Novel Mg-Bi-Mn wrought alloys: The effects of extrusion temperature and Mn addition on their microstructures and mechanical properties

Qinghng Wng ,Howei Zhi ,Linto Liu ,Hongo Xi ,Bin Jing,* ,Jun Zho ,Dolun Chen,Fusheng Pn

aSchool of Mechanical Engineering,Yangzhou University,Yangzhou 225127,China

b National Engineering Research Center for Magnesium Alloys,Chongqing University,Chongqing 400044,China

c School of Mechanical and Electrical Engineering,Hunan City University,Yiyang 413002,China

d Department of Mechanical and Industrial Engineering,Ryerson University,Toronto,ON M5B 2K3,Canada

Abstract Designing and developing the Mg alloys with low cost and high performance is of the great significance.Novel Mg-1Bi-xMn (x=0,1 and 2 wt.%) extruded alloys,in this work,were fabricated at different extrusion temperatures (220,250 and 300 °C).The effects of extrusion temperature and Mn addition on the microstructures and mechanical properties of extruded alloys at room temperature were investigated.The results showed that decreasing the extrusion temperature could refine the average grain size,weaken the basal fiber texture intensity and improve the microstructural homogeneity of extruded alloys.When the Mn element was added to the Mg-1Bi alloy,the average grain size further reduced.Simultaneously,the number fraction of low angle grain boundaries (LAGBs) increased,along with the occurrence of regions without dynamic recrystallization (unDRX).The combined effects of grain refinement and coarse unDRXed structure made the textures of the extruded Mg-1Bi-xMn alloys never obviously change.Besides few large size un-dissolved second phases,fine Mg3Bi2 and α-Mn phases were precipitated in the extruded Mg-1Bi-xMn alloys and partial nano-scale α-Mn particles pined at grain boundaries (GBs) to effectively impede the migration of GBs for grain refinement.Microstructural variations determined the extruded Mg-1Bi-2Mn alloy to exhibit the highest yield strength of～319.2 MPa with the appropriate elongation-to-failure of～13% at the extrusion temperature of 220 °C,and they enabled the extruded Mg-1Bi-1Mn alloy to show the highest elongation-to-failure of～26% without the obvious loss of yield strength of～252.1 MPa.

Keywords: Mg-Bi-Mn alloys;Extrusion temperature;Mn addition;Microstructure;Mechanical property.

1.Introduction

Room-temperature (RT) strength-ductility trade-off dilemma in magnesium (Mg) alloys is a troublesome problem that needs to be solved urgently [1].In the past decade,massive studies have attempted to use rare earth(RE) alloying method [2–4] and severe plastic deformation techniques (like equal channel angular pressing,ECAP)[5–7] to overcome this issue.However,the high cost of production and the complexity of technological procedure to a large extent restrict the industrial applications of Mg alloys.Therefore,designing and developing the low-cost and high-performance Mg alloys is of the essence making use of a simple and high-effective processing approach.

Recently,Mg-Bi based alloys are becoming attractive as new wrought Mg alloy series because of their high costcompetitiveness,processing efficiency and RT mechanical properties [8–11].For instance,Mg-5Bi-3Al (wt.%) alloy,which was comprised of inexpensive alloying elements Bi and Al,could be successfully extruded at a high speed of 67 m/min without any hot cracking and showed a relatively high yield strength of 188 MPa [9].Meng et al.[8] also reported that an extruded Mg-8Bi-1Al-1 Zn (wt.%) alloy had a better balance between the strength (～291 MPa) and the ductility (～14.6%) at RT.The enhancement of mechanical properties in Mg-Bi based alloys was closely related to the Mg3Bi2phases with a high melting temperature of 823 °C that exhibited a relatively good thermal stability and played an important role in the suppression of dynamic recrystallized(DRXed) grain growth during hot forming process.However,the high Bi content containing Mg-Bi based alloys tended to face a drawback that these hard and large micro-scale Mg3Bi2particles were easily broken during deformation and became the sources of cracking in the Mg matrix,worsening the RT ductility [8,11].Reducing the Bi content is the most effective way to eliminate the disadvantage of brittle Mg3Bi2particles,while it is inevitably at the expensive of decreasing the RT strength.To address this contradiction,Mn element is considered to be added to the low-content Mg-Bi alloys for enhancing their RT strength and ductility simultaneously.Previous studies have reported the role of Mn addition on the mechanical properties of extruded Mg-Zn [12],Mg-Sn[13],Mg-Ca [14],Mg-Gd [15,16],Mg-Al [17,18] alloys.A large number of nano-scaleα-Mn particles existed in the Mg matrix by low-temperature extrusion (below 300 °C),remarkably inhibiting the growth of DRXed grains and the motion of dislocations for obvious grain refinement and precipitation strengthening.The Mn addition is expected to make up for the deficiency in the RT strength of the low-content Mg-Bi alloys with the appropriate ductility under the low-temperature extrusion condition.

Based on the aforementioned alloy design strategy,the objective of this work is to produce the novel low-content Mg-Bi-Mn wrought alloys and investigate the effects of extrusion temperature and Mn addition on the microstructure and tensile mechanical properties of Mg-Bi-Mn wrought alloys.We successfully fabricate the extruded low-content Mg-1Bi-xMn(x=0,1 and 2 wt.%) alloys under different extrusion temperature conditions (220,250 and 300 °C,respectively).The results show that the extruded Mg-1Bi-xMn alloys have a simultaneous improvement with respect to the strength and the ductility at RT with decreasing extrusion temperature,compared with the extruded Mg-1Bi alloys.Under the extrusion temperature of 220 °C,when 1 wt.% Mn content is added to the Mg-1Bi alloy,tensile yield strength (TYS) reaches～252.1 MPa and elongation-to-failure (εf) can be up to～26%in the extruded Mg-1Bi-1Mn alloy.However,by the continuous addition of Mn content from 1 to 2 wt.%,the TYS continuously increases to～319.2 MPa,but theεfdecreases to～13% in the extruded Mg-1Bi-2Mn alloy.

2.Experimental procedures

The as-cast Mg-1Bi-xMn (x=0,1 and 2 wt.%) alloys,named as B1,BM11 and BM12 alloys,were prepared in an electric resistance furnace under the atmosphere of SF6and CO2mixed gas (mixing volume ratio was 1:99).The mate-rials used in this study were commercial pure Mg ingot (≥99.99 wt.%),Mg-10Bi (wt.%) and Mg-10Mn (wt.%) master alloys.The melt was held at 720 °C for 20 min after alloying elements were dissolved into the Mg matrix,and then it was poured into a steel mold of Ф80 × 200 mm in size(diameter × height),which was preheated to 350 °C.The chemical compositions of the as-cast B1,BM11 and BM12 alloys were measured by inductively coupled plasma-atomic emission spectroscopy (ICP-AES) and the result is shown in Table 1.

Table 1 Detailed chemical compositions of the as-cast B1,BM11 and BM12 alloys measured by ICP-AES.Unit: wt.%.

