《金属材料强韧化与组织调控》教学资源（参考文献）On mechanical properties and superplasticity of Mg–15Al–1Zn alloys processed by reciprocating extrusion

MATERIALS SEIENGE ENGINEERING A ELSEVIER Materials Science and Engineering A 464(2007)76-84 www.elsevier.com/locate/msea On mechanical properties and superplasticity of Mg-15Al-1Zn alloys processed by reciprocating extrusion Shih-Wei Leea,Yu-Liang Chen2,Hsiao-Yun Wang, Chih-Fu Yangb,Jien-Wei Yeh . Department of Materials Science and Engineering.National Tsing Hua University.101.Sec.2.Kuang Fu Rd..Hsinchu,Taiwan bDepartment of Materials Engineering,Tatung University.No.40.Sec.3.Thongshan N.Rd.Thongshan District,Taipei.Taiwan Received 30 November 2005;received in revised form 17 January 2007;accepted 19 January 2007 Abstract A superior combination of specific strength and low-temperature high-strain-rate superplasticity of two-phase Mg-15Al-1Zn alloys could be achieved with reciprocating extrusion directly from as-cast billets.In the as-extruded state,the yield strength and ultimate tensile strength were 306 and 376 MPa,respectively.The optimum elongation of at least 1610%in company with a high m value of 0.7 was obtained at the strain rate of 1x 10-2s-when tested at 325C.The pronounced strengthening mechanism was attributed to the high volume fraction of fine-grained hard Mg7Al phase in the refined a-Mg matrix.The excellent superplasticity at high strain rates was also mainly contributed by the high volume fraction of MgAl2 phase which displays easier grain boundary sliding than a-Mg phase.The apparent activation energy for superplastic flow is thus smaller than that of the boundary diffusion in Mg.In addition,the mechanisms for the filament formation and cooperative grain boundary sliding are discussed. 2007 Elsevier B.V.All rights reserved Keywords:Extrusion:Magnesium alloys:Superplasticity 1.Introduction forming temperature,high forming rate and high superplasticity [141. In materials selection,Mg alloys are very attractive for struc- Because magnesium has a strong oxidation tendency in air at tural applications due to their lowest density among metals.It temperatures above 673 K,the superplastic deformation temper- was reported that a decrease in density is 3-10 times more effi- ature is usually confined to less than 673 K [15].Unfortunately. cient in reducing structural weight than increases in strength, previous investigations suggested that most fine-grained Mg damage and durability tolerance assessment or modulus of elas- alloys at such low temperature displayed their superplasticity ticity [1].However,Mg alloys do not exhibit good strength only at low strain rate ranging from 10-5 to 10-3s-1 [16-18]. corrosion resistance,weldability,and formability as compared Moreover,those Mg alloys reported to exhibit high-strain-rate to Al alloys.As a result,their applications in portable and trans- superplasticity requires high amount of expensive elements, portation products are not so extensive as Al alloys. such as Ga,Y and RE,and the processing of RS+P/M or TMP Recently,superplastic forming has become a promising route [6,7]. for Mg alloys to broaden their applications with the merits of In commercial applications,9 wt%is the limiting Al con- higher formability,strength and toughness than those of cast tent for cast Mg-Al-Zn alloys and 8 wt%is that for wrought parts and coarse-grained wrought products [2-13].However, Mg-Al-Zn alloys.This limitation is due to that high Al content many efforts are still required to realize this technology in a yields a lot of coarse B phase,Mg17Al12,and renders Mg alloys viable way.The key point for this viability is to develop fine- too brittle to be processed or utilized [19].However,eutectic grained alloys with low cost,low density,low forming stress,low Mg-33Al alloys produced by direct extrusion with high extru- sion ratio exhibited a large elongation over 2000%at the strain rate of 2x 10-3 s-testing at 400C,indicating high volume Corresponding author.Tel.:+886 3 5719558;fax:+886 3 5722366. fraction of B phase was beneficial for the superplasticity of Mg E-mail address:jwyeh@mx.nthu.edu.tw (J.-W.Yeh). alloys [20]. 0921-5093/S-see front matter 2007 Elsevier B.V.All rights reserved doi:10.1016j.msea.2007.01.119

Materials Science and Engineering A 464 (2007) 76–84 On mechanical properties and superplasticity of Mg–15Al–1Zn alloys processed by reciprocating extrusion Shih-Wei Lee a, Yu-Liang Chen a, Hsiao-Yun Wang a, Chih-Fu Yang b, Jien-Wei Yeh a,∗ a Department of Materials Science and Engineering, National Tsing Hua University, 101, Sec. 2, Kuang Fu Rd., Hsinchu, Taiwan b Department of Materials Engineering, Tatung University, No. 40, Sec. 3, Jhongshan N. Rd., Jhongshan District, Taipei, Taiwan Received 30 November 2005; received in revised form 17 January 2007; accepted 19 January 2007 Abstract A superior combination of specific strength and low-temperature high-strain-rate superplasticity of two-phase Mg–15Al–1Zn alloys could be achieved with reciprocating extrusion directly from as-cast billets. In the as-extruded state, the yield strength and ultimate tensile strength were 306 and 376 MPa, respectively. The optimum elongation of at least 1610% in company with a high m value of 0.7 was obtained at the strain rate of 1 × 10−2 s−1 when tested at 325 ◦C. The pronounced strengthening mechanism was attributed to the high volume fraction of fine-grained hard Mg17Al12 phase in the refined -Mg matrix. The excellent superplasticity at high strain rates was also mainly contributed by the high volume fraction of Mg17Al12 phase which displays easier grain boundary sliding than -Mg phase. The apparent activation energy for superplastic flow is thus smaller than that of the boundary diffusion in Mg. In addition, the mechanisms for the filament formation and cooperative grain boundary sliding are discussed. © 2007 Elsevier B.V. All rights reserved. Keywords: Extrusion; Magnesium alloys; Superplasticity 1. Introduction In materials selection, Mg alloys are very attractive for struc￾tural applications due to their lowest density among metals. It was reported that a decrease in density is 3–10 times more effi- cient in reducing structural weight than increases in strength, damage and durability tolerance assessment or modulus of elas￾ticity [1]. However, Mg alloys do not exhibit good strength, corrosion resistance, weldability, and formability as compared to Al alloys. As a result, their applications in portable and trans￾portation products are not so extensive as Al alloys. Recently, superplastic forming has become a promising route for Mg alloys to broaden their applications with the merits of higher formability, strength and toughness than those of cast parts and coarse-grained wrought products [2–13]. However, many efforts are still required to realize this technology in a viable way. The key point for this viability is to develop fine￾grained alloys with low cost, low density, low forming stress, low ∗ Corresponding author. Tel.: +886 3 5719558; fax: +886 3 5722366. E-mail address: jwyeh@mx.nthu.edu.tw (J.-W. Yeh). forming temperature, high forming rate and high superplasticity [14]. Because magnesium has a strong oxidation tendency in air at temperatures above 673 K, the superplastic deformation temper￾ature is usually confined to less than 673 K [15]. Unfortunately, previous investigations suggested that most fine-grained Mg alloys at such low temperature displayed their superplasticity only at low strain rate ranging from 10−5 to 10−3 s−1 [16–18]. Moreover, those Mg alloys reported to exhibit high-strain-rate superplasticity requires high amount of expensive elements, such as Ga, Y and RE, and the processing of RS + P/M or TMP [6,7]. In commercial applications, 9 wt% is the limiting Al con￾tent for cast Mg–Al–Zn alloys and 8 wt% is that for wrought Mg–Al–Zn alloys. This limitation is due to that high Al content yields a lot of coarse phase, Mg17Al12, and renders Mg alloys too brittle to be processed or utilized [19]. However, eutectic Mg–33Al alloys produced by direct extrusion with high extru￾sion ratio exhibited a large elongation over 2000% at the strain rate of 2 × 10−3 s−1 testing at 400 ◦C, indicating high volume fraction of phase was beneficial for the superplasticity of Mg alloys [20]. 0921-5093/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2007.01.119

