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Superplasticity and microstructure in Mg–Gd–Y–Zr rolled sheet
Journal of Alloys and Compounds 485 (2009) 295–299
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jallcom
Superplasticity and microstructure in Mg–Gd–Y–Zr rolled sheet Li Li a,b , Xinming Zhang a,∗ , Yunlai Deng a , Changping Tang a a b
School of Material Science and Engineering, Central South University, Lushan South Road, Changsha 410083, Hunan Province, PR China Department of Mechanical Engineering, Hunan Institute of Technology, Hengyang 421002, PR China
a r t i c l e
i n f o
Article history: Received 21 May 2009 Received in revised form 14 June 2009 Accepted 16 June 2009 Available online 24 June 2009 Keywords: Mg–Gd–Y–Zr alloy Hot rolling Superplastic mechanism Second phases
a b s t r a c t Superplastic behavior and microstructure of the rolled Mg–Gd–Y–Zr alloy sheet with an initial grain size of 66 m were investigated. For the purposes, tensile tests at various temperatures and strain rates were conducted, which revealed that the sheet exhibited a maximum elongation of 380% at 435 ◦ C and 0.0005 s−1 . The high ductility was attributed to grain boundary sliding accommodated by dislocation motion assisted by lattice diffusion. It is suggested from microstructural analysis results that the second phases exhibited the significant effect of pinning grain boundaries, and that the  phase was deformed and the strain was partly transferred from the matrix to the  phase. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Metal sheet forming is an approach to develop thin-walled and lightweight products. Magnesium alloy products have a great potential for aircraft, automotive and electronic industries due to low density and high specific strength [1]. The addition of gadolinium (Gd) and other rare earths (Re) can remarkably improve the heat resistance of magnesium alloy due to solution hardening and precipitation hardening [2]. The mechanical properties meet the challenges of the heat-resistant components in those fields at elevated temperatures that may be up to 200–300 ◦ C. Recently, superplastic forming has been applied to shape magnesium alloys into complex geometries [3]. It can provide higher strength and better reliability in contrast to die casting [4]. So far, although hot rolling was utilized to develop superplastic microstructures of the conventional magnesium alloys [5–8], few research have been conducted on the superplasticity of those containing the heavy RE elements. The aim of this contribution is to explore the superplastic mechanism and microstructural evolution in the Mg–Gd–Y–Zr alloy prepared by hot rolling. 2. Experimental details The experimental material was Mg–9.0Gd–4.0Y–0.4Zr (mass fraction, %) alloy. The as-cast ingot was homogenized at 520 ◦ C for 8 h with subsequent cooling in air (Fig. 1). The billets were rolled at 350 ◦ C with intermediate anneals at 500 ◦ C for 15 min. The rollers were heated by a gasoline torch. The reduction in pass was less than 10%. The sheet with the thickness of 1.3 mm was rolled from original 4 mm.
∗ Corresponding author. Tel.: +86 731 88830265. E-mail address: [email protected] (X. Zhang). 0925-8388/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2009.06.113
To investigate superplastic mechanism, strain rate change tests were carried out to deduce strain rate sensitivity coefficient, m-value. The strain rates ranged from 0.01 s−1 to 0.0001 s−1 at the temperatures from 400 ◦ C to 500 ◦ C in air. The sheet specimens for tensile tests were machined directly from the rolled sheet and had a gauge dimension 10 mm long and 3.5 mm in diameter. The tensile axe was parallel to the rolling direction. The tensile tests were performed with MTS universal testing machine equipped with electrical resistance furnace. The tensile specimen was heated for 1800 s prior to initiation of straining. Optical microscope was used to examine the grain microstructure. The morphology of second phases and fracture surfaces was examined by scanning electron microscopy (SEM). The constituent phases were identified by X-ray diffraction (XRD) and energy dispersive X-ray spectrometry (EDS). Transmission electron microscope (TEM) was also used to investigate the post-deformation microstructures. Foils for TEM were prepared by electro-polishing followed by brief low-energy ion beam milling. The final ion-milling step led to an improvement in the foil surface quality without introducing any detrimental artifacts.