Table 2 Tensile mechanical properties at RT including tensile yield strength (TYS),ultimate tensile strength (UTS) and elongation-to-failure (εf) for each extruded B1,BM11 and BM12 series alloy.

Table 3 Summaries of average grain size,number fraction of “coarse”-grained structure and number fraction of low angle grain boundary in each extruded B1,BM11 and BM12 series alloy.

All the as-cast B1,BM11 and BM12 alloys with a size of 80 × 80 mm in diameter and height machined by wireelectrode cutting were subjected to solid-solution treatment at 500 °C for 12 h,and then they were hot-extruded into rods at 220,250 and 300 °C,respectively,with an extrusion ratio of 28:1 and a ram speed of 5 mm/s after preheating at the corresponding extrusion temperature for 2 h.The extruded B1,BM11 and BM12 alloy rods corresponding to the extrusion temperatures of 220,250 and 300 °C were labeled as the B1/BM11/BM12–220,the B1/BM11/BM12–250 and the B1/BM11/BM12–300 alloys,respectively.Samples used to microstructural observations and RT tensile tests were machined into 5 × 5 × 2 mm in size (length × width × thickness) and 25 × 5 mm in size (initial gage length × gage diameter) from alloy rods,respectively.

The microstructural observation and micro-texture analysis of all the samples were carried out by electron backscattered diffraction (EBSD) technique using JEOL JSM-7800F device.EBSD preparation consisted of grinding on SiC papers of grit sizes of 280,400,600,800,1000 and 1200#,washing,blowdrying as well as electro-polishing at a voltage of 20 V and an electric current of 0.03 A for 90 s at a temperature of-25 °C with a special electrolyte named as AC2.The step sizes of EBSD scan for initial and deformed samples were set as 0.1 and 0.08μm,respectively.All EBSD data were analyzed using Channel 5 software.The average grain sizes of all the samples were measured by Image-pro plus software.The morphologies and chemical compositions of second phases of samples in solid-solution treated and extruded states could be measured by scanning electron microscopy(SEM,TESCAN VEGA 3 LMH SEM) equipped with energy dispersive spectrometer (EDS) and transmission electron microscopy (TEM) using FEI TECNAI G2 F20 device with energy dispersive X-ray (EDX) spectroscopy (operation voltage of 200 keV).In addition,X-ray diffraction (XRD,Rigaku D/Max 2500) was also used to further identify phase constitutions of extruded samples in a larger area.The foils for TEM observations were prepared by mechanical polishing to 50μm and then punched into disks of 3 mm in diameter.Subsequently,the treated foils were further thinned by ion beam using GATAN,PIPS 691 device.The thin regions in the TEM samples are about 30 nm in thickness.In-grain misorientation axes (IGMA) analysis from EBSD data was used to determine the slip system activities of deformed grains from the extruded and tensile samples.Based on the previous IGMA analysis for Mg alloys [19],the minimum and maximum misorientation angles for IGMA analysis were taken as 0.5 and 2.0°,respectively.In order to reflect the interaction between precipitates with dislocations during tensile deformation,geometric phase analysis(GPA)method was applied to reveal the strain fields of close-packed (0001) and (10–10) planes from high-resolution TEM (HRTEM) images around these precipitates.The RT tensile properties of extruded alloys were measured using the CMT6305–300 KN universal tensile testing machine at a strain rate of 1.2 × 10-3s-1.Each tensile test was performed for three times to guarantee the accuracy of experiments.

Fig.1.(a) True stress-strain curves at RT of the extruded B1,BM11 and BM12 series alloys at different extrusion temperatures of 220,250 and 300 °C;(b) comparison of the tensile yield strength (TYS) and the elongation-to-failure (εf) of the BM series alloys with these of the reported RE-free wrought Mg alloys (including Mg-Al,Mg-Zn,Mg-Sn and Mg-Bi based alloys [1,8,9,11,20-22]);(c,d) correlations between the TYS,the εf and the alloying content for the BM series alloys and the above mentioned RE-free wrought Mg alloys,respectively.

3.Results

3.1.Tensile mechanical properties

Fig.2.(a,b) SEM image of the B1 alloy in solid-solution state and the corresponding SEM-EDS results of points A and B,respectively;(c,d) SEM images of the BM11 and BM12 alloys in solid-solution state,respectively;(e,f) TEM images of the BM11 and BM12 alloys in solid-solution state,respectively.

Fig.3.EBSD inverse pole figure (IPF) maps and (0002) pole figure (PF) maps of the extruded B1,BM11 and BM12 series alloys from cross section: (a-c)the B1–220,B1–250 and B1–300 alloys,respectively;(d-f) the BM11–220,BM11–250 and BM11–300 alloys,respectively;(g-i) the BM12–220,BM12–250 and BM12–300 alloys,respectively.

Fig.1a shows the true stress-strain curves at RT of the extruded B1,BM11 and BM12 series alloys subjected to extrusion at different temperatures of 220,250 and 300 °C,and their corresponding tensile data are summarized in Table 2.With decreasing extrusion temperature,the tensile mechanical properties including TYS,ultimate tensile strength (UTS)andεfof extruded alloys gradually increase.As the Mn content add from 0 to 2 wt.%,the TYS and UTS still show a monotonically increasing trend at the same extrusion temperature,while theεfreaches a peak value in 1 wt.% Mn addition.Fig.1b compares the TYS andεfvalues of the extruded BM11 and BM12 series alloys with these of the reported RE-free wrought Mg alloys (including Mg-Al,Mg-Zn,Mg-Sn and Mg-Bi based alloys [1,8,9,11,20-22]) at RT.It can be seen that there is a remarkable strength-ductility tradeoff dilemma among these wrought Mg alloys.For example,Pan et al.[23] developed an ultra-high strength Mg-2Sn-2Ca(wt.%) alloy at RT showing a TYS of～443 MPa,but theεfof this alloy was only～1.2%.Also,Zhang et al.[24] revealed that in an extruded Mg-1.0Zn-0.5Ca (wt.%) alloy the achievement of highεf(～44%)was at the expense of the loss of TYS (only～105 MPa).In this work,the newly developed BM11–220 and BM12–220 alloys have a higher strengthductility synergy,Namely,the BM11–220 alloy shows a TYS of～252.1 MPa,a UTS of～281.7 MPa and anεfof 26%;the BM12–220 alloy has a TYS of～319.2 MPa,a UTS of～337.3 MPa and anεfof 13% (see Table 2).Fig.1c and d demonstrates the correlations between the TYS (and theεf)and the alloying content of the extruded BM series alloys and the above mentioned RE-free wrought Mg alloys.Generally,as the alloying content increases,the TYS gradually improves being illustrated better in Fig.1c.As for the reported RE-free wrought Mg alloys,it is necessary for obtaining the TYS beyond 200 MPa to add total alloying content up to～6% at least.However,the TYS above 200 MPa in the extruded BM series alloys can be achieved by only adding～2–3% total alloying content.Theεfappears to have no connection with the alloying content (see Fig.1d).The high-content wrought Mg alloys still present the highεfvalues,in spite of there exist a lot of second phases as tensile cracking sources being harmful to theεf,as reported by Sasaki et al.[25].From the consideration of cost,the low-content BM series alloys especially for the BM11–220 and BM12–220 alloys showing an appropriateεfwith the relatively high strength provide a huge potential for the more extensive industrial applications.