S.-W.Lee et al.Materials Science and Engineering A 464 (2007)76-84 公 Based on the above consideration,the novel grain-refining ther investigation was conducted on a SEM.JEOL JSM-5410. method“reciprocating extrusion”[2l-2刀was exploited to fab- for the distribution of phases and surface morphology,and on ricate Mg-Al-Zn alloys with moderate addition of 15 wt%Al an EPMA,JEOL JXA-8800M,for chemical compositions of the The results shows that a suitable combination of low cost.low matrix and second phases. density.high strength,moderate ductility,low forming stress, Tensile specimens were machined from extruded rods and low forming temperature,high forming rate and large superplas- had the gauge dimensions 25 mm long and 6.25 mm in diameter. tic elongation might be obtained.The alloy is hence promising Tensile tests were performed with an Instron universal machine, for HSRS forming of high performance parts for transporta- modeled 4505,at a crosshead speed of 1.2 mm/min. tion and 3C products.In order to get a deeper understanding The specimens for superplastic test were also machined from of the mechanisms of improved mechanical properties and extruded rods and had the gauge dimensions 10mm long and high-strain-rate superplasticity(HSRSP)of such alloys,refined 4mm in diameter.They were tensile-tested at 275,300 and microstructure and surface morphology of superplastic speci- 325C within a temperature variation +2C and under differ- mens were investigated and discussed. ent constant initial strain rates,1×l0-3,2×l0-3,5×10-3, 1×10-2,2×10-2and5×10-2s-l,on the same Instron 2.Experimental machine. The experimental alloy Mg-15Al-1Zn was prepared by melt- 3.Results ing pure Mg,Al and Zn in a crucible under a protection gas, dry air plus 0.3%SF6 gas and then cast at 680C into a 3.1.Microstructure stainless mold.The billets were machined to have the final dimensions,34 mm in diameter and 100 mm in length,for extru- Fig.2 shows the micrographs of the Mg-15Al-1Zn alloy sion.The chemical composition of the experimental alloy was in the as-cast condition.The microstructure consisted of o-Mg Mg-15.1Al-0.97Zn analyzed by ICP-AES. dendrite phase and interdendrite eutectic in which a-Mg phase The machined cast billets were repetitively extruded for 10 particles were dispersed in the B phase,Mg17Al12,confirmed passes on a reciprocating extruder as shown in Fig.1.The extru- by EPMA.The width of secondary dendrite arm of o phase was sion ratio was 9:1 and the extrusion temperature was 325C.In about 40 um.After 10-pass extrusion,the average grain size of o the first nine passes,the billet was extruded with a back pressure phase was reduced to 3 um,and that of the B phase to 1.8 um as from the opposite ram so that the billet shape was recovered shown in Fig.3.The grain refinement and redistribution of both in the opposite container.Each pass is in fact a combination of phases was resulted from the overall effect of fragmentation, two deformations:an extrusion to form a rod,and a following accumulated dynamic recrystallization and mass flow mixing as compression to recover its original billet shape.Since the plas- discussed in other reciprocally extruded alloys [23,24].Frag- tic strain is the same for extrusion and compression,the total mentation means a grain is disintegrated into several pieces. strain in a pass is twice that of a simple extrusion or compres- This is especially possible during the first few passes.In the first sion.Therefore,the total true strain for each pass is equal to pass,the grains of dendritic a phase and interdendritic B phase 2 x In(9+1)=4.6 [24].This true strain is equivalent to that of are elongated into fibrous shape during passing through the die simple extrusion with a large extrusion ratio around 100:1.In hole and then flow laterally due to the compression by the back the 10th pass,the billet was extruded into a long rod,13 mm in pressure.In the second pass,intersections between adjacent o or diameter,without using the opposite ram. B grains are unavoidable since the plastic flow is not reversible The metallographic specimens were prepared by cutting, relative to the first pass.The grains thus may be further divided cold mounting,grinding and polishing.Cast samples were etched in a reagent:10 ml phosphoric acid +90 ml alcohol,while extruded samples done in a reagent:5ml CH3COOH+2.1g picric acid+5 ml water+35 ml alcohol.They were examined with a microscope,Olympus BX51M.Grain size was measured on micrographs by the linear intercept method.Volume fraction of each phase was measured by the point-counting method.Fur- Die Container B Container A Ram B Ram A 50μm Billet Fig.1.The schematic diagram of the reciprocating extrusion apparatus. Fig.2.Optical microstrucure of Mg-15Al-12n alloy in the as-cast condition