3. Results and discussion 3.1. Superplastic behaviors Undeformed and fractured tensile specimens are shown in Fig. 2(a). The maximum elongation of 380% was achieved from the specimen (indicated by the rectangle) tested at 435 ◦ C and 0.0005 s−1 , in which the diffusional necking took place within the whole gauge length. Fig. 2(b) and (c) shows the true stress–strain curves obtained at various strain rates with the stable temperature of 435 ◦ C and at various temperatures with the stable strain rate of 0.0005 s−1 . It can be observed from Fig. 2(b) that at the stable temperature of 435 ◦ C, and at below the strain rate of 0.0005 s−1 the curves exhibited relatively weaker strain hardening without obvious stress peaks. Fig. 2(c) shows that at the fixed strain rate of 0.0005 s−1 , the
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Fig. 1. Optical microstructure of the homogenized Mg–9.0Gd–4.0Y–0.4Zr billet with the average grain size of 200 m.
Fig. 3. The variations of flow stress (a), and elongation-to-failure (b) as a function of strain rate.
curve at 400 ◦ C displayed more obvious effect of strain hardening than other curves at higher temperatures. The variation of flow stress at a fixed strain of 0.10 as a function of strain rate is plotted in Fig. 3(a). The strain of 0.1 was selected so that grain growth during the initial superplastic flow stage was negligible. It has been demonstrated that the flow stress increased with strain rate in a typical sigmoidal curve as well as the conventional superplastic materials. The m-value is defined as
m=
∂log ∂log ε˙
(1) ε,T
which was measured to be 0.5 at the intermediate strain rate range of the curves for 400–435 ◦ C. Fig. 3(b) shows the variation of elongation-to-failure as a function of strain rate. The maximum elongation corresponded well with the highest m-value. For the rolled Mg–Gd–Y–Zr alloy, the maximum elongation of 380% was achieved at 435 ◦ C and 0.0005 s−1 , corresponding to the high mvalue of 0.56. The m-value of 0.5 implies that the grain boundary sliding (GBS) has made substantial contributions to the superplastic deformation. It has been recognized that GBS is accommodated by the slip assisted by diffusion [9]. In order to understand the mechanism during superplastic process, the superplastic deformation activation energy Q was calculated under constant strain rate by the following equation [10]: Fig. 2. (a) Undeformed and fractured tensile specimens of the rolling Mg–Gd–Y–Zr alloy; (b) true stress–strain curves tested at 0.0005 s−1 ; (c) true stress–strain curves tested at 435 ◦ C.
Q = nR
∂ ln ∂(1/T )
(2)
L. Li et al. / Journal of Alloys and Compounds 485 (2009) 295–299
297
Fig. 4. Activation energy curves of ln(flow stress) vs. 1/T for the superplastic deformation of the rolled Mg–Gd–Y–Zr sheet.
where is the flow stress, n is the stress exponent (n = 1/m), R is the gas constant (R = 8.314 J/K), T is the absolute temperature and ∂ln /∂(1/T) is estimated from the slope of the curve in Fig. 4. It is demonstrated that the activation energy was determined to be 222 kJ mol−1 , which was higher than the activation energy for grain boundary self-diffusion (75 kJ mol−1 ) or lattice diffusion (134 kJ mol−1 ) of magnesium alloys, respectively [11]. 3.2. Microstructures of the homogenized Mg–Gd–Y–Zr billet and the initial sheet Optical microstructure of initial sheet after repeating rolling was shown in Fig. 5, where the average grain size was 66 m (according to d = (dr × dt × dn)1/3 , where dr, dt and dn are the grain sizes of three cross-sections for the sheet). After repeating rolling and annealing, there apparently existed lots of deformation twins and recrystallized grains with some second phases about 1–2 m in both the gain boundaries and the grain interiors. Fig. 6 shows that SEM photographs of the homogenized Mg–Gd–Y–Zr billet and the initial sheet after repeating rolling. The corresponding EDS in Table 1 indicate that the second phases (marked by I) residing in the homogenized billet were RE-rich compounds. Moreover, the results of EDS distinguished two types of the second
Fig. 5. Optical microstructure of the initial sheet after repeating rolling with the average grain size of 66 m.
