability at room temperature[5,6],which restricts the application and promotion of Mg alloy products.It was perceived that the thermo-mechanical processing could effectively refin microstructure and eliminate casting defects,which basically improved the mechanical properties of Mg alloy materials[7-9].Particularly,severe plastic deformation(SPD)methods have been studied to improve the comprehensive mechanical properties of Mg alloys through not only refin the microstructure but also intensely modifyweaken the crystallographic texture[10-12].

As typical hollow-shaped cylindrical components,shell ube parts are fundamental form of products in numerous industrial applications.Over several decades,the backward extrusion(BE)method has played an important role in fabricating this kind of parts.It provides economical and environmental advantages such as material consumption,high surface quality and productivity[13].Even so,the unsteady deformation zone during the BE creates an uneven strain distribution in the extruded part[14].Inhomogeneous microstructures and mechanical properties in the extruded part have been observed[15,16],and it can be predicted that this heterogeneity will become significan with increasing dimensions of the products.When forming larger-scale parts,the load is no longer the only limitation which restricts the BE method,the unstable mechanical properties further limit their applications as the key load-bearing structures.The increased demand for the fabrication of this type of parts stimulated development of the combination of SPD and extrusion technology.Several combined methods such as tubular channel angular pressing(TCAP)[17],parallel tubular channel angular pressing(PTCAP)[18],tube cyclic expansion extrusion(TCEE)[19],tube cyclic extrusion-compression(TCEC)[20],and hydrostatic tube cyclic expansion extrusion(HTCEE)[21]were proposed to produce the high-quality tubes.Faraji et al.[17]investigated the effects of the TCAP on the microstructure and mechanical properties of AZ91 Mg alloy.They reported that TCAP introduced three shear zones during equal diameter tubular extrusion,and refine the grain to～1.5μm after one pass,which enhanced the tension properties of AZ91 Mg alloy.Other studies studied the deformation behavior of the AZ31 Mg alloy tubes,processed by multi-pass PTCAP[18],through tensile tests at different temperatures and strain rates.They revealed that grain refinemen and homogeneity of the microstructure were enhanced as the number of deformation passes increased,and the fines average grain size of 6.8μm was obtained from the four-pass PTCAP processed sample.Abdolvand et al.[22]applied PTCAP and tube backward extrusion(TBE)on the AZ31 Mg alloy,and microhardness in the processed tube was about two times higher than the as-received condition.However,limited number of studies has been contributed to utilize cylindrical billets to form tubeshell parts.Compared with tubular extrusion,the potential of directly forming large-scale hollow parts with small diameter billets is more promising,which can further reduce the material supply costs.In this aspect,Jamali et al.[23]carried out a study on the effect of radial-forward extrusion(RFE)on the microstructure and metallurgical properties of AZ91 Mg alloy.The results demonstrated that average grain size of the extruded tube could be greatly refine to 3μm(150μm in the initial cylindrical billets),and high YS of 235MPa,UTS(ultimate tensile strength)of 325MPa with an appropriate EL to failure of 7.7% were achieved concurrently.Recently,Shatermashhadi et al.[14]firs proposed a novel method for the BE,annular channel angular extrusion(ACAE)process,which applied a higher shear plastic strain and achieved a better strain homogeneity during the deformation.In their research,the feasibility analysis was performed using a lead blank with initial diameter of 20mm,and eventually a larger diameter cup-shaped part with the large outer diameter of 63mm was successfully extruded.Also,Hosseini et al.[24]successfully applied ACAE method to extrude a pure aluminum(Al)part.The result showed obvious grain refinement and uniform microhardness with more than 52% improvement was achieved in the wall of the extruded part.

Among a wide variety of processing methods,ACAE is of particular interest for Mg alloy hollow-shaped components(shell ube parts)because it has several advantages:(1)Due to the limited deformation zone in mandrel,triaxial compressive deformation is facilitated and deformability of Mg alloy can be immensely improved[14].(2)Smaller billet diameters also sharply reduces the forming load requirements and greatly improves the forming stability,especially compared to traditional BE method[24].(3)Furthermore,ACAE is an effective single-pass SPD method,which introduces two shear zones in the mold structure,can further promote microstructure refinemen and well coordinate the deformation texture of Mg alloys.And it is believed that it is very beneficia to improve comprehensive mechanical properties of the Mg alloy materials.

Based on such advantages of ACAE process,the aim of the current research is to investigate the effect of the this novel SPD process on the microstructure homogeneity,texture,mechanical properties of typical commercial AZ80 Mg alloy.For comparison,conventional BE method was also conducted on processing AZ80 part with same specification This research is of significanc in promoting the ACAE technology towards the processing of Mg alloy shell part with superior mechanical properties.