S.-W. Lee et al. / Materials Science and Engineering A 464 (2007) 76–84 77 Based on the above consideration, the novel grain-refining method “reciprocating extrusion” [21–27] was exploited to fab￾ricate Mg–Al–Zn alloys with moderate addition of 15 wt% Al. The results shows that a suitable combination of low cost, low density, high strength, moderate ductility, low forming stress, low forming temperature, high forming rate and large superplas￾tic elongation might be obtained. The alloy is hence promising for HSRS forming of high performance parts for transporta￾tion and 3C products. In order to get a deeper understanding of the mechanisms of improved mechanical properties and high-strain-rate superplasticity (HSRSP) of such alloys, refined microstructure and surface morphology of superplastic speci￾mens were investigated and discussed. 2. Experimental The experimental alloy Mg–15Al–1Zn was prepared by melt￾ing pure Mg, Al and Zn in a crucible under a protection gas, dry air plus 0.3% SF6 gas and then cast at 680 ◦C into a stainless mold. The billets were machined to have the final dimensions, 34 mm in diameter and 100 mm in length, for extru￾sion. The chemical composition of the experimental alloy was Mg–15.1Al–0.97Zn analyzed by ICP-AES. The machined cast billets were repetitively extruded for 10 passes on a reciprocating extruder as shown in Fig. 1. The extru￾sion ratio was 9:1 and the extrusion temperature was 325 ◦C. In the first nine passes, the billet was extruded with a back pressure from the opposite ram so that the billet shape was recovered in the opposite container. Each pass is in fact a combination of two deformations: an extrusion to form a rod, and a following compression to recover its original billet shape. Since the plas￾tic strain is the same for extrusion and compression, the total strain in a pass is twice that of a simple extrusion or compres￾sion. Therefore, the total true strain for each pass is equal to 2 × ln (9 + 1) = 4.6 [24]. This true strain is equivalent to that of simple extrusion with a large extrusion ratio around 100:1. In the 10th pass, the billet was extruded into a long rod, 13 mm in diameter, without using the opposite ram. The metallographic specimens were prepared by cutting, cold mounting, grinding and polishing. Cast samples were etched in a reagent: 10 ml phosphoric acid + 90 ml alcohol, while extruded samples done in a reagent: 5 ml CH3COOH + 2.1 g picric acid + 5 ml water + 35 ml alcohol. They were examined with a microscope, Olympus BX51M. Grain size was measured on micrographs by the linear intercept method. Volume fraction of each phase was measured by the point-counting method. Fur￾Fig. 1. The schematic diagram of the reciprocating extrusion apparatus. ther investigation was conducted on a SEM, JEOL JSM-5410, for the distribution of phases and surface morphology, and on an EPMA, JEOL JXA-8800M, for chemical compositions of the matrix and second phases. Tensile specimens were machined from extruded rods and had the gauge dimensions 25 mm long and 6.25 mm in diameter. Tensile tests were performed with an Instron universal machine, modeled 4505, at a crosshead speed of 1.2 mm/min. The specimens for superplastic test were also machined from extruded rods and had the gauge dimensions 10 mm long and 4 mm in diameter. They were tensile-tested at 275, 300 and 325 ◦C within a temperature variation ±2 ◦C and under differ￾ent constant initial strain rates, 1 × 10−3, 2 × 10−3, 5 × 10−3, 1 × 10−2, 2 × 10−2 and 5 × 10−2 s−1, on the same Instron machine. 3. Results 3.1. Microstructure Fig. 2 shows the micrographs of the Mg–15Al–1Zn alloy in the as-cast condition. The microstructure consisted of -Mg dendrite phase and interdendrite eutectic in which -Mg phase particles were dispersed in the phase, Mg17Al12, confirmed by EPMA. The width of secondary dendrite arm of phase was about 40m. After 10-pass extrusion, the average grain size of phase was reduced to 3m, and that of the phase to 1.8 m as shown in Fig. 3. The grain refinement and redistribution of both phases was resulted from the overall effect of fragmentation, accumulated dynamic recrystallization and mass flow mixing as discussed in other reciprocally extruded alloys [23,24]. Frag￾mentation means a grain is disintegrated into several pieces. This is especially possible during the first few passes. In the first pass, the grains of dendritic phase and interdendritic phase are elongated into fibrous shape during passing through the die hole and then flow laterally due to the compression by the back pressure. In the second pass, intersections between adjacent or grains are unavoidable since the plastic flow is not reversible relative to the first pass. The grains thus may be further divided Fig. 2. Optical microstrucure of Mg–15Al–1Zn alloy in the as-cast condition.