Fig. 6. SEM micrographs of (a) the homogenized Mg–Gd–Y–Zr billet and (b) the initial sheet after repeating rolling. The corresponding compositions of the second phases identified by EDS are listed in Table 1. Table 1 The corresponding results of EDS in Figs. 6 and 8. Position
Gd (at.%)
Y (at.%)
Zr (at.%)
O (at.%)
Mg (at.%)
I II III IV V VI
23.78 2.22 23.10 10.94 27.15 22.71
1.2 2.13 0.11 1.52 41.08 2.53
1.09 69.12 0.22 0.09 – –
0.57 2.69 – – 1.26 0.35
73.36 23.84 76.57 87.45 30.51 74.41
Fig. 7. Grain structure evolution of the rolled Mg–9.0Gd–4.0Y–0.4Zr sheet under the static annealing at 435 ◦ C.
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L. Li et al. / Journal of Alloys and Compounds 485 (2009) 295–299
Fig. 8. SEM microstructures of the post-deformation Mg–Gd–Y–Zr rolling sheet: (a) the specimen stretched at 435 ◦ C and 0.0005 s−1 , (b) the grip of the specimen, and (c) the fracture surface of the specimen. The corresponding compositions of the second phases identified by EDS are listed in Table 1.
phases in the initial sheet. The spherical particle (marked by II) was zirconium-rich core, because the EDS shows that the spherical particle was enriched in Zirconium that never reacts to other elements in the alloy. Examination of more than five the cuboidal compounds (marked by III) shows that the average atomic percentage for Mg:Gd was equal to 76.13:24.32. This indicates the cuboidal particle was Mg3 Gd compound generating during cast processing [12]. 3.3. Microstructures of post-deformation specimen Fig. 7 shows grain structure evolution of the rolled Mg–9.0Gd–4.0Y–0.4Zr sheet under the static annealing at 435 ◦ C. The grains, after heating at 435 ◦ C for about 2.5 h, still kept in an average size of 74 m, which suggests grain growth speed was expected to be much lower than that in conventional magnesium alloys. Fig. 8 presents SEM images of the fractured specimen stretched at 435 ◦ C and 0.0005 s−1 . It is evident in Fig. 8(a) that the average grain size, 12 m, within the gauge region was much finer than that of the initial sheet. Assuming that the isochronous grains in grip of the specimen has not undergone the influence of stress coarsened, finer grains observed within the gauge region can be attributed to dynamic recrystallization (DRX) during tensile deformation, in contrast to the coarse grains found in the grip (see Fig. 8(b)). The post-deformation equiaxed grain about 10 m was characteristic of GBS and consistent with the m-value of 0.56 [13]. Compared to the microstructure of the initial sheet, as shown in Fig. 5, more second phases appeared as irregular-shaped blocks. The compositions of the second phases in Fig. 8(a) and (b), determined by EDS, are listed in Table 1. From the stoichiometric ratio of the constituent elements, the irregular-shaped phase (marked by IV) was identified as Mg5 (Gd, Y), i.e. the  phase. These results agreed fairly well with the phases identified by XRD as presented in Fig. 9. It is evident that more small diffraction peaks of Mg5 (Gd, Y) marked by “䊉”appeared after tensile deformation at the elevated temperature, which implies that more  phase precipitated during tensile deformation. Except for the  phase, other RE-rich phases with a higher Re content (marked by V and VI) were identified as Mg2 Y3 Gd2 and Mg3 Gd, respectively. In Fig. 8(c), voids caused by the second phases were observed on the fracture surface of the specimen. Fig. 10 shows TEM microstructures of the fractured specimen tested at 435 ◦ C and 0.0005 s−1 . In Fig. 10(a), the grain boundaries (GB) and subgrain boundaries (SGB) are pinned by the second phases. Fig. 10(b) displays that a strip recrystallized grain, with dislocations retaining in the grain, developed between second particles. Meanwhile, dislocations at a high density are observed in the region adjacent to the second particles. It has been recognized that dynamic recrystallization (DRX) proceeds continuously, when dislocations remain in the recrystallized grains
Fig. 9. XRD patterns for (a) the initial sheet and (b) the fractured specimen tested at 435 ◦ C and 0.0005 s−1 .