2.Experimental procedures

In this study,a commercial as-cast AZ80 Mg(Mg-8.0Al-0.5Zn-0.11Mn,wt%)alloy was employed.The initial material was provided in the form of an as-cast ingot with a diameter of 450mm.Two different specification of bar with Φ200×85 mm2(for BE deformation)andΦ90×360 mm2(for ACAE deformation)in size were machined from the ascast ingot,and then homogenized at 415°C for 16h so as to be extruded.As shown in Fig.1a,the as-cast AZ80 alloy exhibited a typical dendritic structure consisting predominantly ofα-Mg matrix and interdendritic eutecticβ-Mg17Al12phases distributing mostly at the grain boundaries[25].After homogenization treatment,the intergranular network eutectic phases have basically redissloved into the Mg matrix,and a uniform microstructure with an average grain size of～350μm was developed(Fig.1b).Fig.2 shows the schematic of ACAE and BE methods as well as respective mold parameters applied in this research.The specifi implementation principles and detailed operating method for ACAE process has also been reported by Shatermashhadi et al.[14].Prior to extrusion,the billets were preheated at 350°C for 3h to enhance their deformability and reduce the risk of cracking.At the same time,the molds were also preheated up to 360°C to ensure the extrusion temperature still met the requirement after the temperature radiation loss during the installation.Then,the extrusions were conducted at 4-THP-630 hydraulic press with a punch speed of～1mm/s(350°C),together with the MoS2paste as lubrication.Finally,two high quality shell parts with 200mm in outer diameter,168mm in inner diameter,and～210mm in height were prepared.

Fig.1.SEM microstructures of(a)as-cast and(b)as-homogenized AZ80 Mg alloy.

Fig.2.Schematic and specifi mold parameters of(a,c)ACAE and(b,d)BE methods applied in this research(TD,transverse direction;ND,normal direction;ED,extrusion direction).

The microstructures of the two extruded parts were observed using optical microscopy(OM,Zeiss A1m),scanning electron microscopy(SEM,Hitachi SU5000),and electron backscattering diffraction(EBSD).The average grain size and phase size were measured using Image-Pro Plus 5.0 software via optical and SEM microstructures,respectively.EBSD observation was conducted on the Hitachi SU5000 SEM equipped with an EDAX-Team EBSD system operating at 20kV with a scan step length of 1.0-1.2μm(depended on the minimum grain size).For microtexture analysis,however,a large area containing high enough grains of each sample was examined with an appropriate larger scan step length.The EBSD data were analyzed by Orientation Imaging Microscopy(OIM)Analysis software.In OIM software,particularly,the intragranular misorientation axis(IGMA)analysis was applied to identify the geometrically necessary dislocation(GND)types caused by plastic deformation[26-28].According to the literatures[26,27],the dislocation type could be basically identifie by the corresponding Taylor axis(the axis of plastic rotation due to the activity of potential slip system).In this study,the activity Taylor axes,including basal＜a+c＞and prismatic dislocations,were investigated using IGMA analysis.Besides,the grain orientation spread(GOS)define by calculating the average misorientation between all pixel point within each grain was applied to count the dynamic recrystallization(DRX)fraction of tested samples[28].In this work,the DRXed grain was identifie by the GOS value smaller than 2°,while the unDRXed grain greater than 2° Tensile tests were conducted at ambient temperature using an Instron 5967 testing machine under a cross-head speed of 1mm/min.Tension specimens,whose gauges were 18mm in length and 5mm in width,were extracted from the extruded parts,and finall ground to a thickness of 2mm.At least three samples were tested in each group of experiments.Besides,the hardness was determined using a Brinell indenter with a load of 62.5 Kg as well as a loading time of 15s.To ensure the reliability of the results,5 indentations were measured under each condition.

The DEFORM 3D-v11 software was used for FEM simulation.Due to the axisymmetric of BE and ACAE methods,the FEM analysis of deformation behavior was carried out using Deform-2D model.The material and the die components model were define as plastic and rigid bodies,respectively.The material used in FEM simulation was considered to be the same as the practical experiment,and it was meshed to 10,000 square elements with four nodes.The shear friction coefficien was assumed to be 0.3 between the billet and the mold,which was frequently used in the hot forming with lubrication.

3.Results and discussions

3.1.FEM analysis

In order to quantitatively compare the effects of BE and ACAE processes on deformation characteristics of materials,FEM analysis was conducted to simulate the deformation of billets during these two processes by using actual mold parameters.Fig.3a,b and d present the distribution of effective strain on the cross-sections(ED-TD plane)of the FEM simulated BE and ACAE processed parts.Generally,the effective deformation positions provided by the BE process are irregular upsetting deformation zone on the bottom,additional shear deformation zone near the punch corner and potential friction zone at the punch edge[15,16].The irregular upsetting process is the primary driving factor for upward f ow of metal.These actually gave rise to inhomogeneous plastic deformation as well as the dramatically drop in strain level as it moves towards the outside of the part.It can be seen that there was an evident strain gradient has formed between the inner and outer walls of BE processed part.The maximum effective strain of the selected nodes near the inner wall was～2.33 in this calculated simulation,while that approach the outer wall was rapidly reduced to～1.37,which was almost 1.7 times for the difference.Consistent with the report by Chalay-Amoly et al.[15],the effective strain along the height of the part has also been changed in a wide range(～0.67-1.58 in this simulation).It should be note that such inhomogeneous strain distribution derived from the incipient stage of the billet deformation,which was concomitant with uneven upsetting deformation.In contrast,in the ACAE process,the billet undergoes radial fl w through the initial upsetting deformation,followed by two intense shear deformations at the entrance and end of the annular channel,and finall the part is formed around the mandrel.Apparently,the supplementary shear deformation at two corners has made the effective strain in the wall regions of the extruded part markedly higher than that obtained by the traditional BE method(Fig.3d).More critically,since the variation of forming mode,more adequate material fl w was achieved via the annular channel and thus the hard-to-deform areas has been eliminated effectively.As shown in Fig.3a,the strain values of the selected intermediate nodes(p3-p7),representing the most of the wall thickness direction,were relatively well-distributed,and their fluctuation were far lower than that obtained by BE method.For the inner and outer walls of the part,however,the higher effective strain actually derived from the extra friction strain during the whole deformation history.The material withstood additional friction stress at the regions of radial channel and corner sizing belt,which ultimately led to higher cumulative strain on both sides of the wall.Furthermore,it can be seen that the variation of the effective strain along the height of the part by ACAE method was also far more homogeneous in comparison to the BE processed one,and this variation range was～0.47 in the present simulation(Fig.3b).The result is concurrent with other studies by using different mold parameters[14,24],such as reported by Shatermashhadi et al.[14],revealing this range was～0.62.It is well established that different strain patterns inevitably lead to distinct microstructure features of materials,which in turn induces different responses on their mechanical properties.