3.2.Microstructures and micro-textures

3.2.1.Solid-solution treated alloys

Fig.2a shows the SEM image of the B1 alloy in solidsolution state.A small amount of residual second phases are never dissolved into the Mg matrix during solid-solution treatment.Based on the SEM-EDS results of points A and B shown Fig.2a (see Fig.2b),these un-dissolved second phases are identified as Mg3Bi2.Similarly,a few un-dissolved Mg3Bi2phases are also found in the solid-solution treated BM11 and BM12 alloys,as shown in Fig.2c and d.Because of the low solid solubility of Mn element in the Mg matrix,α-Mn phases generally can be measured in as-cast state and exist in a form of elementary substance,i.e.,in Mg-4Zn-0.6Ca-0.7Mn (wt.%) alloy [26].After solid-solution treatment,there are still someα-Mn particle phases observed in the BM11 and BM12 alloys.Solid-solution temperature of 500°C can hardly make residualα-Mn phases completely dissolve into the Mg matrix.The average sizes of these observed un-dissolved Mg3Bi2andα-Mn phases are about 3～4μm.The number fractions are～0.5% in three alloys in solidsolution state.By TEM view,we never observe any nanoscale second phases in the solid-solution treated BM11 and BM12 alloys,as shown in Fig.2e and f.

3.2.2.Extruded alloys

Fig.3 exhibits the EBSD inverse pole figure (IPF) maps and (0002) pole figure (PF) maps of the extruded B1,BM11 and BM12 series alloys from cross section.It is well known that hot-extrusion can refine the grain sizes of alloys via the DRX process.The extrusion parameters (i.e.,temperature and speed) tend to be vital factors to impact the microstructural feature of DRX.Under the condition of same extrusion speed,increasing the extrusion temperature can bring about grain coarsening [27].This statement is fairly consistent with the microstructural observation results of extruded alloys in this work,where the average grain size shows an increasing trend with the extrusion temperature from 220 to 300 °C reflected in Fig.4a.The corresponding data are listed in Table 3.The average grain sizes of the extruded B1 series alloys seem to be more sensitive to the extrusion temperature than the extruded BM11 and BM12 series alloys,since it exhibits a larger increment in the extruded B1 series alloys with the extrusion temperature.In other words,the addition of Mn remarkably weakens the sensitivity of extrusion temperature to the grain size.At the same extrusion temperature,the average grain sizes of the extruded BM11 and BM12 series alloys tend to be lower than these of the extruded B1 series alloys,and as the Mn content adds from 1 to 2 wt.% the average grain size further reduces.

Fig.4.(a) Average grain sizes,(b) number fractions of “coarse”-grained structure and (c) number fractions of low angle grain boundaries (LAGBs)in the extruded B1,BM11 and BM12 series alloys.

Fig.5.(a-c) SEM images from longitudinal section of the B1–220,BM11–220 and BM12–220 alloys showing the distribution of second phases,respectively;(e-f) SEM-EDS results corresponding to second phases labeled in (a-c),respectively.

In addition,note that bimodal-grained structure is increasingly obvious with the extrusion temperature in extruded alloys.In the extruded B1 series alloys,this structure mainly consists of the fine and coarse DRXed grains (coarse DRXed grains are marked by white arrows in Fig.3),while in the extruded BM11 and BM12 series alloys it is mainly composed of the fine DRXed grains and the coarse un-DRXed regions (coarse unDRXed regions are labeled by red dotted circles in Fig.3).We define the segment beyond the average grain size as the “coarse”-grained microstructure.The number fraction of “coarse”-grained microstructure in each extruded alloy is shown in Fig.4b and Table 3.It can be seen that the number fractions of“coarse”-grained microstructures in the extruded B1 series alloys far exceed these in the extruded BM11 and BM12 series alloys,meaning that the Mn addition can effectively decreases the microstructural homogeneity of extruded alloys.However,when the Mn content adds from 1 to 2 wt.%,such a microstructural homogeneity seems to be deteriorated.The number fraction of low angle grain boundaries (LAGBs) labeled by white lines in Fig.3 in each extruded alloy is shown Fig.4c and Table 3.Obviously,with increasing Mn content,the number fraction of coarse unDRXed regions gradually increases,along with the high proportion of LAGBs among grain boundaries (GBs).This result suggests that the Mn addition delays the DRX process during extrusion.

Micro-texture is also one of crucial microstructure features in Mg alloys with a hexagonal close-packed (hcp) structure.All the extruded B1,BM11 and BM12 series alloys show a typical basal fiber texture characteristic with c-axes of most of grains perpendicular to the extrusion direction (ED) in the(0002) PF maps,and they have a common tendency for the texture intensity to strengthen with the extrusion temperature.Numerous researches have reported that the high DRX level and the fine grain size gave rise to the low texture intensity [28,29].It is worth noting that even though the extruded B1 series alloys show the higher DRX fractions,the larger average grain sizes still make the texture intensity become stronger than the extruded BM11 and BM12 series alloys at the same extrusion temperature.However,as the Mn content adds from 1 to 2 wt.%,the texture intensity in turn hardly alters,although the Mn addition reduces the grain size.

3.3.Second phases in extruded alloys

Fig.6.TEM images showing dynamic precipitates of the B1–220,BM11–220 and BM12–220 alloys: (a,d,g) views of low magnification in the B1–220,BM11–220 and BM12–220 alloys,respectively.Green and red arrows mark irregular-shaped and rod-shaped precipitates,respectively;(b,e,h) views of high magnification inside grains in the B1–220,BM11–220 and BM12–220 alloys,respectively.Blue circles label spherical precipitates;(c,f,i) views of high magnification at grain boundaries in the B1–220,BM11–220 and BM12–220 alloys,respectively.Blue and black arrows present spherical precipitates and grain boundaries,respectively.