[14]. The effects of second particles on DRX depend on those sizes. On one hand, the interaction between the large particles (∼1 m) and dislocations during superplastic tensile increases the driving force for recrystallization; On the other hand, the small particles, particularly if closely spaced, may exert a significant pinning effect on grain (subgrain) boundaries. The first effect tends to promote recrystallization, whereas the second tends to hinder recrystallization. In this work, the strip DRX grain occurred between two large second particles about 1 m, suggesting that the first effect might gain the advantage over the second one. In general, GBS occurs readily in magnesium alloys, but becomes blocked or inhibited from further sliding at regions such as triple points or grain boundary ledges. The generation of dislocations, however, can remove these impediments [10]. The occurrence of intergranular phase (Fig. 10(a)) was expected to further block GBS, due to the fact that dislocation gliding (or climbing) is hard to get across phase boundary, which explains that the activation energy for superplastic flow was much higher than the activation energy of grain boundary diffusion or lattice diffusion of magnesium alloy. Fig. 10(c) depicts that dislocation network was distributed within the  phase, which implies that the  phases had undergone deformation along with the matrix. The inset of Fig. 10(c) presents the selected area electron diffraction pattern for the  phase with the zone axis parallel to [1¯ 1 2]ˇ . At the elevated temperature up to 435 ◦ C, the  phase (its melting point is 658 ◦ C, close to that of ˛Mg ) was deformed. Therefore, the strain was partly transferred from the matrix to the  phase. Furthermore, it is noted within the  phase that other phases precipitated, which seems to be Mg2 (Y3 Gd2 ) [15].
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Fig. 10. TEM images of the fractured specimen tested at 435 ◦ C and 0.0005 s−1 . (a) An example of grain boundary and subgrain boundary pinned by second phases; (b) a strip recrystallized grain developed between second phases; (c) dislocation network within the  phase.
4. Conclusions
References
Superplasticity and microstructure of the rolled Mg–Gd–Y–Zr sheet with an initial grain size of 66 m were investigated systematically at various strain rates and deformation temperatures. The results are summarized as follows: The material exhibited a maximum elongation of 380% at 435 ◦ C and 0.0005 s−1 . The maximum m-value was equal to 0.56. The activation energy for the superplastic flow was much higher than that of grain boundary diffusion or lattice diffusion of magnesium. Thus, the predominant superplastic mechanism was grain boundary sliding accommodated by dislocation motion assisted by lattice diffusion, which is characterized by post-deformation equiaxed grains. The second phases facilitate dynamic recrystallization during superplastic deformation. The irregular-shaped  phase was deformed and the strain was partly transferred from the matrix to the  phase.
[1] A.A. Luo, Int. Mater. Rev. 49 (2004) 13–30. [2] X. Yang, Z. Xinming, C. Jianmei, J. Hao, Trans. Nonferrous Met. Soc., China 16 (2006) 1888–1894. [3] J.J. Blandin, Superplast. Adv. Mater. 551–552 (2007) 211–217. [4] B.L. Mordike, T. Ebert, Mater. Sci. Eng. A 302 (2001) 37–45. [5] H. Somekawa, K. Hirai, H. Watanabe, Y. Takigawa, K. Higashi, Mater. Sci. Eng. A 407 (2005) 53–61. [6] H. Fujii, H. Iwasaki, J.K. Araki, Magnesium—Sci. Technol. Appl. 488–489 (2005) 571–574. [7] Y.H. Wei, Mater. Sci. Eng. A 360 (2003) 107–115. [8] W.J. Kim, S.W. Chung, C.S. Chung, D. Kum, Acta Mater. 49 (2001) 3337–3345. [9] O.D. Sherby, J. Wadsworth, Prog. Mater. Sci. 33 (1989) 169–221. [10] X. Wu, Y. Liu, Scripta Mater. 46 (2002) 269–274. [11] H.J. Frost, M.F. Ashby, Deformation Mechanism Maps, Pergamon Press Oxford, 1982. [12] X. Zhang, L. Li, Y. Deng, N. Zhou, J. Alloys Compd. 481 (2009) 296–300. [13] T.G. Langdon, Acta Metal. Mater. 42 (1994) 2437–2443. [14] T. Mohri, M. Mabuchi, M. Nakamura, T. Asahina, H. Iwasaki, T. Aizawa, K. Higashi, Mater. Sci. Eng. A 290 (2000) 139–144. [15] Y. Gao, Q. Wang, J. Gu, Y. Zhao, Y. Tong, Mater. Sci. Eng. A 459 (2007) 117–123.