Besides,Fig.3c presents the FEM analysis calculated force-displacement curve recorded during the BE and ACAE processes.The curve of BE process was consisted of a simple stage,in which the force increased to the f ow stress of the material and then the load gradually reduced until the end of process.The curve of the ACAE process was characterized by three typical stages,corresponding to the initial upsetting deformation and the subsequent two shear deformations when passing through the mold corners,respectively,as shown in the figure It can be noted that for the two different methods to produce the shell parts with same specification the highest driving force requirement for the ACAE method was far lower than that of BE method,and their ratio was～3:1 in the current calculated simulations.This means that ACAE method shows incomparable advantages in preparing shell parts with larger-diameter and thinner-walled.

Fig.3.The comparisons of FEM calculated(a,b,d)effective strain and(c)load demand for the parts processed by BE and ACAE methods(p1-p8 and n1-n13 are equidistant nodes,respectively,where p1 and n13 are points 1mm and 25mm away from the inner wall and the bottom of the extruded parts).

3.2.Microstructure characteristics

As illustrated by the FEM analysis,the material exhibited distinct deformation characteristics under the two deformation methods owing to the discrepant deformation modes.In most areas of the wall for the two extruded parts(except their tops),strain unevenness mainly occurred along the wall thickness direction,especially for BE extruded one.To evaluate deformation uniformity and microstructures of applied alloy after these two extrusion methods,the optical microstructures of the entire wall thickness along the extrusion direction(ED)of the two extruded parts were specificall detected(both sampling positions were close to 1/2 of the parts height,labeled as“BE sample”and“ACAE sample”,respectively),as shown in Figs.4 and 5,respectively.

In Fig.4a,the macroscopic optical microstructure of BE sample is displayed.It is apparent that the macroscopic metal fl w of the BE sample substantially conformed to the strain distribution simulated by the FEM analysis.The metal f ow was evidently clustered in the inner side of the part due to the higher accumulate strain applied.As moving to the outside region,however,the metal streamline has widened and gradually reduced,conforming to the exerted poor strain.As shown in Fig.4c,this variation in streamline morphology can also be well reflecte from the f ow grid variation by the FEM analysis.In the inner region,the grid was severely deformed,representing the more effective microstructure refinement However,with the movement to the outside of the wall,the degree of grid enrichment has greatly decreased,reflectin the uneven metal fl w and the obvious reduction in strain level from inside to outside.In Fig.4b,the average grain size at different locations along the whole wall thickness and the optical microstructures of three selected regions of BE sample are presented.For better understanding,the corresponding inverse pole figure(IPF)maps obtained by EBSD are also displayed in Fig.6a-c.As is seen,the average grain size near the inner wall region of the sample has been dramatically refine to～22.9μm,in which the microstructure was composed of residual deformed grains and fin DRXed grains with relatively high DRX ratio.However,the DRX fraction continued to decrease,and the average grain size rapidly increased to～39.2μm as it approaches the outside side of the wall.It is evident that the microstructure converted into a necklace-like grain structure with dispersed DRXed fin grains interact with plenty of coarse deformed grains.Besides,it can be seen from the grain boundary structure distribution(Fig.6a-c)that there was always a certain number of low angle grain boundaries(LAGBs)distributed inside and around the coarse deformed grains,representing the incomplete DRX structure with the appearance of dynamic recover(DRV)[29].It is well known that LAGBs are proliferated by the consecutive cross-slip and climbing of dislocations and the formation of three-dimensional network structures,which can further trap the moving dislocations and transform them into high angle grain boundaries(HAGBs)(i.e.,continuous DRX(CDRX))[28,30].At this moment,however,the DRX process has not yet finishe fully,as the DRX nucleation is a dynamic and continuous process.Hence,in the BE deformation,insufficien strain in most regions scarcely induced ample DRX structure of the material.With a large quantity of DRV structures spread in various degrees,the grain distribution along the wall thickness of the part exhibited an uneven state.