Fig.5a-c shows the second phase features in the extruded B1–220,BM11–220 and BM-220 alloys through SEM images.After extrusion,those scattered un-dissolved second phases in solid-solution state are broken into small clusters with the average size of～2μm,and they are in chain along the ED.By SEM-EDS analysis,these small clusters consist of Mg3Bi2phases in the extruded B1–220 alloy,and the coexistence of Mg3Bi2andα-Mn phases occurs in the extruded BM11–220 and BM12–220 alloys (see Fig.5d-f).Besides these “l(fā)arge size” phases,a certain amount of “small size”dynamic precipitates also can be found by TEM observation.Taking the B1–220,BM11–220 and BM12–220 alloys as examples,Fig.6 shows the morphologies,the sizes,the number fractions and the distributions of “small size” dynamic precipitates in these extruded alloys.In the B1–220 alloy,two types of precipitates with the irregular and rod shapes(labeled by green and red arrows,respectively) are observed by TEM view of low magnification (see Fig.6a).They distribute randomly on the matrix.The irregular-shaped ones have an average diameter of～100 nm and a number fraction of～4.5%.The rod-shaped items show an average length of～600 nm and an average width of～30 nm,and they account for～3.2%.Fig.6b and c corresponds to the high-magnified TEM images located inside the grain and at the GB,respectively,where no any finer nano-scale phases are detected.In both the BM11–220 and BM12–220 alloys,we still find that there are the similar characteristics of precipitates with the B1–220 alloy,as shown in Fig.6d and g.However,the third type of spherical precipitates marked by blue circles is measured at the high-magnified TEM images.They not only distribute dispersedly on the matrix (see Fig.6e and h),but also segregate at the GBs (see Fig.6f and i).Both the BM11–220 and BM12–220 alloys exhibit the similar size of～10 nm for these spherical precipitates,but their number fractions are～4.6 and～7.3% for the BM11–220 and BM12–220 alloys,respectively.

Fig.7.(a-c) TEM images and SAED patterns of the irregular-shaped,rod-shaped and spherical precipitates at the high magnification in the BM12–220 alloy,and the corresponding measured items marked by A,B and C,respectively;(d-f) TEM-EDX results of precipitates labeled by A,B and C,respectively;(g,h) XRD results ranged from 41 to 45° of all extruded alloys and the corresponding highlighted observation of diffraction peaks of α-Mn phase,respectively.

The BM12–220 alloy,as an experimental group,is selected to obtain the precise chemical compositions of the above mentioned three kinds of precipitates by TEM-EDX measurement and selected area electron diffraction (SAED) pattern,as shown in Fig.7.Fig.7a-c exhibits the TEM images and the SAED patterns of the irregular-shaped,rod-shaped and spherical precipitates at high magnification,and the corresponding measured items are marked by circles A,B and C,respectively.TEM-EDX results display that the irregularshaped(marked by green circle A)and rod-shaped(labeled by red circle B) precipitates may be made up of the same chemical constitutions including Mg,Bi and Mn elements,where Mg and Bi are the main elements,and Mn is minimal (see Fig.7d and e).From SAED results,we can confirm that both the irregular-shaped and rod-shaped precipitates are Mg3Bi2(hcp structure,where a and c are 0.4667 and 0.7401 nm [8],respectively),as shown in Fig.7a and b.This result indicates there are two orientations of existence for Mg3Bi2phase in the extruded BM series alloys.It is consistent with the report of Sun et al.[30].In addition,the nano-scale spherical precipitate (labeled by blue circle C) is detected for Mg and Mn elements (see Fig.7f).Since its size is extremely small(～10 nm),it is inevitable to introduce Mg element during TEM-EDX measurement.This deviation was also found in the study of Peng et al.[15].From SAED result,we can clearly find two diffraction patterns with the spherical precipitate and the Mg matrix showing an orientation relationship of [11–20]Mg//[001]Mnand (0001)Mg//(111)Mn(see Fig.7c),which further verify a fact that these nano-scale spherical precipitates areα-Mn phases.Besides micro-level TEM-EDX measurement,macro-scale XRD analysis is also carried out for all extruded alloys in Fig.7g.The extruded B1 series alloys contain two phase constitutions:α-Mg and Mg3Bi2.With the addition of Mn,α-Mn phases occur in the extruded BM series alloys,which in accordant with the above mentioned TEM-EDX results.The diffraction peak intensities of Mg3Bi2phases are similar in all extruded alloys,while forα-Mn phases,the peak intensities in the extruded BM12 series alloys are higher than in the extruded BM11 series alloys shown in Fig.7h.This indicates that the volume fraction ofα-Mn phases is gradually large with the addition of Mn.

4.Discussion

The mentioned above show that the extrusion temperature and the Mn content have a significant impact on the microstructures and tensile mechanical properties of the extruded Mg-Bi alloys at RT.Low-temperature extrusion promotes the DRX process and refines the grain sizes of alloys,which further leads to the reduced texture intensity.The addition of Mn is also beneficial for the grain refinement,while it inevitably accelerates the more unDRXed regions.With respect to second phases,the extruded BM series alloys contain“l(fā)arge size” un-dissolved phases and “small size” dynamic precipitates.With the increase of Mn content,the number fraction of nano-scaleα-Mn precipitates gradually increases,and partialα-Mn particles can segregate the GBs.These observed results directly give rise to the variation of tensile mechanical properties in extruded alloys.To further explore the effects of extrusion temperature and Mn addition on the microstructures and mechanical properties of the extruded Mg-Bi alloys,it is necessary to highlight two crucial aspects: 1) DRX behavior,and 2) strengthening-plasticity mechanisms.

4.1.DRX behavior

4.1.1.Microstructure evolution

Compared with face-centered cubic structural metals (i.e.,aluminum alloys),Mg alloys are easier to induce the DRX behavior under the hot deformation condition because of their finite slip systems and low staking fault energy [31].In order to better understand the DRX behavior of the BM series alloys,we take the BM12–220 alloy as an example to illustrate its microstructure evolution during extrusion,as shown in Fig.8.

(a) The solid-solution treated alloy shows an approximately equiaxial-grained microstructure with the serrated GBs,along with a few un-dissolved Mg3Bi2andα-Mn phases distributed at the GBs and inside the grains,before extrusion (see Fig.8a).

(b) {10–12}extension twins are active in the original grains at the early stage of extrusion.These twins reorient the corresponding original grains via～86° orientation (see Fig.8b).No DRX occurs within {10–12} twins during extrusion due to the high mobility of {10–12} twin boundary,but these twins facilitate the subsequent continuous DRX (CDRX) process by the consumption of original grains [32].

(c) With the development of extrusion process,those large un-dissolved second phases are crushed to many small clusters (beyond 1μm) distributed along the ED.The clusters accumulate high density of dislocations around them,providing the driving force of DRX,as shown in Fig.8c.It has been reported that the clusters are more effective than single particles in causing the DRX and the static recrystallization [33,34].

Fig.8.Schematic illustration of the DRX process for the BM12 alloy during extrusion at 220 °C.

(d) At the medium stage of extrusion,sufficient dislocation densities drive the preferred DRX nucleation occurred around these clusters,known as particle-stimulated nucleation (PSN) effect [33,34].The formation of sub-GB inside the parent grains depends on dislocation rearrangement by the accumulation of dislocations.The higher boundary diffusion rate enables the dislocation accumulated on sub-GBs to be absorbed into LAGBs finally leading to the formation of high angle grain boundaries (HAGBs) [31].This process is known as CDRX.In addition,GB bulging is also observed.This phenomenon derives from strain induced boundary migration(SIBM)[35].Some new recrystallized grains are nucleated at the initial GBs by the migration of original flat boundaries to the side with high dislocation density,resulting in the occurrence of a crystalline core.The mentioned DRX mechanisms can be reflected in Fig.8d.Besides the DRX behavior,“small size” Mg3Bi2andα-Mn phases are also precipitated and randomly distribute on the Mg matrix,and partial precipitatedα-Mn particles segregate at the GBs impeding the migration of GBs.