Acknowledgement The authors would like to appreciate the financial supports from the National Security Basic Research Program (No. 5133001E), China.
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jallcom
Superplasticity and microstructure in Mg–Gd–Y–Zr rolled sheet Li Li a,b , Xinming Zhang a,∗ , Yunlai Deng a , Changping Tang a a b
School of Material Science and Engineering, Central South University, Lushan South Road, Changsha 410083, Hunan Province, PR China Department of Mechanical Engineering, Hunan Institute of Technology, Hengyang 421002, PR China
a r t i c l e
i n f o
Article history: Received 21 May 2009 Received in revised form 14 June 2009 Accepted 16 June 2009 Available online 24 June 2009 Keywords: Mg–Gd–Y–Zr alloy Hot rolling Superplastic mechanism Second phases
a b s t r a c t Superplastic behavior and microstructure of the rolled Mg–Gd–Y–Zr alloy sheet with an initial grain size of 66 m were investigated. For the purposes, tensile tests at various temperatures and strain rates were conducted, which revealed that the sheet exhibited a maximum elongation of 380% at 435 ◦ C and 0.0005 s−1 . The high ductility was attributed to grain boundary sliding accommodated by dislocation motion assisted by lattice diffusion. It is suggested from microstructural analysis results that the second phases exhibited the significant effect of pinning grain boundaries, and that the  phase was deformed and the strain was partly transferred from the matrix to the  phase. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Metal sheet forming is an approach to develop thin-walled and lightweight products. Magnesium alloy products have a great potential for aircraft, automotive and electronic industries due to low density and high specific strength [1]. The addition of gadolinium (Gd) and other rare earths (Re) can remarkably improve the heat resistance of magnesium alloy due to solution hardening and precipitation hardening [2]. The mechanical properties meet the challenges of the heat-resistant components in those fields at elevated temperatures that may be up to 200–300 ◦ C. Recently, superplastic forming has been applied to shape magnesium alloys into complex geometries [3]. It can provide higher strength and better reliability in contrast to die casting [4]. So far, although hot rolling was utilized to develop superplastic microstructures of the conventional magnesium alloys [5–8], few research have been conducted on the superplasticity of those containing the heavy RE elements. The aim of this contribution is to explore the superplastic mechanism and microstructural evolution in the Mg–Gd–Y–Zr alloy prepared by hot rolling. 2. Experimental details The experimental material was Mg–9.0Gd–4.0Y–0.4Zr (mass fraction, %) alloy. The as-cast ingot was homogenized at 520 ◦ C for 8 h with subsequent cooling in air (Fig. 1). The billets were rolled at 350 ◦ C with intermediate anneals at 500 ◦ C for 15 min. The rollers were heated by a gasoline torch. The reduction in pass was less than 10%. The sheet with the thickness of 1.3 mm was rolled from original 4 mm.
∗ Corresponding author. Tel.: +86 731 88830265. E-mail address: [email protected] (X. Zhang). 0925-8388/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2009.06.113
To investigate superplastic mechanism, strain rate change tests were carried out to deduce strain rate sensitivity coefficient, m-value. The strain rates ranged from 0.01 s−1 to 0.0001 s−1 at the temperatures from 400 ◦ C to 500 ◦ C in air. The sheet specimens for tensile tests were machined directly from the rolled sheet and had a gauge dimension 10 mm long and 3.5 mm in diameter. The tensile axe was parallel to the rolling direction. The tensile tests were performed with MTS universal testing machine equipped with electrical resistance furnace. The tensile specimen was heated for 1800 s prior to initiation of straining. Optical microscope was used to examine the grain microstructure. The morphology of second phases and fracture surfaces was examined by scanning electron microscopy (SEM). The constituent phases were identified by X-ray diffraction (XRD) and energy dispersive X-ray spectrometry (EDS). Transmission electron microscope (TEM) was also used to investigate the post-deformation microstructures. Foils for TEM were prepared by electro-polishing followed by brief low-energy ion beam milling. The final ion-milling step led to an improvement in the foil surface quality without introducing any detrimental artifacts.