Fig.4.(a)The macroscopic optical microstructure along the wall thickness of the BE extruded part(BE sample),(b)the corresponding average grain size distribution and the typical optical microstructure of three selected regions(I,II,and III)from the sample,and(c)the f ow grid distribution during BE process simulated by FEM analysis(regions I,II,and III refer to areas～1mm,～8mm,and～15mm from the inner wall of the extruded part,respectively).

Fig.5.(a)The macroscopic optical microstructure along the wall thickness of the ACAE extruded part(ACAE sample),(b)the corresponding average grain size distribution and the typical optical microstructure of three selected regions(I,II,and III)from the sample,and(c)the f ow grid distribution during ACAE process simulated by FEM analysis.

Fig.6.IPF maps of the regions I,II,and III from(a-c)BE and(d-f)ACAE samples.

In contrast,Fig.5a shows the macroscopic optical microstructure of ACAE sample,exhibited the metal streamline became much more slender,with a relatively identical distribution in the middle regions,revealing more uniform deformation of the material during the extrusion.Obviously,this is in good agreement with the result of FEM analysis,the effective strain distribution along the middle region of the wall was relatively uniform(Fig.3a).By measurement,the metal streamline in the middle regions was approximately at an angle of～25-40°to the ED.The more sloping streamline distribution well elucidated that strong shear deformation has effectively implemented during ACAE process.However,it is prominent that the streamline distribution has also changed to varying degrees on the inner and outer walls due to the presence of stronger friction stress than BE process.Again,such metal f ow characteristics can be further understood via the fl w grid distribution simulated by FEM analysis,as shown in Fig.5c.It is clear that the additional frictional stress appeared on the radial channel and sizing belts of the second corner during the entire deformation history caused the metal fl w near both sides was progressively impeded,which in turn resulted in metal fl w agglomerating inward and straightening outward respectively.Furthermore,it can be seen from the selected optical and EBSD microstructures(Figs.5b and 6d-f)that the material exhibited fine grain size and more complete DRX structure without any elongated deformed grains after ACAE treatment.The microstructures were characterized by the typical bimodal grain structure that composed of coarse and fin DRXed grains with significantl higher HAGBs ratio in most regions,while the frictional stress occurred in both sides of the wall also ineluctable altered such pattern to some extent.Compared to BE processed part,it is obvious that the maximum average grain size differentiation along the wall with a thickness of 16mm has sharply reduced to～3.4μm,with a relatively identical grain distribution in the middle region(±2.6μm).This shows that ACAE process by achieving more sufficien metal f ow and introducing higher effective strain could effectively promote grain refinemen and markedly reduce the deformation inhomogeneity of the extruded part than BE process.It is well known that fine grain and more identical grain size distribution tend to represent higher and more uniform mechanical properties of the materials.Hence,this result potential clarifie that ACAE treatment is a feasible single pass SPD method for preparing homogeneous shell members with superior mechanical properties.Nevertheless,it should also be noted that due to the stronger friction occurred during deformation than the BE method,the microstructure was also inevitable affected to varying degrees on both sides of the wall.So how to reduce the impact of friction in ACAE process is still an important issue to be considered.

Fig.7.Representative SEM microstructures were selected from(a)BE and(b,c)ACAE samples;(d)the scanning result of elements distribution from the selected area in(c)showing the dynamic precipitation ofβ-Mg17Al12 phase particles.

In Fig.7a-c,the representative SEM images selected from two samples are appended,to reveal the phase characteristics after extrusions.As shown,there were plenty of fin phase particles(identifie asβ-Mg17Al12phases,Fig.7d)which did not existed in the homogenization state have notably occurred after extrusions,and they were mainly distributed in the fin DRXed grain regions,especially the grain boundary(GB)regions.Similar phenomenon has also been found in various wrought Mg alloys,which was referred to dynamic precipitation[10,31,32].By measuring,the average size of these precipitated particles in the two samples was comparable,and their values were～780nm(BE sample)and～830nm(ACAE sample),respectively.These were larger than the reports by Kang et al.[25]and Zhou et al.[10]during slow extrusion of AZ91-0.5Ca alloy(～168-241nm)and multidirectional forging(MDF)process of AZ80 alloy(～370-470nm),probably owing to higher deformation temperature catalyzed their coarsening.Generally,plastic deformation easily induces high density dislocations and vacancies in the matrix,which provides favorable location for rapid solute diffusion and subsequent heterogeneous phase precipitation[33].In this aspect,the DRXed regions or regions are undergoing DRX absorbed profuse deformation energy and created huge quantity of LAGBsHAGBs during deformation,it is therefore rationalized that these regions becoming the favorable site for phase nucleation.Furthermore,the precipitated fin phase particles distributed around GBs usually also can provide a strong barrier for impeding the GB migration and grain growth via Zener pinning effect[10,25,34],thus promoting the grain refinement Recently,Pan et al.[35]developed an ultra-high strength Mg-2Sn-2Ca alloy via conventional extrusion,and they found that dynamic precipitation actually occurred more frequently in high density dislocation stacking zones during extrusion,which generally earlier than the fully DRX process.And subsequently,the generation of the nanoscale Mg2Ca phases could enormously restrict growth of the subgrainsDRXed grains by pinning effect,contributing to excellent tensile strength.However,on the other hand,although these precipitatedβphase particles could effectively hinder the growth of the DRXed grains,it should also be considered that their appearance would also consecutively consume the solute atoms(Al atoms)in the surrounding areas by GB diffusion,which led to the formation of solute-depleted areas.Especially with the coarsening of these phase particles,the surrounding solute content would be further consumed.Hence,the subsequent nucleation of DRXed grains in the regions a little bit away from these phase particles had a less content of solute,and thus led to a weak precipitation driving force at the GBs[36].Resultantly,the precipitate rare regions with relative coarse DRXed grains possibly appeared and eventually aggravated the uneven grain distribution,as shown in Fig.7b.