(e) When extrusion process completes,tremendous fine DRXed grains occupy the majority of Mg matrix,accompanied with the unDRXed regions (see Fig.8e).This is because those nano-scaleα-Mn particles act as a double-edge role.On the one hand,they strongly prevent the boundaries of DRXed grains from migrating.The increasing volume fraction ofα-Mn particles located at the recrystallized GBs enlarges the curvature of GB migration,resulting in the low migration rate and promoting the grain refinement [36].On the other hands,the pinning effect of original GBs acts a cause of retarding the DRX nucleation.Reducing the migration rate of original GBs goes against the occurrence of discontinuous DRX (DDRX),such as the mentioned SIBM mechanism,thereby making the more unDRXed regions.

According to the relationship between the average grain size and the deformation temperature under the hotdeformation condition [37],it is reasonable that the average grain sizes of extruded alloys gradually increase with the extrusion temperature (see Fig.3).Moreover,we find that there are coarse unDRXed regions embedded into massive fine DRXed grains in the BM series alloys.This may be closely related to the deformation mechanism of alloy during extrusion.At the high temperature of 300 °C,non-basal slips are activated easily in the parent grains.Partial parent grains are inclined to deform (i.e.,via prismatic 〈a〉 slip [19]) for accommodating the strains,except for the DRX process,giving rise to the heterogeneous microstructure (especially for the BM12–300 alloy).In another aspect,increasing the Mn content facilitates the more precipitation of nano-scaleα-Mn particles,which pins the GBs and decreases the GB migration rate.This result causes the finer DRXed grains and the more unDRXed regions in the extruded BM series alloys with high Mn content.

4.1.2.Texture modification

Fig.9.(a-c) Recrystallization maps of the B1–300,BM11–300 and BM12–300 alloys,respectively.Blue and yellow colors represent the DRXed and unDRXed regions,respectively;(d) (0002) PF maps of the total,DRX and unDRX regions in the B1–300,BM11–300 and BM12–300 alloys.

Common Mg alloys,i.e.,AZ31[38],Mg-Mn[39]and Mg-Bi [21] alloys,tend to show a strong 〈10–10〉 fiber texture with the basal planes and〈10–10〉directions of the crystallites parallel to the ED.Unlike some special elements,i.e.,RE elements,they would change the ratio of lattice parameters c and a of Mg alloys,and further activate the more non-basal slips and the double twinning,which offered a potential approach to randomize the recrystallization texture.A large amount of literatures have reported that as regards common Mg alloys without RE elements,with the variation of extrusion temperature their basal fiber texture features never changed,and the only difference was their texture intensities [40–43].The study of Du et al.[43]exhibited that the basal fiber texture intensity gradually reduced with the extrusion temperature from 300 to 350 °C in an extruded Mg-4.5Zn-1.1Ca (wt.%) alloy.They found there were vast elongated grains during extrusion of 300 °C having a stronger<10-10>//ED texture component,leading to texture strengthening.Nevertheless,this result is the opposite in the reports of Borkar et al.[40],Huang et al.[41] and Niu et al.[42] on extruded Mg-1Mn-1Sr,Mg-9Al-1Zn-2Ca and Mg-2Zn-1Al (wt.%) alloys at different temperatures,respectively.In the present study,our result is consistent with the mentioned finding that the basal fiber texture is reinforced with the extrusion temperature (see Fig.3).The grain refinement is one of most important reasons for the low texture intensity induced by decreasing the extrusion temperature,since the finer DRXed grains tend to show the more profound grain orientations.In another respect,these DRXed grains with“soft”orientation easily deform,when they are applied to an extrusion force,to make grain rotate through basal〈a〉 slip,until they show a “hard” orientation (basal fiber texture) of basal 〈a〉 slip.As the extrusion temperature increases,the resistance of grain rotation becomes weak and the basal planes of massive “soft” orientation grains are parallel to the ED,resulting in the enhanced basal fiber texture.

The effect of Mn content on the texture is also considered.Taking the B1–300,BM11–300 and BM12–300 alloys as examples,Fig.9a-c shows the recrystallization maps in the mentioned three alloys.Blue and yellow colors represent the DRXed and unDRXed regions,respectively.From the (0002)PF maps corresponding to the DRXed regions in three alloys shown in Fig.9d,the B1–300 alloy with the large-sized DRXed structure exhibits a strong basal fiber texture,and it is weakened by the addition of Mn in the BM11–300 and BM12–300 alloys.However,obvious preferred basal-oriented components in the coarse unDRXed regions enhance the texture intensities of the BM11–300 and BM12–300 alloys (see Fig.9d).Therefore,under the combined effects of grain refinement and unDRX phenomenon,the total texture intensity of the B1–300 alloy slightly decreases by the addition of 1 wt.% Mn,but it almost never changes even if the Mn further adds to 2 wt.%.

Note that it has widely been reported that the role of the PSN effect was to provide the additional nucleation sites for recrystallization that generated various grain orientations other than the<10-10>//ED orientation [44].However,in this work,the sizes of most of particles are below 1μm (see Fig.3),which they are hard to activate the remarkable PSN effect,so that the randomization induced by the PSN effect can be neglected.

4.2.Strengthening-plasticity mechanisms

The mechanical properties of an Mg alloy are determined by the alloy composition and its microstructure characteristic (i.e.,grain size,texture and second phase).These quantities can indeed also be significantly modified by thermomechanical processing,which gives additional degrees of freedom for material researchers to control the microstructure features and further enhance the mechanical properties of alloys.In this section,the strengthening-plasticity mechanisms of the investigated extruded alloys are discussed.

4.2.1.Strengthening mechanism

In Mg alloys,there are mainly five kinds of strengthening mechanisms including GB strengthening,texture strengthening,precipitation strengthening,solid solution strengthening,and dislocation strengthening.In order to quantify the roles of different strengthening mechanisms on the strength of alloy,the contribution values of these main strengthening mechanisms to the TYS are calculated as follows [45]:

whereσTYSis the tensile yield stress,ΔσGBandΔσTare the enhancements from GBs and grain orientation on the stress,respectively,Mis the Taylor factor (～2.5) [46],τ0is the critical resolved shear stress (CRSS),ΔτP,ΔτssandΔτρa(bǔ)re the CRSSs contributed by precipitation strengthening,solid solution strengthening and dislocation strengthening,respectively.M·(τ0+Δτss) is considered to beσ0that is the intrinsic strength taken which is taken as the as-cast yield strength [26],because solid solution strengthening is limited and not obviously changed after extrusion.Therefore,in this work,theσ0values are～40.5,～43.2 and～46.1 MPa for the extruded B1,BM11 and BM12 series alloys,respectively.However,given that the variation on the strengthening modes in the DRXed and unDRXed regions of extruded alloys,Eq.(1) is divided into two parts as follows:

whereΔσHAGBandΔσLAGBare the contributions of HAGBs in the DRXed regions and LAGBs in the unDRXed regions on the strength,respectively.