3. Results and discussion 3.1. Superplastic behaviors Undeformed and fractured tensile specimens are shown in Fig. 2(a). The maximum elongation of 380% was achieved from the specimen (indicated by the rectangle) tested at 435 ◦ C and 0.0005 s−1 , in which the diffusional necking took place within the whole gauge length. Fig. 2(b) and (c) shows the true stress–strain curves obtained at various strain rates with the stable temperature of 435 ◦ C and at various temperatures with the stable strain rate of 0.0005 s−1 . It can be observed from Fig. 2(b) that at the stable temperature of 435 ◦ C, and at below the strain rate of 0.0005 s−1 the curves exhibited relatively weaker strain hardening without obvious stress peaks. Fig. 2(c) shows that at the fixed strain rate of 0.0005 s−1 , the
296
L. Li et al. / Journal of Alloys and Compounds 485 (2009) 295–299
Fig. 1. Optical microstructure of the homogenized Mg–9.0Gd–4.0Y–0.4Zr billet with the average grain size of 200 m.
Fig. 3. The variations of flow stress (a), and elongation-to-failure (b) as a function of strain rate.
curve at 400 ◦ C displayed more obvious effect of strain hardening than other curves at higher temperatures. The variation of flow stress at a fixed strain of 0.10 as a function of strain rate is plotted in Fig. 3(a). The strain of 0.1 was selected so that grain growth during the initial superplastic flow stage was negligible. It has been demonstrated that the flow stress increased with strain rate in a typical sigmoidal curve as well as the conventional superplastic materials. The m-value is defined as
m=
∂log ∂log ε˙
(1) ε,T
which was measured to be 0.5 at the intermediate strain rate range of the curves for 400–435 ◦ C. Fig. 3(b) shows the variation of elongation-to-failure as a function of strain rate. The maximum elongation corresponded well with the highest m-value. For the rolled Mg–Gd–Y–Zr alloy, the maximum elongation of 380% was achieved at 435 ◦ C and 0.0005 s−1 , corresponding to the high mvalue of 0.56. The m-value of 0.5 implies that the grain boundary sliding (GBS) has made substantial contributions to the superplastic deformation. It has been recognized that GBS is accommodated by the slip assisted by diffusion [9]. In order to understand the mechanism during superplastic process, the superplastic deformation activation energy Q was calculated under constant strain rate by the following equation [10]: Fig. 2. (a) Undeformed and fractured tensile specimens of the rolling Mg–Gd–Y–Zr alloy; (b) true stress–strain curves tested at 0.0005 s−1 ; (c) true stress–strain curves tested at 435 ◦ C.
Q = nR
∂ ln ∂(1/T )
(2)
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297
Fig. 4. Activation energy curves of ln(flow stress) vs. 1/T for the superplastic deformation of the rolled Mg–Gd–Y–Zr sheet.
where is the flow stress, n is the stress exponent (n = 1/m), R is the gas constant (R = 8.314 J/K), T is the absolute temperature and ∂ln /∂(1/T) is estimated from the slope of the curve in Fig. 4. It is demonstrated that the activation energy was determined to be 222 kJ mol−1 , which was higher than the activation energy for grain boundary self-diffusion (75 kJ mol−1 ) or lattice diffusion (134 kJ mol−1 ) of magnesium alloys, respectively [11]. 3.2. Microstructures of the homogenized Mg–Gd–Y–Zr billet and the initial sheet Optical microstructure of initial sheet after repeating rolling was shown in Fig. 5, where the average grain size was 66 m (according to d = (dr × dt × dn)1/3 , where dr, dt and dn are the grain sizes of three cross-sections for the sheet). After repeating rolling and annealing, there apparently existed lots of deformation twins and recrystallized grains with some second phases about 1–2 m in both the gain boundaries and the grain interiors. Fig. 6 shows that SEM photographs of the homogenized Mg–Gd–Y–Zr billet and the initial sheet after repeating rolling. The corresponding EDS in Table 1 indicate that the second phases (marked by I) residing in the homogenized billet were RE-rich compounds. Moreover, the results of EDS distinguished two types of the second
Fig. 5. Optical microstructure of the initial sheet after repeating rolling with the average grain size of 66 m.