3.3.Texture characteristics

Fig.8.(0001)basal pole maps of the regions I,II,and III from(a-c)BE and(d-f)ACAE samples.

Fig.8 present the overall micro-texture features of the two samples.It is evident that due to the different forming modes applied,the basal poles of the two extruded parts along the wall thickness exhibited discrepant components.In most regions of the BE sample,due to the dominated distribution of uniaxial extrusion stress(could also be affected by friction stress in the inner region),the c-axis of most grains was preferentially aligned to the thickness direction(TD),with certain similarities to the rolling texture in Mg alloys.However,the c-axis of all these grains was not always remained at TD,and some of them have also deviated from TD in a certain range of angles,which eventually dispersed the basal poles.And it can be seen that this was more notable in the inner region with higher DRX fraction and more fin DRXed grains.This is consistent with the fact that the fin DRXed grains usually have a relatively random orientation,which weakens the deformation texture[37].However,considering the aspect of deformation mechanism,the dispersion of basal poles away from TD could be caused by the enhanced activation of pyramidal＜c+a＞slip at high deformation temperature.As evident in Fig.8e,this trend has got more obvious in the ACAE extruded part,in which two separated texture components already developed remarkably(middle region)under the shear deformation.It is obvious that the texture component with its maximum basal pole tilted by～30-45 ° toward the ED was the most common shear texture,originating from dominated activation of basal slip system[38].However,the appearance of another notable texture component with its maximum basal pole aligned～0-20 ° toward the ED can be considered caused by the enhanced activation of＜c+a＞slip[39,40].It has been revealed when deformed at elevated temperatures,the non-basal slips will be greatly activated owing to more energetically beneficia as compared to the activation of basal slip is barely temperature dependence.The activation of non-basal slips,especially the pyramidal＜c+a＞slip,can enormously improve the deformation capacity of Mg alloys.Generally,screw dislocations of＜c+a＞slip type can move to the next slip planes by cross-slip and climb,which plays a vital role in DRV behavior of Mg alloys[41].At this studied ACAE process with high extrusion temperature(350°C),it can be noted that the texture intensity caused by pyramidal＜c+a＞slip was already comparable to the shear texture,showing that the dislocation glide in pyramidal slip system already established significantl.The activation of＜c+a＞slip consumed the shear texture and thus led to the formation of bimodal texture,which reduced the overall texture intensity to a certain extent.Interestingly,the current results also present that owing to the effect of strong friction stress,both the texture component and intensity also varied to different degrees on ACAE sample.There were two new texture components appeared in the inner side with their maximum basal poles aligned almost perpendicular ilted～30 ° to the ED as an expense of previous components(Fig.8d).However,the two previous textures were agglomerated in the outer side,forming a stronger single-peak texture component with its maximum basal poles tilted～0-30° to the ED(Fig.8f).That is,the c-axis of grains was gradually rotated to the two sides of the wall under the function of uniaxial friction stress,respectively.Similar result was also reported by Sagapuram et al.[42]in studying the effect of large strain extrusion machining on the texture evolution of AZ31 Mg alloy.

3.4.Microstructure evolution during ACAE process

Excellent grain refinin and texture modificatio abilities have clarifie that the ACAE technology can be considered as an effective SPD method.To better understand the grain refinemen behavior and texture evolution of applied material during this novel single pass process,microstructure evolution by EBSD was examined from three typical positions(i.e.,upsetting zone(UZ),shear zone I(SZ I)and shear zone II(SZ II))along the annular channel.The corresponding IPF figures pole figure and sampling schematic are presented in Fig.9,respectively.

Fig.9.IPF maps of the UZ,SZ I and SZ II and the corresponding(0001)and(10-10)pole f gure maps revealing the microstructure evolution during ACAE process(FD:metal fl w direction or shear direction).