(i) Calculation of strengthening mechanisms for the DRXed regions in extruded alloys

With respect to HAGB strengthening,the strengthening increment can be expressed by the modified Hall-Petch equation as follows [47]:

wherekis the strengthening factor,daverageis the average grain size andfHAGBis the number fraction of HAGBs.Thekvalue is about 184 MPa ·μm1/2in basal-oriented AZ31 alloys [48].Combined with the statistical data as regards thedaverageandfHAGBvalues shown in Table 3,the contributions of GB to the TYSs (ΔσHAGB) are calculated as～97.6,～71.7,～64.6,～100.2,～88.1,～75.3,～128.6,～100.8 and～86.5 MPa for the B1–220,B1–250,B1–300,BM11–220,BM11–250,BM11–300,BM12–220,BM12–250 and BM12–300 alloys,respectively.

Basal slip,as the most significant deformation mode,accommodates the strains at RT.In general,using the Schmid factor (SF) for basal slip represents the level of basal slip.The low SF for basal slip suggests that basal slip is so difficult.Reaching the yield strain requires to the higher yield stress in the lower SF alloys.A revised form of Hall-Petch relationship can calculate the texture in the DRXed regions on the TYS by Eq.(5) as follows [49,50]:

In this work,there exist “l(fā)arge size” un-dissolved second phases and “small size” dynamic precipitates.The effect of these “l(fā)arge size” second phases on the yield stress tend to be tiny and can be ignored,according to the Orowan law [51].The contribution of these“small size”precipitates on the yield stress can be estimated.In the extruded B1 series alloys,only irregular-shaped and rod-shaped Mg3Bi2phases are observed.We consider these irregular-shaped phases as approximately a spherical shape of uniform diameterdp=～100 nm.The Orowan increment on the stress (ΔσP,irregular Mg3Bi2) produced by the need for dislocations to by-pass these obstacles is given as follows [51]:

whereGis the shear modulus of the Mg matrix phase (～17 GPa),bis the magnitude of the Burgers vector of the slip dislocations (～0.32 nm),vis the ratio of Poisson (～0.3),andfpis the volume fraction of precipitates.Based on the statistical results for the irregular-shaped Mg3Bi2phases,the precipitation hardening value is calculated as～1.6 MPa for each extruded B1 series alloy,respectively.With respect to the rod-shaped Mg3Bi2phases grew along the [0001]Mgdirection (see Fig.7b),the Orowan criterion for [0001]Mgprecipitate rods (ΔσP,rod Mg3Bi2) assuming that the lengths of precipitatesdp=～ 600 nm far exceed their uniform diameters can be expressed as follows [51]:

Using Eq.(7),the stress increment from the rod-shaped Mg3Bi2phases is～0.2 MPa for each extruded B1 series alloy.Thus,the contribution of precipitated Mg3Bi2phases on the TYS is～1.8 MPa for each extruded B1 series alloy.Compared with the extruded B1 series alloys,the extruded BM series alloys contain the more sphericalα-Mn particles,besides Mg3Bi2phases.Using Eq.(6),the precipitation hardening values (ΔσP,spherical α-Mn) of sphericalα-Mn particles are calculated as～30.1 and～42.2 MPa for each extruded BM11 and BM12 series alloy,respectively.Therefore,combined with the mentioned calculated results,the contributions of dynamic precipitates including Mg3Bi2andα-Mn phases on the TYSs are～31.9 and～44.0 MPa for each extruded BM11 and BM12 series alloy,respectively.As mentioned above,theΔσP,DRXedvalues are～1.8,～31.9 and～44.0 MPa for each B1,BM11 and BM12 series alloy,respectively.

In order to qualitatively observe the effect of precipitates on the strength,Fig.10 shows the interaction between precipitates and dislocations after tension of 3%strain along the ED in the BM11–220 alloy.Fig.10a,g and m shows the HRTEM images of three types of precipitates,and their local-magnified images marked by yellow frames correspond to Fig.10b,h and n,respectively.It can be seen that the Mg lattice deforms near these precipitates.To better explore the Mg lattice distortion derived from the interaction between precipitates and dislocations,the inverse Fast Fourier Transformation (FFT) images obtained from the (0001) and (10–10) planes are shown in Fig.10c,i and o,as well as Fig.10d,j and p,respectively.We can see some incomplete lattice fringes near the interface of Mg matrix and precipitates,indicating that there exist many edge dislocations marked by yellow/red symbols“⊥”,and these precipitates effectively impede the dislocation motion.When the investigated alloy is applied to the tensile strain,dislocation slip is the main deformation mode,since most of grains are “hard” orientations for {10–12} extension twinning.The occurrence of dislocations on the (0001) and(10–10) planes proves that basal 〈a〉 and prismatic 〈a〉 slips are activated simultaneously.To evaluate these local strains from the (0001) and (10–10) planes,GPA analysis from the inverse FFT images are demonstrated in Fig.10e,k and q as well as Fig.10f,l and r,respectively.Obviously,the high local strains occur near the dislocations.The strain (εxx) on the (0001) plane is apparently higher than that (εyy) on the(10–10) plane,indicating that the inhibition of precipitates for basal dislocations is larger than that for prismatic dislocations.Compared with two types of Mg3Bi2phases,the sphericalα-Mn particles more availably hamper the dislocation motion to generate the larger local strains and improve the strength.This observation is highly consistent with the calculated results via the Orowan strengthening.

To sum up,theσTYS,DRXedvalues are～221.4,～185.2,～178.1,～253.8,～241.1,～224.5,～321.1,～288.3 and～263.2 MPa for the B1–220,B1–250,B1–300,BM11–220,BM11–250,BM11–300,BM12–220,BM12–250 and BM12–300 alloys,respectively.

(ii) Calculation of strengthening mechanisms for the un-DRXed regions in extruded alloys

In the unDRXed regions,the strengthening induced by LAGBs is expressed as follows [47]:

whereαis the constant (～0.2) andθLAGBis the misorientation angle (in radian) of LAGBs.Using Eq.(8),the LAGBinduced strengthening values (ΔσLAGB) are～7.3,～5.9,～4.8,～9.4,～11.1,～12.6,～12.3,～14.7 and～15.9 MPa for the B1–220,B1–250,B1–300,BM11–220,BM11–250,BM11–300,BM12–220,BM12–250 and BM12–300 alloys,respectively.