S.-W.Lee et al.Materials Science and Engineering A 464 (2007)76-84 where o is the flow stress of a metal,oo the flow stress if there is no resistance to slip across grain boundaries,K a con- stant and d is the grain size,by which magnesium alloys have much higher K value,about 210kNm-3/2 [30]than the K value, 63kNm-32 [31],of 5000 series Al alloys.In addition,the present Mg-15Al-1Zn alloy processed by reciprocating extru- sion exhibits the best yield strength.ultimate tensile strength and specific strength in Table I and a moderate elongation of 5%.This strengthening is attributable to the dense and uniform fine-grained structure,and the high volume fraction of hard B phase resulting from reciprocating extrusion. Since this two-phase structure could be regarded as an aggre- gate of two phases,o and B,the strength of B phase might be calculated based on isostrain condition [35,36]: 20μm Ototal faCo fBoB (2) Fig.3.Optical microstructure of Mg-15Al-1Zn alloy after 10-pass extrusion. where f and o are volume fraction and strength,respectively. into smaller pieces by large shear.Besides this,many sites are As the volume fractions of B phase for Mg-9Al-1Zn and induced for dynamic recrystallization since dynamics recrystal- Mg-15Al-1Zn were 0.18 and 0.34,respectively,the yield lization is easily activated in grain boundaries and at inclusions strength of each phase for such a grain-size level can be cal- with large strain concentration.Thus,dynamic recrystallization culated to be o=106 MPa and oB=694 MPa by substituting rapidly prevails throughout the whole materials in the first few A-Z-212 MPa and SAI-Z-306 MPa.This passes.As for the redistribution of grains,it is resulted from the demonstrates that B phase being an intermetallic compound was repeated exchange of different extrusion and compression flow stronger than o phase by about six-folds at room temperature. patterns.This lets the whole structure be more uniform,isotropic The strengthening by B phase is thus important in the AZ-series and equiaxed.Furthermore,a dual phase structure is effective Mg alloys. in inhibiting grain growth during reciprocating extrusion.The Mg-15Al-1Zn alloy could be regarded as a two-phase alloy 3.3.Superplasticity since the volume fraction of B phase is about 34%as measured by the point-counting method and in consistent with Al-Mg Fig.4 shows the flow stress as a function of strain rate from binary phase diagram.The well mixing of these two phases 1×10-3to5×10-2s-1for275,300and325C.Fig.5 shows was helpful in stabilizing the grain structure during superplastic the elongation-to-failure as a function of strain rate.The pictures deformation. in Fig.6 display the untested and tested specimens tested at 275, 300 and 325C.It can be seen that the present Mg-15Al-1Zn 3.2.Mechanical properties alloy exhibited excellent HSRSP at 300 and 325C.The max- imum elongation tested at 325C and 1 x 10-2s-was above Table 1 lists the mechanical properties of the reciprocally 1610%in company with a high m value of 0.7.Table 2 com- extruded Mg-15Al-1Zn alloy,in which other high-strength AZ- pares the superplasticity of several Mg-Al alloys with similar or series Mg alloys and commercial medium-strength Al alloys are smaller grain size.It can be seen the superplasticity was largely compared.It can be seen that all fine-grained Mg alloys display improved by more Al addition in combination with reciprocat- pronounced grain-size strengthening.This has been discussed ing extrusion [26].The enhanced mechanism of this excellent and reflected in Hall-Petch relation [28,29]. HSRSP will be demonstrated with cooperative grain boundary sliding(CGBS)and the active grain boundary sliding of B phase a=00+Kd-1/2 1) in Section 4. Table 1 Mechanical properties of the present alloys and several comparative Mg and Al alloys Alloy Method d(μm) dy (MPa) Outs (MPa) E(%) p(g/cm3) aylp Mg-15Al-1Zn RE 2.8 306 376 5.0 1.85 165 Mg-9Al-1Zn RE 3.8 212 328 23.0 1.81 117 Mg-15AI[32 RS+extrusion 365 2.5 AZ80A[15] Extrusion 250 340 1.81 138 AZ91D,T6[8] Extrusion 7.6 261.7 360.6 5.3 1.81 145 Mg-9A-1Zn[331 ECAE 1 277 318 2.5 1.81 154 5083-H34[341 Cold rolled 283 345 9 2.66 106 6061-T6[34] Extrusion 276 310 17 2.70 102

78 S.-W. Lee et al. / Materials Science and Engineering A 464 (2007) 76–84 Fig. 3. Optical microstructure of Mg–15Al–1Zn alloy after 10-pass extrusion. into smaller pieces by large shear. Besides this, many sites are induced for dynamic recrystallization since dynamics recrystal￾lization is easily activated in grain boundaries and at inclusions with large strain concentration. Thus, dynamic recrystallization rapidly prevails throughout the whole materials in the first few passes. As for the redistribution of grains, it is resulted from the repeated exchange of different extrusion and compression flow patterns. This lets the whole structure be more uniform, isotropic and equiaxed. Furthermore, a dual phase structure is effective in inhibiting grain growth during reciprocating extrusion. The Mg–15Al–1Zn alloy could be regarded as a two-phase alloy since the volume fraction of phase is about 34% as measured by the point-counting method and in consistent with Al–Mg binary phase diagram. The well mixing of these two phases was helpful in stabilizing the grain structure during superplastic deformation. 3.2. Mechanical properties Table 1 lists the mechanical properties of the reciprocally extruded Mg–15Al–1Zn alloy, in which other high-strength AZ￾series Mg alloys and commercial medium-strength Al alloys are compared. It can be seen that all fine-grained Mg alloys display pronounced grain-size strengthening. This has been discussed and reflected in Hall–Petch relation [28,29], σ = σ0 + Kd−1/2 (1) where σ is the flow stress of a metal, σ0 the flow stress if there is no resistance to slip across grain boundaries, K a con￾stant and d is the grain size, by which magnesium alloys have much higher K value, about 210 kN m−3/2 [30] than the K value, 63 kN m−3/2 [31], of 5000 series Al alloys. In addition, the present Mg–15Al–1Zn alloy processed by reciprocating extru￾sion exhibits the best yield strength, ultimate tensile strength and specific strength in Table 1 and a moderate elongation of 5%. This strengthening is attributable to the dense and uniform fine-grained structure, and the high volume fraction of hard phase resulting from reciprocating extrusion. Since this two-phase structure could be regarded as an aggre￾gate of two phases, and , the strength of phase might be calculated based on isostrain condition [35,36]: σtotal = fσ + fσ (2) where f and σ are volume fraction and strength, respectively. As the volume fractions of phase for Mg–9Al–1Zn and Mg–15Al–1Zn were 0.18 and 0.34, respectively, the yield strength of each phase for such a grain-size level can be cal￾culated to be σ = 106 MPa and σ = 694 MPa by substituting σMg–9Al–1Zn total = 212 MPa and σMg–15Al–1Zn total = 306 MPa. This demonstrates that phase being an intermetallic compound was stronger than phase by about six-folds at room temperature. The strengthening by phase is thus important in the AZ-series Mg alloys. 3.3. Superplasticity Fig. 4 shows the flow stress as a function of strain rate from 1 × 10−3 to 5 × 10−2 s−1 for 275, 300 and 325 ◦C. Fig. 5 shows the elongation-to-failure as a function of strain rate. The pictures in Fig. 6 display the untested and tested specimens tested at 275, 300 and 325 ◦C. It can be seen that the present Mg–15Al–1Zn alloy exhibited excellent HSRSP at 300 and 325 ◦C. The max￾imum elongation tested at 325 ◦C and 1 × 10−2 s−1 was above 1610% in company with a high m value of 0.7. Table 2 com￾pares the superplasticity of several Mg–Al alloys with similar or smaller grain size. It can be seen the superplasticity was largely improved by more Al addition in combination with reciprocat￾ing extrusion [26]. The enhanced mechanism of this excellent HSRSP will be demonstrated with cooperative grain boundary sliding (CGBS) and the active grain boundary sliding of phase in Section 4. Table 1 Mechanical properties of the present alloys and several comparative Mg and Al alloys Alloy Method d (m) σy (MPa) σuts (MPa) ε (%) ρ (g/cm3) σy/ρ Mg–15Al–1Zn RE 2.8 306 376 5.0 1.85 165 Mg–9Al–1Zn RE 3.8 212 328 23.0 1.81 117 Mg–15Al [32] RS + extrusion 2 – 365 2.5 – – AZ80A [15] Extrusion – 250 340 7 1.81 138 AZ91D, T6 [8] Extrusion 7.6 261.7 360.6 5.3 1.81 145 Mg–9Al–1Zn [33] ECAE 1 277 318 2.5 1.81 154 5083–H34 [34] Cold rolled – 283 345 9 2.66 106 6061–T6 [34] Extrusion – 276 310 17 2.70 102