In the initially billet deformed region,deficien strain rarely induced sufficien DRX nucleation of the alloy,and a necklace-like grain structure with fin DRXed grains in a size range of～6-12μm was developed.Usually,owing to the large grain size in the starting material,the imposed strain could not be well delivered within the grains,and more stress incompatibility could be preferentially accommodated in the periphery of initial GBs[43].With the influenc of their plastic compatibility stress,the basal on-basal cross-slips tend to be activated and local shearing could easily occur.This would yield intense lattice rotation and subgrain nucleation near initial GB,resulting in progressively development of necklacelike grain structure[44].However,the striking feature in this stage is that there were many lenticular structures,identifie as typical{10-12}extension twin,with different morphologies and sizes have nucleated within the coarse grains,ending at the GBs or intersecting each other[45].It is well established that{10-12}extension twins usually possess lower critical resolved shear stress(CRSS)and susceptible to activate at the initial stage of deformation with coarse grain size and favorable orientation.Different from the{10-11}contraction and{10-11}-{10-12}double twins in Mg,the extension twin boundaries(TBs)are extremely easy to expand and migrate under successive deformation for their intrinsic structure[46,47].In this involved extrusion,twin refinemen behavior can be divided into two types.Firstly,it can be seen that the similarly oriented twin variants T1,T2 and T4 have evidently expanded and extended,traversing and gradually engulfin matrix grains,which eventually refine the coarse grains.And it should be believed that these oriented variants have a strong propensity to activate and expand preferentially during this deformation.However,when variant T3 and T5 that similar to the above types encountered distinct variant T6,TBs migration was enormously hindered,but these interacting twins could also effectively subdivide the matrix grains and refine them into different portions.Obviously,this type of twin refinemen contradicts the firs one,and it has also been reported by Kadiri et al.[48].They presented that the cross of different twin variants tended to improve the twin nucleation rate,but rapidly reduced the growth ability for single twin.Thus,both twining engulfin and segmenting of coarse grains greatly contributed to grain refinemen in initial stage with lower strain.

Fig.10.The distribution of average grain size,LAGBs fraction,and average misorientation angle of UZ,SZ I and SZ II obtained by EBSD.

With shifting to SZ I,twins basically disappeared since the great reduction of coarse grains.Instead,under high amplitude of shear strain,a mixed grain structure appeared successively.This included profuse fin(～4-12μm)DRXed grains,relative coarse DRXed grains,and some elongated deformed grains,and all of them were distributed along the f ow direction(FD).Obviously,compared with initial stage,the average grain size have been reduced,while the proportion of LAGBs has greatly increased since the strong shear deformation induced massive dislocation proliferation(Fig.10).As mentioned earlier,such dislocation proliferation and the development of LAGBs were involved with early stage of CDRX.In more detail it can be noted that the high ratio of LAGBs in this stage were mainly clustered inside the elongated grains.And the prominent feature is that they were distributed tended to parallel to each other,which divided the grains and promoted the development of new grain with different basal orientation(Area 1,G2).To clarify this interesting process,the representative grain(G3)in Area 2 with much more severe strain accumulation was extracted,as shown in Fig.11.Obviously,the G3 consisted of original purple and new developed light green colored regions with distinct different basal plane orientations.The local misorientation map and the point-to-pointorigin profil line(L1)in Fig.11a and d present high residual plastic strain concentration have remarkably formed within the grain.This could provide suffcient strain gradient for stimulating DRX nucleation.Moreover,the possible dislocation types caused by these near parallel LAGBs were further conducted through the IGMAs analysis,revealing dominant arrangement of basal＜c+a＞GNDs with the distinctive characteristic of the[10-10]rotation(Taylor)axis mostly occurred,while prismatic dislocations with the rotation of[0001]axis were scarcely distributed.This well implies the meaningful activation of basal＜c+a＞slip within the deformed grains was crucial factor for promoting nucleation of the different basal oriented grains under this shear deformation stage.Furthermore,it can be more clearly seen from the texture evolution that the newly nucleated grain(G2)or the matrix with light green and their corresponding elongated deformed grain(G1 and G3)obviously belonged to two different texture components that have formed in this stage(labeled as“1”and“2”).It is evident that the basal pole of texture 1 was almost parallel to FD(or shear direction),corresponding to the shear texture dominated by basal slip.Hence,combining with the result of IGMA analysis,the texture 2 actually caused by enhanced activation of＜c+a＞slip.As a result,it is further clarifie that the nucleation of DRXed grains within the deformed grains originated from the accumulation and release of dominated＜c+a＞dislocations.Also,it is suggested that the formation of bimodal texture in the ACAE extruded part was closely related to the enhanced activation of＜c+a＞slip on the basis of basal slip.

Fig.11.The DRX behavior of the unDRXed grain(G3)selected from Area 2 in Fig.9:(a)the local misorientation map,(b,c)the distribution of basal＜c+a＞and prismatic GNDs obtained by IGMA analysis,(d)line profil of point-to-pointorigin misorientation angle along the L1,and(e,f)the corresponding discrete(0001)and(10-10)pole f gures.

When shifting to SZ II,the residual deformed grains have been gradually consumed under the action of further shear strain,and more sufficien DRXed grains were developed with the continuous transition from LAGBs to HAGBs(CDRX)(Fig.10).Comparatively,the fin grains nucleated in earlier stage grew to different degrees and some fine DRXed grains have also nucleated on their boundaries owing to a new round of DRX.Besides,it should be noted a small amount of elongated deformed grains were still retained and distributed along the FD.Just like the G4,parallel LAGBs already appeared within the grain,representing the chief stacking of the＜c+a＞dislocations has occurred notably.And it is obvious that these grains would be wholly consumed after the material passed through this stage entirely(Fig.6e).In addition,it can be seen from the texture evolution of the three stages that the texture intensity gradually decreased with the increase of the DRX ratio and the grain refinement And at this stage,it seems that the rudiments of two textures that appeared in the wall region(Fig.8e)have formed,representing the activation and transmitted of basal and＜c+a＞slips with the development of plastic deformation.As a result,the significan grain refinemen of applied AZ80 alloy in the ACAE process was closely associated with the twin engulfin or segmenting refinemen in the coarse grain stage.And subsequently,two powerful shear deformations refinemen that enormously catalyzed both the fin and coarse DRXed grains nucleation,involving the chief mechanism of CDRX.And in this process,the texture intensity was markedly reduced.In particular,the rearrangement of dominated＜c+a＞GNDs divided and refine the elongated deformed grains into different basal oriented DRXed grains has been greatly found,which actually contributed to the formation of bimodal texture.