Fig. 6. SEM micrographs of (a) the homogenized Mg–Gd–Y–Zr billet and (b) the initial sheet after repeating rolling. The corresponding compositions of the second phases identified by EDS are listed in Table 1. Table 1 The corresponding results of EDS in Figs. 6 and 8. Position
Gd (at.%)
Y (at.%)
Zr (at.%)
O (at.%)
Mg (at.%)
I II III IV V VI
23.78 2.22 23.10 10.94 27.15 22.71
1.2 2.13 0.11 1.52 41.08 2.53
1.09 69.12 0.22 0.09 – –
0.57 2.69 – – 1.26 0.35
73.36 23.84 76.57 87.45 30.51 74.41
Fig. 7. Grain structure evolution of the rolled Mg–9.0Gd–4.0Y–0.4Zr sheet under the static annealing at 435 ◦ C.
298
L. Li et al. / Journal of Alloys and Compounds 485 (2009) 295–299
Fig. 8. SEM microstructures of the post-deformation Mg–Gd–Y–Zr rolling sheet: (a) the specimen stretched at 435 ◦ C and 0.0005 s−1 , (b) the grip of the specimen, and (c) the fracture surface of the specimen. The corresponding compositions of the second phases identified by EDS are listed in Table 1.
phases in the initial sheet. The spherical particle (marked by II) was zirconium-rich core, because the EDS shows that the spherical particle was enriched in Zirconium that never reacts to other elements in the alloy. Examination of more than five the cuboidal compounds (marked by III) shows that the average atomic percentage for Mg:Gd was equal to 76.13:24.32. This indicates the cuboidal particle was Mg3 Gd compound generating during cast processing [12]. 3.3. Microstructures of post-deformation specimen Fig. 7 shows grain structure evolution of the rolled Mg–9.0Gd–4.0Y–0.4Zr sheet under the static annealing at 435 ◦ C. The grains, after heating at 435 ◦ C for about 2.5 h, still kept in an average size of 74 m, which suggests grain growth speed was expected to be much lower than that in conventional magnesium alloys. Fig. 8 presents SEM images of the fractured specimen stretched at 435 ◦ C and 0.0005 s−1 . It is evident in Fig. 8(a) that the average grain size, 12 m, within the gauge region was much finer than that of the initial sheet. Assuming that the isochronous grains in grip of the specimen has not undergone the influence of stress coarsened, finer grains observed within the gauge region can be attributed to dynamic recrystallization (DRX) during tensile deformation, in contrast to the coarse grains found in the grip (see Fig. 8(b)). The post-deformation equiaxed grain about 10 m was characteristic of GBS and consistent with the m-value of 0.56 [13]. Compared to the microstructure of the initial sheet, as shown in Fig. 5, more second phases appeared as irregular-shaped blocks. The compositions of the second phases in Fig. 8(a) and (b), determined by EDS, are listed in Table 1. From the stoichiometric ratio of the constituent elements, the irregular-shaped phase (marked by IV) was identified as Mg5 (Gd, Y), i.e. the  phase. These results agreed fairly well with the phases identified by XRD as presented in Fig. 9. It is evident that more small diffraction peaks of Mg5 (Gd, Y) marked by “䊉”appeared after tensile deformation at the elevated temperature, which implies that more  phase precipitated during tensile deformation. Except for the  phase, other RE-rich phases with a higher Re content (marked by V and VI) were identified as Mg2 Y3 Gd2 and Mg3 Gd, respectively. In Fig. 8(c), voids caused by the second phases were observed on the fracture surface of the specimen. Fig. 10 shows TEM microstructures of the fractured specimen tested at 435 ◦ C and 0.0005 s−1 . In Fig. 10(a), the grain boundaries (GB) and subgrain boundaries (SGB) are pinned by the second phases. Fig. 10(b) displays that a strip recrystallized grain, with dislocations retaining in the grain, developed between second particles. Meanwhile, dislocations at a high density are observed in the region adjacent to the second particles. It has been recognized that dynamic recrystallization (DRX) proceeds continuously, when dislocations remain in the recrystallized grains
Fig. 9. XRD patterns for (a) the initial sheet and (b) the fractured specimen tested at 435 ◦ C and 0.0005 s−1 .