S.-W.Lee et al.Materials Science and Engineering A 464 (2007)76-84 79 3.4.Surface morphology 100 The surface observation of the deformed sample has been an effective way to realize the grain boundary sliding behavior which is the dominant mechanism of superplasticity [41-46]. 0.25 Fig.8 shows the typical irregular surface structure of the super- ■275°c plastic sample deformed at 325C and 1x 10-2s-1.A plenty of apparent filaments,ligaments and cavities between grains are ● 300°c observed,which reflects large deformations occurred in these ▲325°C regions before rupture.The local deformation resulted in fila- 10 10 ments on the surface of the superplastic tensile test samples was Strain rate,s called as"micro-superplasticity"[47]and had been reported in many aluminum alloys and its MMCs,such as AA2091 [48]. Fig.4.Flow stress as a function of strain rate for the present alloy tested at 275 AA5083[491,AA7475[50]and AA5182/AA6090-25%SiCp 300and325C. [51].However,this phenomenon has been reported in a few 1800 magnesium alloys,such as AZ31 [52].AZ91 [53],Mg-33Al 。-Mg-15A1Znat275C [20]and ZK60 [54].It is also noted that the temperature for 1600 ◆-Mg-15A1-1Znat300c micro-superplasticity is much higher in aluminum alloy than 1400 Mg-15Al-1Zn at 325'C in magnesium alloy.The details of filament formation will be 1200 discussed in the next section. 1000 800 4.Discussion 600 400 4.1.On the formation of filaments 200 The minimun elongation for superplasticity 0 The micro-superplasticity was suggested to correlate with 10 102 10 three possible mechanisms [51]:(1)viscous flow due to par- Strain rate,s tial melting [50,55-60]:(2)single crystalline plasticity [61-63]: (3)superplastic flow in micro-volume [64].However,the high- Fig.5.The variation of elongation-to-failure with strain rate tested at 275,300 est testing temperature,325C,for the present alloy was only and 325C. about 0.64 Tm of pure Mg and much lower than the eutectic The apparent activation energy a could be derived from the temperature of Mg-Al alloy,437C [15].Thus,the viscous following equation [40]: flow resulted from incipient melting was impossible to occur in the present alloy.Moreover,the present alloy did not con- alogo tain the grain refiners,such as Zr and Sc for aluminum alloys, Qa=2.3nR a(1/T) (3) to produce a lot of fine dispersoids,thus the possibility to have local ultrafine grains for micro-superplasticity was very where n is the stress exponent (n=1/m);R the gas constant; low.Therefore,single crystalline plasticity is the most proba- o the flow stress;T is the temperature in Kelvin's scale. ble mechanism to develop the filaments in the present alloy.To Fig.7 shows the plots of logo against 1/T at a constant strain check this,a more detailed examination of surface features was rate of 2x 10-3s-1.The apparent activation energy a was made.The close-up of Fig.8(b)reveals the evidence for sin- 83.8 kJ/mol.which is close to but smaller than the activation gle crystalline plasticity since the folds existed on the surface of energy of boundary diffusion,92 kJ/mol in Mg[16].This implies grains and ligaments was the slip-line steps relieved on the grain that there exists some other mechanism for the energy decrease surface. The mechanism will be also related with the active grain bound- From the fact that ligaments had a high aspect ratio even ary sliding of B phase as discussed in the Section 4. under the high strain rate of 1x 10-2s-1,one could image Table 2 Superplasticity of several Mg-Al alloys Materials Method o(MPa) (s-1) e(%) d(pm) T(C) Mg-15Al-1Zn Reciprocating Extrusion 45.9 1×10-2 1610 2.8 325 0.7 AZ31[37J Extrusion 1×10-3 740 2.9 300 0. AZ61[381 ECAE ~6 8×10-5 ~1100 1 250 0.3 AZ91[39] ECAE(annealed) 25 8×10-5 950 0.5 250 0.5 Mg-33A1[20] Extrusion ~15 33×10-4 2100 2.2 400 0.6-0.8 o is flow stress:is strain rate;er is the elongation to failure:d is the average grain size;T is the testing temperature and m is the strain rate sensitivity