As observed in Fig.4 and Table 3,the Mn addition promotes the more unDRXed regions in the extruded BM11 and BM12 series alloys,especially at the high extrusion temperature.These unDRXed regions tend to show a strong basal fiber texture,leading to a texture strengthening effect.According to Eq.(5),a revised form of Hall-Petch relationship can calculate the texture in the unDRXed regions on the TYS(ΔσT,unDRXed) by Eq.(9) as follows [49,50]:

Fig.10.Interactions between different dynamic precipitates and dislocations in the BM11–220 alloy during tension of 3%: (a,g,m) HRTEM images showing the irregular-shaped Mg3Bi2,rod-shaped Mg3Bi2 and spherical α-Mn particles,respectively;(b,h,n) HRTEM images and FFT images corresponding to yellow frames from (a,g,m),respectively;(c,i,o) inverse FFT images obtained from (0001) planes corresponding to yellow frames from (b,h,n),respectively;(d,j,p) inverse FFT images obtained from (10–10) planes corresponding to yellow frames from (b,h,n),respectively;(e,k,q) GPA strain distributions in(0001) planes corresponding to (c,i,o),respectively;(f,l,r) GPA strain distributions in (10–10) planes corresponding to (d,j,p),respectively.Dislocations are marked by symbols “⊥”. 2601

Considering the uniform distribution of dynamic precipitates in extruded alloys,we assume that the precipitation strengthening effect from the unDRXed regions is equal to that from the DRXed regions.Consequently,theΔσP,unDRXedvalues also are～1.8,～31.9 and～44.0 MPa for each extruded B1,BM11 and BM12 series alloy,respectively.

The residual dislocations in the unDRXed regions can exert an indispensable impact on the TYS during tension by the dislocation strengthening (Δσρ,unDRXed),which is evaluated as follows [52]:

whereρGNDis the geometrically necessary dislocation density.During extrusion,most of grains release the stress via the DRX,while for the unDRXed regions they can accommodate the strain through the residual dislocations that are known as the geometrically necessary dislocations (GNDs).Based on the EBSD data,the averageρGNDvalues for the unDRXed regions are～1.6,～1.3,～1.1,～16.1,～17.8,～18.9,～17.9,～18.2 and～18.6 (× 1014)m-2for the B1–220,B1–250,B1–300,BM11–220,BM11–250,BM11–300,BM12–220,BM12–250 and BM12–300 alloys,respectively.Thus,the contributions of residual dislocations on the TYSs (σρ,unDRXed) can be obtained as～34.4,～31.0,～28.5,～109.5,～114.7,～118.2,～115.1,～116.0 and～117.3 MPa,respectively.In conclusion,theσTYS,unDRXedvalues are～92.9,～86.6,～81.8,～208.7,～218.3,～225.6,～236.9,～244.0 and～248.4 MPa for the B1–220,B1–250,B1–300,BM11–220,BM11–250,BM11–300,BM12–220,BM12–250 and BM12–300 alloys,respectively.

In this sense,the rule of mixture shown in Eq.(11) is employed to predict the average TYSs in extruded alloys[46]:

wherefDRXis the number fraction of DRXed regions.Based on the measured results offDRXedvalues (see Fig.4 and Table 3) and the calculated data on theσTYS,DRXedandσTYS,unDRXed,theσTYSvalues are～218.8,～183.2,～177.1,～252.4,～240.4,～224.6,～315.1,～283.8 and～260.1 MPa for the B1–220,B1–250,B1–300,BM11–220,BM11–250,BM11–300,BM12–220,BM12–250 and BM12–300 alloys,respectively.To better clearly reflect the role of each strengthening mechanism on the TYS,Fig.11a shows the accumulative graph of calculated contribution of each strengthening mechanism on extruded alloys.The corresponding data are listed in Table 4.As we can see,HAGB and texture strengthening mechanisms play the most important roles in the calculated TYS.They are enhanced as the extrusion temperature decreases.In addition,with increasing Mn addition,the contributions of precipitation and dislocation strengthening mechanisms on the TYS remarkably increase,as a result of the occurrence of numerous nano-scaleα-Mn particles and coarse unDRXed regions.Besides,LAGB strengthening is also activated,but its effect seems to be tiny due to the limited number fraction.By comparison of the calculated TYSs with the experimental ones shown in Fig.11b and Table 4,they are almost equal within the margin of error.

Fig.11.(a) Accumulative graph of calculated contribution of each strengthening mechanism on the extruded B1,BM11 and BM12 series alloys;(b)comparison of the calculated yield stresses with the experimental ones.

Table 4 Contributions of strengthening mechanisms including intrinsic strength (σ0),HAGB strengthening (ΔσHAGB),LAGB strengthening (ΔσLAGB),texture strengthening (ΔσT),precipitation strengthening (ΔσP),dislocation strengthening(Δσρ)on the tensile yield stress in each extruded B1,BM11 and BM12 series alloys.Unit: MPa.

4.2.2.Plasticity mechanism

Ductility is another important material property for most of the structural materials.In this section,we emphasize qualitatively the effect of microstructural change induced by the extrusion temperature and the Mn addition on the ductility of extruded alloys.As we know,refining the grain size is an effective approach to improve simultaneously the strength and the ductility of Mg alloys.When the metallic material is applied to an external force,the plastic deformation of grains accommodates the macroscopic strains.The greater the number of grains is,the stronger the ability to coordinate the macroscopic strain gets,along with the lower level of stress concentration.Therefore,the fine-grained materials show the higher work-hardening ability,leading to the higher ductility than the coarse-grained materials.However,it is also well accepted that the ductility of the ultra-fine-grained or nanograined materials tend to be extremely low due to the limited work-hardening ability [53–55].This is attributed to the fully saturated defects formed during processing and the diminished defects activities in the fine grains,which would restrict the crystalline defects accumulations [56].Lu [57] also pointed out that the work-hardening rate decreases obviously when the grain size becomes lower than the critical value of～1μm for most of the metals.Consequently,the ultra-fine-grained alloys only have a low uniform plastic deformation.In this work,decreasing the extrusion temperature refines the average grain size of extruded alloys,thereby making theεfimprove.At the same extrusion temperature,i.e.,220 °C,the Mn addition reduces the average grain size from～3.6 to～2.1μm,leading to theεfof the BM11–220 alloy (～26%) having more than twice as much as that of the B1–220 alloy(～12%).The continuous Mn addition in turn deteriorates theεffrom～26 to～13% in the BM12–220 alloy.This is mainly attributed to the finer grain size (～1.3μm) being lack of enough spaces for dislocation storage,showing the low work-hardening ability.Therefore,we believe that the appropriate grain size(～2μm)in this work can provide the sufficient grain fractions and spaces to coordinate the inter-granular deformation and accumulate the dislocations,giving rise to the enhanced ductility.