3.5.Mechanical properties

3.5.1.Hardness

Fig.12.The Brinell hardness(HB)distribution along the(a)wall thickness and(b)height directions on the cross-section of the BE and ACAE extruded parts(where these test positions were purposely chosen to basically follow the node positions in Fig.3).

Fig.13.The engineering stress-strain curves of tested samples were extracted along the wall thickness of(a)BE and(b-c)ACAE extruded parts(sampling position was close to 12 height of the extruded parts).

Fig.12 presents the Brinell hardness(HB)distribution along the wall thickness and height directions on the crosssection of the BE and ACAE processed parts.Concurrent with the microstructure observation results in Figs.4 and 5,the hardness of the two extruded parts has changed to varying degrees along the wall thickness,with relatively higher values in the inner side.And it is clear that this trend was more pronounced in BE extruded parts owing to more signifi cant grain size difference.In contrast,the hardness difference of ACAE extruded part along the height became less obvious,with only an abrupt decline appeared at the top regions(～66.8-63.1 HB),and this good testimony to the FEM calculated effective strain distribution(Fig.3).However,owing to the uneven upsetting deformation at the initial stage during BE process,such a trend became much more apparent in BE extruded part,and there was a larger and wider range of hardness reduction already appeared at the top regions of the part(～59.8-53.5 HB),representing poor performance in this region.Similar to the ACAE extruded one,for the relatively lower areas of the BE extruded part,the hardness along the height direction was relatively uniformly distributed.As a whole,in most areas of the wall of the two extruded parts,the uneven hardness primarily appeared along the wall thickness direction,and the ACAE extruded part displayed higher and relative more uniform hardness than BE extruded one.

3.5.2.Tensile properties

The engineering stress-strain curves of the tested samples which were extracted from the wall thickness direction of two extruded parts when tensile(ambient temperature)along the ED are shown in Fig.13.Tensile properties of all these tested samples in different regions are summarized in Table 1.It is no wonder that the part processed by ACAE process exhibited significantl higher and more uniform tensile properties.The highest tensile properties of BE extruded part,with YS of～173MPa,UTS of～305MPa and elongation to failure of～15.0%,appeared on the inside of the wall region,corresponding to the fines grain size.But as moving toward to outward of the part,all of these properties gradually decreased.It should be noted that the maximum YS,UTS,and EL differences along the side wall of extruded part reached to～19MPa,～35MPa,and～3.9%,respectively,caused by the uneven plastic rheology.In contrast,these differences were remarkably shrunk in ACAE extruded part,in which YS fluctuate between～169MPa and～174MPa,UTS between～306MPa and～315MPa,with EL differences meaningfully downed to～1.5%.The more uniform tensile properties of ACAE processed part than the BE processed one could be primarily attributed to much more identical grain size has been obtained.

Table 2 shows the EBSD calculated(0001)＜11-20＞basal slip Schmid factor(SFbasal)distributions at the regions I,II and III of BE and ACAE samples when the tensile stress applied along the ED.It is evident that due to more non-basal texture components has developed in the ACAE sample,the SFbasalvalues were relatively higher than that of the BE sample,with highest value of～0.30 occurred in the inner region.However,although a certain degree of basal texture components appeared in most regions of BE extrudedpart(especially the outer wall region),the relatively dispersed grain orientations caused by the refine grains and enhanced activation of＜c+a＞slip have also weakened them.And the calculated SFbasalvalues were also higher than that of the typical extruded olled Mg alloys that have been frequently reported(～0.17-0.20)[25,49,50].By comparing the different wall regions of the two extruded parts,it can be seen that in the case of relatively weak basal texture,the YS of ACAE extruded part was basically higher than that of BE processed one,regardless of the effect of texture strengthening.Therefore,the improved YS of the extruded parts can be considered primarily from the GB strengthening,as generally explained by Hall-Petch relationship[35,51].Besides,it has been reported that the migration of dislocation could also be effectively pinned and hindered by relative finβphase particles[25],which actually also contributing to the material strength as a supplement.Generally,the interface ofβ-Mg17Al12phase with Mg matrix is not coherent,and the phase strengthening in Mg-Al alloys conform to Orowan mechanism(incoherent strengthening)[33].However,as a result of texture weakening,especially for ACAE extruded part,the improved YS was actually lower than that of the textured Mg alloy materials.For example,the YS of AZ31 sheet processed by porthole die extrusion,with typical basal texture and average grain size of～17.2μm,fell into the range of～200-235MPa(depended on tension direction)[50],which was higher than the～169-174MPa of this ACAE extruded AZ80 part with average grain size of～16.5-19.9μm.But then again,as an expense of YS,texture weakening plays a vital role in improving the room temperature plasticity and concurrently restrain the tension-compression asymmetry of the Mg alloy materials,because much more grains that are conducive to basal slip can fully participate in plastic deformation at an earlier stage[30].Hence,for ACAE extruded part containing more non-basal texture components,the weaker anisotropy and superior ductility could be potentially achieved.