S.-W. Lee et al. / Materials Science and Engineering A 464 (2007) 76–84 79 Fig. 4. Flow stress as a function of strain rate for the present alloy tested at 275, 300 and 325 ◦C. Fig. 5. The variation of elongation-to-failure with strain rate tested at 275, 300 and 325 ◦C. The apparent activation energy Qa could be derived from the following equation [40]: Qa = 2.3nR ∂ logσ ∂(1/T ) (3) where n is the stress exponent (n = 1/m); R the gas constant; σ the flow stress; T is the temperature in Kelvin’s scale. Fig. 7 shows the plots of log σ against 1/T at a constant strain rate of 2 × 10−3 s−1. The apparent activation energy Qa was 83.8 kJ/mol, which is close to but smaller than the activation energy of boundary diffusion, 92 kJ/mol in Mg [16]. This implies that there exists some other mechanism for the energy decrease. The mechanism will be also related with the active grain bound￾ary sliding of phase as discussed in the Section 4. 3.4. Surface morphology The surface observation of the deformed sample has been an effective way to realize the grain boundary sliding behavior which is the dominant mechanism of superplasticity [41–46]. Fig. 8 shows the typical irregular surface structure of the super￾plastic sample deformed at 325 ◦C and 1 × 10−2 s−1. A plenty of apparent filaments, ligaments and cavities between grains are observed, which reflects large deformations occurred in these regions before rupture. The local deformation resulted in fila￾ments on the surface of the superplastic tensile test samples was called as “micro-superplasticity” [47] and had been reported in many aluminum alloys and its MMCs, such as AA2091 [48], AA5083 [49], AA7475 [50] and AA5182/AA6090–25%SiCp [51]. However, this phenomenon has been reported in a few magnesium alloys, such as AZ31 [52], AZ91 [53], Mg–33Al [20] and ZK60 [54]. It is also noted that the temperature for micro-superplasticity is much higher in aluminum alloy than in magnesium alloy. The details of filament formation will be discussed in the next section. 4. Discussion 4.1. On the formation of filaments The micro-superplasticity was suggested to correlate with three possible mechanisms [51]: (1) viscous flow due to par￾tial melting [50,55–60]; (2) single crystalline plasticity [61–63]; (3) superplastic flow in micro-volume [64]. However, the high￾est testing temperature, 325 ◦C, for the present alloy was only about 0.64 Tm of pure Mg and much lower than the eutectic temperature of Mg–Al alloy, 437 ◦C [15]. Thus, the viscous flow resulted from incipient melting was impossible to occur in the present alloy. Moreover, the present alloy did not con￾tain the grain refiners, such as Zr and Sc for aluminum alloys, to produce a lot of fine dispersoids, thus the possibility to have local ultrafine grains for micro-superplasticity was very low. Therefore, single crystalline plasticity is the most proba￾ble mechanism to develop the filaments in the present alloy. To check this, a more detailed examination of surface features was made. The close-up of Fig. 8(b) reveals the evidence for sin￾gle crystalline plasticity since the folds existed on the surface of grains and ligaments was the slip-line steps relieved on the grain surface. From the fact that ligaments had a high aspect ratio even under the high strain rate of 1 × 10−2 s−1, one could image Table 2 Superplasticity of several Mg–Al alloys Materials Method σ (MPa) ε˙ (s−1) ef (%) d (m) T ( ◦C) m Mg–15Al–1Zn Reciprocating Extrusion 45.9 1 × 10−2 1610 2.8 325 0.7 AZ31 [37] Extrusion – 1 × 10−3 ∼740 2.9 300 0. AZ61 [38] ECAE ∼6 8 × 10−5 ∼1100 <1 250 0.3 AZ91 [39] ECAE (annealed) ∼25 8 × 10−5 ∼950 0.5 250 0.5 Mg–33Al [20] Extrusion ∼15 3.3 × 10−4 2100 2.2 400 0.6–0.8 σ is flow stress; ε˙ is strain rate; ef is the elongation to failure; d is the average grain size; T is the testing temperature and m is the strain rate sensitivity.

0 S.-W.Lee et al.Materials Science and Engineering A 464 (2007)76-84 (a) Mg-15Al-1Zn alloys of 10-pass Tested at275℃ Untested =1x10351 516% E=2x103s1 1270% i=5x103s1 728% E=1x102s 370% E=2x102s1 285% E=5x102s 261% 2cm () Mg-15Al-1Zn alloys of 10-pass Tested at300℃ Untested 2=1x108 1093% i=2x103r' 1406% =5x103g1 1087% E=1×102g 986% E=2x102s1 623% E=5x102s1 2cm 417% (c) Mg-15Al-1Zn alloys of 10-pass Tested at325'℃ Untested E=1x103s1 695% e=2x103g1 963% e=5x103s1 1297% E=1×102g1 1610% E=2x103g4 608% E=5x10251 2 cm, 227% Fig.6.Untested and tested Mg-15Al-1Zn specimens tested at:(a)275C:(b)325C:(c)325C. that these regions had underwent a high-strain-rate superplas- thus the amount of B phase.The second is the fact that B tic deformation [43].Based on this,the morphology evolution phase has a melting point of 450C much lower than that of of a grain from equiaxed shape to filament shape on the sur- a phase,about 650C,and hence a higher mobility of dis- face of the superplastic tensile sample is proposed as shown in location during plastic deformation.Fig.10 shows two direct Fig.9.Two reasons conclude that B phase is the most prob- evidences to support the proposed mechanism by examin- able phase contributing these filaments.The first is the fact ing the surface and center regions of the fractured portion that the amount of filaments increases with Al content and which was mounted with the resin and polished to reveal the

80 S.-W. Lee et al. / Materials Science and Engineering A 464 (2007) 76–84 Fig. 6. Untested and tested Mg–15Al–1Zn specimens tested at: (a) 275 ◦C; (b) 325 ◦C; (c) 325 ◦C. that these regions had underwent a high-strain-rate superplas￾tic deformation [43]. Based on this, the morphology evolution of a grain from equiaxed shape to filament shape on the sur￾face of the superplastic tensile sample is proposed as shown in Fig. 9. Two reasons conclude that phase is the most prob￾able phase contributing these filaments. The first is the fact that the amount of filaments increases with Al content and thus the amount of phase. The second is the fact that phase has a melting point of 450 ◦C much lower than that of phase, about 650 ◦C, and hence a higher mobility of dis￾location during plastic deformation. Fig. 10 shows two direct evidences to support the proposed mechanism by examin￾ing the surface and center regions of the fractured portion which was mounted with the resin and polished to reveal the