In the particle-containing ductile materials,fracture may take place prematurely due to the void formation at secondphase particles,and this could be closely associated with the particle decohesion in soft matrix or the particle cracking in a hard matrix [58].Based on the report of Huang et al.[36],it is assumed that particle is regarded as a spherical shape,and thus a critical stress for the void nucleation is proportional to ther-1/2(rpresents the radius of spherical particle).This means that the interfacial voids need a larger stress in the materials with fine particles.A small particle size minimizes the possibility for multiple slip-band pileups that will increase the local stress and create the fracture.In this work,a very small amount of large-sized particles can be observed,and most of fine nano-scale precipitates distribute on the matrix in extruded alloys.Thus,the effect of second phase on the ductility can be neglected.

Due to the strong basal fiber textures in extruded alloys,the activities of basal slip and {10–12} extension twinning are difficult during tension along the ED.In this case,nonbasal slips,as extra deformation modes,can be launched and participate in the plastic deformation,as the flow stress increases.Taking the BM11–220 alloy as an example,Fig.12 shows the slip activities of the BM11–220 alloy after tension of～8%.Fig.12a-c exhibits the SF maps for basal 〈a〉,prismatic 〈a〉 and pyramidal 〈c+a〉 slips,respectively,in extruded states.High SFs for non-basal slips (beyond 0.4)provide the possibility of non-basal slip activities.After tension,IGMA distribution based on the whole deformed region(see Fig.12d)displays two preferences close to Taylor axes of〈0001〉 and 〈uvt0〉,as shown in Fig.12e.By IGMA analysis in Mg alloys [59],the distribution of Taylor axe of 〈0001〉can be interpreted as a result of prismatic 〈a〉 slip activity,and the formation of 〈uvt0〉 IGMA distribution is attributed to the activation of basal 〈a〉 and/or pyramidal 〈c+a〉 slips.Although pyramidal 〈c+a〉 slip mode has the highest CRSS at RT,decreasing the grain size is beneficial for reducing the ratio of CRSSs between basal 〈a〉 slip and pyramidal 〈c+a〉slip (marked by[60].Koike et al.[60] reported that in a fine-grained AZ31 alloy (～6.5μm) subjected to tension at RT,pyramidal 〈c+a〉 dislocations was found to consist of～40% of the total dislocation density at a yield anisotropy factorof only 1.1 instead of an expected value of 100 obtained from single-crystal experiments,showing a high ductility of～47%.This indicates that pyramidal 〈c+a〉 slip is active in the fine-grained alloys during RT deformation.Bright and dark field TEM images after tension are observed by two-beam diffractions in Fig.12f-i.Based on the g ·b=0(b,Burger vector) invisible criterion [61],the straight dislocation lines (labeled by blue arrows) parallel to the (0002) plane are treated as a result of the basal 〈a〉 slip underg=10–10 (see Fig.12f and g),and the segments of dislocations marked by red arrows can be identified as an indicator of the pyramidal 〈c+a〉 slip underg=0002 (see Fig.12h and i).The observed results suggest that the co-activation of multiple slip modes also provides a key reason for the increased ductility of the BM11–220 alloy.

Besides the mentioned above,twinning is also another important factor to influence the ductility of extruded alloys.Barnett[62,63]has clearly revealed the positive effect of{10–12} extension twin formation on the ductility of alloy during deformation,while {10–11} contraction twins and {10–11}-{10–12} double twins might be responsible for a decreased uniform elongation(because of the combination of strain softening inside twins and the local generation of twin sized voids).Moreover,numerous literatures have also reported that{10–11} contraction and {10–11}-{10–12} double twins as the sources of micro-cracks easily form in the coarse-grained alloys with a strong basal texture,leading to the preferential fracture,when they suffer from tension along the ED [64].Therefore,as for the strong basal-oriented B1–300,BM11–300 and BM12–300 alloys with the obvious “coarse”-grained structure,they show the deteriorated ductility which may be closely related to {10–11} contraction twins and/or {10–11}-{10–12} double twins.The twin-induced micro-cracks easily coalesce and propagate along twin boundaries,resulting in the formation of a relatively large macro-crack.On the contrary,the homogeneous and fine grain structure to some extents restricts the occurrence of {10–11} contraction twins(and/or {10–11}-{10–12} double twins) and twin-induced micro-cracks in the BM11–220 alloy,guaranteeing the relatively high ductility.

Fig.12.(a-c) SF maps for basal 〈a〉,prismatic 〈a〉 and pyramidal 〈c + a〉 slips in the BM11–220 alloy,respectively;(d,e) IPF map and IGMA distribution after tension of 8%,respectively;Bright and dark field TEM images of the BM12–220 alloy after tension of 8% under two-beam diffractions of g=10–10(f,g) and g=0002 (h,i),respectively.

5.Conclusions

In summary,we have successfully fabricated the Mg-1BixMn (x=0,1 and 2 wt.%) extruded alloys at different extrusion temperatures (220,250 and 300 °C).The effects of extrusion temperature and Mn addition on the microstructures and mechanical properties of the extruded Mg-1Bi-xMn alloys are investigated.Some of the main conclusions are shown as follows:

(1) With decreasing extrusion temperature,the average grain sizes of all extruded alloys remarkably reduce,accompanied with the weakening of basal fiber texture intensity and the improvement of microstructural homogeneity.With the addition of Mn,the average grain size further decreases,while the number fraction of low angle grain boundaries becomes larger,along with the more unDRXed regions.This indicates that the Mn addition can delay the DRX process of Mg-Bi alloy.Moreover,the continuous Mn addition never shows an obvious impact on the texture because of the combined effects of grain refinement and unDRXed region increment.

(2) Besides a very small amount of large-sized un-dissolved second phases,two types of Mg3Bi2phases with the irregular and rod shapes can be precipitated randomly in the extruded B1 series alloys.With the Mn addition,some nano-scale sphericalα-Mn particles also can be observed and distributed insides the grains and at the grain boundaries in the extruded BM11 and BM12 series alloys.Theseα-Mn particles at grain boundaries effectively pin the motions of recrystallized grain boundaries and original grain boundaries,playing a crucial role in refining the DRXed grains and forming the un-DRXed regions.

(3) High angle grain boundary strengthening and texture strengthening are the uppermost strengthening mechanisms.Precipitation strengthening and dislocation strengthening become active with the Mn addition.Decreasing the extrusion temperature and increasing the Mn addition enhance the contribution of each strengthening on the yield strengths of extruded alloys,making the extruded BM12–220 alloy reach the highest value of～319.2 MPa,along with the appropriate elongationto-failure of～13%.

(4) The extruded BM11–220 alloy shows the highest elongation-to-failure (～26%) without the obvious loss of yield strength (～ 252.1 MPa),which is closely related to the following aspects: (i) fine grain size(～2μm) provides the high grain fraction and the appropriate work-hardening ability;(ii) multiple slip modes including pyramidal 〈c+a〉 slip can be activated during tension;(iii) uniform and fine grain structure restricts the formation of {10–11} contraction twins(and/or{10–11}-{10–12}double twins)and twininduced micro-cracks.

Declaration of Competing Interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

Acknowledgements

The authors are grateful for the financial support from the National Key Research and Development Program of China(U1764253),and the Chongqing Scientific &Technological Talents Program (KJXX2017002).