[14]. The effects of second particles on DRX depend on those sizes. On one hand, the interaction between the large particles (∼1 m) and dislocations during superplastic tensile increases the driving force for recrystallization; On the other hand, the small particles, particularly if closely spaced, may exert a significant pinning effect on grain (subgrain) boundaries. The first effect tends to promote recrystallization, whereas the second tends to hinder recrystallization. In this work, the strip DRX grain occurred between two large second particles about 1 m, suggesting that the first effect might gain the advantage over the second one. In general, GBS occurs readily in magnesium alloys, but becomes blocked or inhibited from further sliding at regions such as triple points or grain boundary ledges. The generation of dislocations, however, can remove these impediments [10]. The occurrence of intergranular phase (Fig. 10(a)) was expected to further block GBS, due to the fact that dislocation gliding (or climbing) is hard to get across phase boundary, which explains that the activation energy for superplastic flow was much higher than the activation energy of grain boundary diffusion or lattice diffusion of magnesium alloy. Fig. 10(c) depicts that dislocation network was distributed within the  phase, which implies that the  phases had undergone deformation along with the matrix. The inset of Fig. 10(c) presents the selected area electron diffraction pattern for the  phase with the zone axis parallel to [1¯ 1 2]ˇ . At the elevated temperature up to 435 ◦ C, the  phase (its melting point is 658 ◦ C, close to that of ˛Mg ) was deformed. Therefore, the strain was partly transferred from the matrix to the  phase. Furthermore, it is noted within the  phase that other phases precipitated, which seems to be Mg2 (Y3 Gd2 ) [15].
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Fig. 10. TEM images of the fractured specimen tested at 435 ◦ C and 0.0005 s−1 . (a) An example of grain boundary and subgrain boundary pinned by second phases; (b) a strip recrystallized grain developed between second phases; (c) dislocation network within the  phase.
4. Conclusions
References
Superplasticity and microstructure of the rolled Mg–Gd–Y–Zr sheet with an initial grain size of 66 m were investigated systematically at various strain rates and deformation temperatures. The results are summarized as follows: The material exhibited a maximum elongation of 380% at 435 ◦ C and 0.0005 s−1 . The maximum m-value was equal to 0.56. The activation energy for the superplastic flow was much higher than that of grain boundary diffusion or lattice diffusion of magnesium. Thus, the predominant superplastic mechanism was grain boundary sliding accommodated by dislocation motion assisted by lattice diffusion, which is characterized by post-deformation equiaxed grains. The second phases facilitate dynamic recrystallization during superplastic deformation. The irregular-shaped  phase was deformed and the strain was partly transferred from the matrix to the  phase.
[1] A.A. Luo, Int. Mater. Rev. 49 (2004) 13–30. [2] X. Yang, Z. Xinming, C. Jianmei, J. Hao, Trans. Nonferrous Met. Soc., China 16 (2006) 1888–1894. [3] J.J. Blandin, Superplast. Adv. Mater. 551–552 (2007) 211–217. [4] B.L. Mordike, T. Ebert, Mater. Sci. Eng. A 302 (2001) 37–45. [5] H. Somekawa, K. Hirai, H. Watanabe, Y. Takigawa, K. Higashi, Mater. Sci. Eng. A 407 (2005) 53–61. [6] H. Fujii, H. Iwasaki, J.K. Araki, Magnesium—Sci. Technol. Appl. 488–489 (2005) 571–574. [7] Y.H. Wei, Mater. Sci. Eng. A 360 (2003) 107–115. [8] W.J. Kim, S.W. Chung, C.S. Chung, D. Kum, Acta Mater. 49 (2001) 3337–3345. [9] O.D. Sherby, J. Wadsworth, Prog. Mater. Sci. 33 (1989) 169–221. [10] X. Wu, Y. Liu, Scripta Mater. 46 (2002) 269–274. [11] H.J. Frost, M.F. Ashby, Deformation Mechanism Maps, Pergamon Press Oxford, 1982. [12] X. Zhang, L. Li, Y. Deng, N. Zhou, J. Alloys Compd. 481 (2009) 296–300. [13] T.G. Langdon, Acta Metal. Mater. 42 (1994) 2437–2443. [14] T. Mohri, M. Mabuchi, M. Nakamura, T. Asahina, H. Iwasaki, T. Aizawa, K. Higashi, Mater. Sci. Eng. A 290 (2000) 139–144. [15] Y. Gao, Q. Wang, J. Gu, Y. Zhao, Y. Tong, Mater. Sci. Eng. A 459 (2007) 117–123.
Acknowledgement The authors would like to appreciate the financial supports from the National Security Basic Research Program (No. 5133001E), China.