S.-W.Lee et al.Materials Science and Engineering A 464 (2007)76-84 1.8 (a) tensile direction 1.7- 1.6 edW 1.5- 5 filaments 1.4- Q=83.8 kJ/mol 1.3- 1.2+ 0.00165 0.00170 0.00175 0.00180 0.00185 X3,500 1/T,K Fig.7.log o vs.I/T from 275 to 325C at 2 x 10-3s-1. (b) tensile direction cross-section.In these pictures,several stretched B grains are folds observed. 7 4.2.On the enhancement of HSRSP by B phase surface cavities filament The excellent HSRSP might be explained based on the mech- anism of cooperative grain boundary sliding(CGBS)and the easy grain boundary sliding of B phase.In contrast to the distribution of B phase particles along grain boundaries of o ligament phase in those low-Al content magnesium alloys processed by equal channel angular or hot extrusion [65,66],the present ×7,500 1μm500031 alloy with a large volume fraction of B phase,34%,displayed a two-phase aggregate structure as shown in Fig.3.Thus,in Fig.8.SEM surface morphology of the deformed sample tested at 325C with three-dimensional space,a and B phases are substantially two 1x 10-2s-:(a)in a region near the central portion:(b)a close-up of (a). intervening fine-grained skeletons in which B skeleton was thin- ner and had a smaller grain size than o skeleton.Fig.11 compares the microstructure near and far from the fracture surface of the that cooperative grain boundary gliding mechanism could apply specimen deformed at 325C with 1 x 10-2s-1,in which light similarly to the present two-phase alloy [42,67]. gray phase is B phase and the dark gray one is o phase.Since the In detailed observation,it reveals that B phase tended to flow grains in these two skeletons did not coarsen to the first order to form thicker horizontal striations and cavities are preferred approximation during superplastic deformation,it is proposed to initiate at o-B grain boundaries.This illustrates that B phase cavity slip band tensile direction Fig.9.The schematic evolution for the formation of a filament on the surface of specimens

S.-W. Lee et al. / Materials Science and Engineering A 464 (2007) 76–84 81 Fig. 7. log σ vs. 1/T from 275 to 325 ◦C at 2 × 10−3 s−1. cross-section. In these pictures, several stretched grains are observed. 4.2. On the enhancement of HSRSP by β phase The excellent HSRSP might be explained based on the mech￾anism of cooperative grain boundary sliding (CGBS) and the easy grain boundary sliding of phase. In contrast to the distribution of phase particles along grain boundaries of phase in those low-Al content magnesium alloys processed by equal channel angular or hot extrusion [65,66], the present alloy with a large volume fraction of phase, 34%, displayed a two-phase aggregate structure as shown in Fig. 3. Thus, in three-dimensional space, and phases are substantially two intervening fine-grained skeletons in which skeleton was thin￾ner and had a smaller grain size thanskeleton. Fig. 11 compares the microstructure near and far from the fracture surface of the specimen deformed at 325 ◦C with 1 × 10−2 s−1, in which light gray phase is phase and the dark gray one is phase. Since the grains in these two skeletons did not coarsen to the first order approximation during superplastic deformation, it is proposed Fig. 8. SEM surface morphology of the deformed sample tested at 325 ◦C with 1 × 10−2 s−1: (a) in a region near the central portion; (b) a close-up of (a). that cooperative grain boundary gliding mechanism could apply similarly to the present two-phase alloy [42,67]. In detailed observation, it reveals that phase tended to flow to form thicker horizontal striations and cavities are preferred to initiate at – grain boundaries. This illustrates that phase Fig. 9. The schematic evolution for the formation of a filament on the surface of specimens.

82 S.-W.Lee et al.Materials Science and Engineering A 464 (2007)76-84 (a) resin specimen (a)个 ligamen ligament 20kUX7,500 1m06g29 (b) (b) 2gkU×7,5 1μmg8g627 Fig.11.SEM microstructure of the deformed sample tested to 1610%elongation Fig.10.The ligament on the cross-section of tensile specimens tested at 325C with 1 x 10-2s:(a)near the free surface:(b)in the center region of deformed at 325Cand 1x 10-2s-:(a)in a region far from the fracture surface:(b)in a region near the fracture surface. sample. was more active in grain boundary sliding than o phase.That Combining the mechanism of CGBS and the easier grain means,under the same strain rate.B skeleton behaved with a bet- boundary sliding of B phase,a model for explaining the supe- ter superplasticity than o skeleton,as a result,it was squeezed rior superplasticity of the present two-phase structure over those from longitudinal regions and flowed towards the horizontal AZ alloys with low B volume fraction is proposed as shown regions to compensate o skeleton with the deviated elongation in Fig.12.For the latter case,since most of B phase parti- from B skeleton.Similar observation was reported for two-phase cles lie discretely along the grain boundaries of o phase after Ti-24%Al-11%Nb alloy tested at 790K and 2 x 10-4s-1 and hot working (extrusion)[65],the CGBS becomes ineffective discussed with the operation of two shear surface systems,i.e. and lack for the active grain boundary sliding of B phase.In CGBS mechanism [43,68].It is noted that the stringers of o2 fact,Kashyap [67]considered such cases as quasi-single-phase particles oriented approximately at 30-60 to the tensile axis materials and stated that the defects present at the surface of the in the Ti alloy,which is different from the present observation particles cannot completely relax the elastic strain energy dur- revealing more B phase existing in the horizontal regions,as ing superplastic deformation.For the present case as shown in shown in Fig.11(a).This is expected since the present elonga- Fig.12(b),the CGBS mechanism works with the active grain tion is much larger than that the Ti-24%Al-11%Nb alloy and boundary sliding of B phase.Since the easiness of grain bound- thus more B phase was sent into the horizontal regions by the ary sliding increases in sequence:o-a,a-B and B-B boundaries, operation of two shear surface systems.As for the reason why B phase skeleton acts like the lubricant for a phase skeleton. B phase had better superplasticity than o phase,the low melt- Moreover,as the large strain of B phase skeleton will induce ing point of B phase as compared to o phase is also the main the stress and strain concentrations at the o-B interface,cavi- cause since a test temperature of 325C has reached 0.83Tm of ties are expected to preferentially nucleate at such boundaries B phase.It is thus expected B grains were easier to glide along as shown in Fig.12(a).That is also why a large bunch of inter- their own boundaries and to accommodate stress concentration connected cavities was found near the fracture surface as seen in by climb of dislocations. Fig.11(b)[69]