Table 1The YS,UTS,and EL of the tested samples were extracted from the wall thickness direction of BE and ACAE extruded parts.

Table 2The EBSD calculated average SFbasal.value of regions I,II and III from BE and ACAE samples.

Fig.14.(a,b)SEM fractographs and(c-f)SEM microstructures taken adjacent to the fractured surface of tensile tested samples from:(a,c,d)BE extruded part and(b,e,f)ACAE extruded part.

Fig.15.EBSD results of fractured tensile samples from(a-b)ACAE and(cd)BE extruded parts:(a,c)IPF maps,(b)selected twinned grains and their image quality(IQ)maps as indicated by white ellipse boxes in(a),and(d)(0001)and inverse pole f gures presenting the highlighted crystallographic orientation of twins and matrix from(c).

Fig.14a and b show the representative SEM fractographies of tensile tested samples from BE and ACAE extruded parts.Both the fracture surfaces of the two tested samples were characterized by the typical brittle-ductile failure mode[36,52],in which various dimples(marked with red arrows),tear ridges together with cleavage planes(marked with yellow oval dotted frames)in different sizes were observed.It is obvious that a higher density and larger size cleavage planes occurred in the fracture surface of BE extruded sample,corresponding to the existence of transgranular fracture caused by coarse DRXed esidual deformed grains.Fig.14c-f show the SEM micrographs of the longitudinal section adjacent to fractured surface of the two samples.As shown,lamellar twins occurred within the coarseunDRXed grains of both samples,and microcracks were initiated at these lamellar TBs,GBs(both fin and coarse DRXed grains),and phase particlesmatrix interfaces.It is clear that the microcracks appeared at the TBs have a large length and this tendency got more apparent as the twin size(grain size)increased.Therefore,it is inferred that these large TB cracks could be the main source of cleavage planes appeared in the fractures,which extremely deteriorated the facture properties(UTS and EL)of the material.As evident in Fig.15,these twins were further identifie as typical{10-11}contraction or{10-11}-{10-12}double twins by EBSD,in which their basal planes were reorientated～56 ° and～38 ° with respect to the adjacent matrix grains,respectively[53,54].Generally,the contraction and double twins possess high CRSS values and tend to be easily activated when the grain is in hard orientation for the basal slip and extension twin[36].The nucleation of these twins easily induces rapid f ow localization and early shear failure owing to joint function of strain softening and localized twinsized microcracks development[34,55].It has been revealed that the CRSS requirement for activating twins generally also adhere to the Hall-Petch relationship,i.e.,the twins are more easily to be activated in larger size grains[54,56].Obviously,for ACAE extruded part with relative fin grain size,the nucleation probability and size of contraction and double twins during tensile deformation have been meaningfully reduced,thereby delaying the cracks development,which eventually contributed to its superior fracture properties than BE processed one.

4.Conclusions

In this paper,the comparisons of the microstructure,texture,mechanical properties,and fracture behavior of the AZ80 Mg alloy shell parts fabricated by the BE and ACAE processes were specificall performed.The main conclusions can be drawn:

(1)By introducing a high magnitude of effective strain and achieving more adequate metal fl w,ACAE process could refin the microstructure and reduced the deformation inhomogeneity of the extruded part more effectively than BE process.The average grain size of the ACAE extruded part along the wall with a thickness of 16mm was refine to～16.5-19.9μm,while that processed by BE method was varied in a wide range of～22.9-39.2μm.

(2)During extrusion process,dynamically precipitated ofβ-Mg17Al12phase particles were initiated at boundaries of DRXed grains.The precipitated phase particles could act as powerful obstacles against the DRXed grains growth via pinning effect.

(3)Due to two strong shear deformations were implemented,ACAE process could weaken the basal texture more remarkable than BE process.In particular,a bimodal texture with its maximum basal poles tilted～30-45°and～0-20°toward ED formed in ACAE extruded part.Further microstructure evolution analysis revealed that this was strongly related to the enhanced synergetic effect of basal slip and secondary＜c+a＞slip caused by applied shear stress.

(4)More relative uniform and high hardness was achieved in the ACAE extruded part than the BE extruded one.Also,the part fabricated by ACAE process also exhibited higher and far more uniform tensile properties with an excellent balance of strength and ductility.Grain refinemen and phase strengthening(βphase particles)significantl increased YS of the parts.Texture weakening meaningfully increased ductility and reduce deformation anisotropy,but as an expense of YS.

(5)Further tensile fracture analysis revealed that the microcracks were mainly initiated at TBs,GBs and phase particles/matrix interfaces during tensile test.The grain refinemen meaningfully reduced the number of large-size twin cracks,which contributed to the distinctly higher fracture properties(UTS and EL)of ACAE extruded part than BE extruded one.

Acknowledgments

The authors gratefully acknowledge the financia supports from the National Natural Science Foundation of China(Grant no.51605448),Natural Science Foundation of Shanxi(Grant no.201701D221093),and“HIGH-GRADE CNC machine tools and basic manufacturing equipment”Major National Science and technology projects(Grant no.2019ZX04022001-004).