82 S.-W. Lee et al. / Materials Science and Engineering A 464 (2007) 76–84 Fig. 10. The ligament on the cross-section of tensile specimens tested at 325 ◦C with 1 × 10−2 s−1: (a) near the free surface; (b) in the center region of deformed sample. was more active in grain boundary sliding than phase. That means, under the same strain rate, skeleton behaved with a bet￾ter superplasticity than skeleton, as a result, it was squeezed from longitudinal regions and flowed towards the horizontal regions to compensate skeleton with the deviated elongation from skeleton. Similar observation was reported for two-phase Ti–24%Al–11%Nb alloy tested at 790 K and 2 × 10−4 s−1 and discussed with the operation of two shear surface systems, i.e. CGBS mechanism [43,68]. It is noted that the stringers of 2 particles oriented approximately at 30–60◦ to the tensile axis in the Ti alloy, which is different from the present observation revealing more phase existing in the horizontal regions, as shown in Fig. 11(a). This is expected since the present elonga￾tion is much larger than that the Ti–24%Al–11%Nb alloy and thus more phase was sent into the horizontal regions by the operation of two shear surface systems. As for the reason why phase had better superplasticity than phase, the low melt￾ing point of phase as compared to phase is also the main cause since a test temperature of 325 ◦C has reached 0.83Tm of phase. It is thus expected grains were easier to glide along their own boundaries and to accommodate stress concentration by climb of dislocations. Fig. 11. SEM microstructure of the deformed sample tested to 1610% elongation at 325 ◦C and 1 × 10−2 s−1: (a) in a region far from the fracture surface; (b) in a region near the fracture surface. Combining the mechanism of CGBS and the easier grain boundary sliding of phase, a model for explaining the supe￾rior superplasticity of the present two-phase structure over those AZ alloys with low volume fraction is proposed as shown in Fig. 12. For the latter case, since most of phase parti￾cles lie discretely along the grain boundaries of phase after hot working (extrusion) [65], the CGBS becomes ineffective and lack for the active grain boundary sliding of phase. In fact, Kashyap [67] considered such cases as quasi-single-phase materials and stated that the defects present at the surface of the particles cannot completely relax the elastic strain energy dur￾ing superplastic deformation. For the present case as shown in Fig. 12(b), the CGBS mechanism works with the active grain boundary sliding of phase. Since the easiness of grain bound￾ary sliding increases in sequence:–,–and–boundaries, phase skeleton acts like the lubricant for phase skeleton. Moreover, as the large strain of phase skeleton will induce the stress and strain concentrations at the – interface, cavi￾ties are expected to preferentially nucleate at such boundaries as shown in Fig. 12(a). That is also why a large bunch of inter￾connected cavities was found near the fracture surface as seen in Fig. 11(b) [69].

S.-W.Lee et al.Materials Science and Engineering A 464 (2007)76-84 83 System A (a) System B tensile direction System A System B (b) System B System A B 寸 tensile direction B 0 System A System B Fig.12.The schematic model of cooperative grain boundary sliding in alloys with (a)low B volume fraction and (b)high B volume fraction,respectively. 5.Conclusions specimens are primarily of B phase displaying an excellent plasticity.The cavities preferentially nucleate at the o-B (1)This work has demonstrated that Mg-15Al-1Zn alloy could interface. be refined to get a fine and uniform two-phase structure (6)This Mg-15Al-1Zn alloy has low forming stress,low form- directly from the as-cast billets by the reciprocating extru- ing temperature,high forming rate and high superplasticity sion method. and superior specific strength.It is promising for HSRS (2)The Mg-15Al-1Zn alloy possesses superior yield and forming of high performance parts for transportation and ultimate strengths,306 and 376 MPa,respectively,and a 3C products. moderate elongation of 5%.This is due to its fine-grained microstructure and high volume fraction of hard B phase. Acknowledgment (3)The maximum elongation of Mg-15Al-1Zn alloy was over 1610%in company with a high m value of 0.7 obtained at The authors would like to thank the National Science Council 325C and 1x 10-2s-1.The apparent activation energy of the Republic of China,Taiwan,for financially supporting this for superplastic flow is 83.8 kJ/mol and smaller than the research under Contract No.NSC 93-2216-E-007-035. boundary diffusion in Mg.This decrease is attributable to the smaller activation energy of B phase. (4)The excellent HSRSP is related with cooperative grain References boundary sliding (CGBS)of refined two-phase structure enhanced by the active grain boundary sliding of B phase. [1]W.E.Quist,R.E.Lewis,in:M.E.Fine,E.A.Starke (Eds.),Rapidly Solid- (5)B phase tends to flow into the horizontal region by the oper- ified Powder Aluminum Alloys,ASTM Publication,Philadelphia,1984. Pp.7-38. ation of CGBS.Ligaments in the surface of superplastic [2]I.C.Hsiao,J.C.Huang.Metall.Mater.Trans.A33A(2002)1373-1384

S.-W. Lee et al. / Materials Science and Engineering A 464 (2007) 76–84 83 Fig. 12. The schematic model of cooperative grain boundary sliding in alloys with (a) low volume fraction and (b) high volume fraction, respectively. 5. Conclusions (1) This work has demonstrated that Mg–15Al–1Zn alloy could be refined to get a fine and uniform two-phase structure directly from the as-cast billets by the reciprocating extru￾sion method. (2) The Mg–15Al–1Zn alloy possesses superior yield and ultimate strengths, 306 and 376 MPa, respectively, and a moderate elongation of 5%. This is due to its fine-grained microstructure and high volume fraction of hard phase. (3) The maximum elongation of Mg–15Al–1Zn alloy was over 1610% in company with a high m value of 0.7 obtained at 325 ◦C and 1 × 10−2 s−1. The apparent activation energy for superplastic flow is 83.8 kJ/mol and smaller than the boundary diffusion in Mg. This decrease is attributable to the smaller activation energy of phase. (4) The excellent HSRSP is related with cooperative grain boundary sliding (CGBS) of refined two-phase structure enhanced by the active grain boundary sliding of phase. (5) phase tends to flow into the horizontal region by the oper￾ation of CGBS. Ligaments in the surface of superplastic specimens are primarily of phase displaying an excellent plasticity. The cavities preferentially nucleate at the – interface. (6) This Mg–15Al–1Zn alloy has low forming stress, low form￾ing temperature, high forming rate and high superplasticity and superior specific strength. It is promising for HSRS forming of high performance parts for transportation and 3C products. Acknowledgment The authors would like to thank the National Science Council of the Republic of China, Taiwan, for financially supporting this research under Contract No. NSC 93-2216-E-007-035. References [1] W.E. Quist, R.E. Lewis, in: M.E. Fine, E.A. Starke (Eds.), Rapidly Solid￾ified Powder Aluminum Alloys, ASTM Publication, Philadelphia, 1984, pp. 7–38. [2] I.C. Hsiao, J.C. Huang, Metall. Mater. Trans. A 33A (2002) 1373–1384.

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