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Development of high-strength magnesium alloys via combined processes of extrusion, rolling and ageing
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Scripta Materialia 64 (2011) 950–953 www.elsevier.com/locate/scriptamat
Development of high-strength magnesium alloys via combined processes of extrusion, rolling and ageing R.G. Li,a J.F. Nie,a,b G.J. Huang,a Y.C. Xina and Q. Liua,⇑ a
School of Materials Science and Engineering, Chongqing University, Chongqing 400045, People’s Republic of China b Department of Materials Engineering, Monash University, Victoria 3800, Australia Received 20 November 2010; revised 26 January 2011; accepted 27 January 2011 Available online 17 February 2011
A remarkably high tensile yield strength of 445 MPa is obtained in a Mg–14Gd–0.5Zr (wt.%) alloy when this alloy is prepared by the combined processes of hot extrusion, cold rolling and ageing. This significant improvement in strength is associated with a denser distribution of precipitates, an intensive basal texture and a refined grain size. Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Mg–Gd–Zr alloy; Cold rolling; Precipitate; Ageing
Magnesium alloys containing Gd have received much attention in recent years [1–18]. The equilibrium solid solubility of Gd in Mg is relatively high [19], and the Mg–Gd system is therefore regarded as an ideal one for developing high-strength alloys via precipitation hardening [2]. However, the 0.2% yield strength that has been achieved in peak-aged samples of the Mg–Gd binary alloys only reach 300 MPa [6], which is significantly lower than that of the corresponding Al alloys. Therefore, considerable efforts have been made to further improve the strength of Mg–Gd–based alloys. One approach for improving strength is to add suitable microalloying elements to enhance the precipitation hardening response during the isothermal ageing treatment [14]. Another approach involves the combined use of thermomechanical processing and ageing. It has been reported [7] that a 0.2% proof strength of about 360 MPa can be achieved in an Mg–16.9Gd–0.51Zr (wt.%) alloy when this alloy is hot rolled and aged. This alloy contains an equiaxed grain of above 100 lm, and it is strengthened predominantly by fine precipitates. A tensile yield strength of 311 MPa has also been achieved in an Mg–10Gd–2Y–0.5Zr (wt.%) alloy that was processed by hot extrusion and ageing [8]. In this case, the alloy contains equiaxed grains of 4–7 lm, and the grain boundary strengthening and precipitation hardening contribute to the alloy’s strength.
⇑ Corresponding author. Fax: +86 23 65111295; e-mail: qingliu@cqu. edu.cn
More recently, a value as high as 473 MPa has been reported for the 0.2% proof strength in a Mg–10Gd– 5.7Y–1.62Zn–0.65Zr (wt.%) alloy that was processed by hot extrusion and ageing [15]. The grain size achieved in this alloy is about 1.1 lm, which is considerably smaller than those typically obtained in magnesium extrusion alloys. In addition, the dense distribution of precipitates inside the individual grains and along the grain boundaries also contributes to the ultrahigh strength achieved in this alloy [15]. Cold rolling of magnesium alloys can introduce a high density of dislocations, which may facilitate precipitation and therefore enhance the age hardening response. Therefore, the application of cold rolling after hot extrusion and before artificial ageing is a useful process in the development of high-strength magnesium alloys. Here we report that an ultrahigh 0.2% proof strength of over 440 MPa, together with an ultimate strength of 482 MPa, can be achieved in an Mg– 14Gd–0.5Zr (wt.%) alloy that is processed by the combined use of hot extrusion, cold rolling and ageing. The total concentration of alloying elements in this alloy is much lower than that in the Mg–10Gd–5.7Y–1.62Zn– 0.65Zr (wt.%) alloy [15]. An alloy of nominal composition Mg–14Gd–0.5Zr (wt.%)/Mg–2.67Gd–0.15Zr (at.%) was prepared from high-purity (99.9%) Mg and Gd and an Mg–33Zr (wt.%) master alloy by induction melting in a mild steel crucible at 775 °C and casting into a steel mold under mixed SF6 and CO2. The cast billet was 230 mm in
1359-6462/$ - see front matter Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2011.01.042
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Table 1. Tensile properties of the extruded, cold-rolled and aged samples.
Figure 1. (a) Ageing curves of the extruded and cold-rolled samples. Engineering stress–strain tensile curves obtained from (b) the extruded and cold-rolled samples and (c) the extruded, cold-rolled and aged samples.
diameter and 630 mm in height. It was machined into cylindrical-shaped samples with a height of 100 mm and a diameter of 80 mm. The billet was homogenized at 505 °C for 24 h and subsequently extruded at 505 °C, with an extrusion ratio of 28:1 and an extrusion rate of 0.6 mm s1, followed by quenching in water of room temperature. The extruded cylindrical bars have a diameter of 16 mm and a length of 250 mm, and they were denoted as E samples. Some extruded samples were also machined into plates 7 mm wide and 2 mm thick, which were subsequently cold, rolled with total thickness reductions of 10%, 19% and 27% and 10% reduction in each pass. They were denoted as E + 10%R, E + 19%R and E + 27%R, respectively. Cracks started to appear at the edges of samples when the amount of cold work was over 27%. Some of the extruded and cold-rolled samples were also aged and letter A was used to represent aged. Vickers hardness testing was carried out using a 1 kg load and a holding time of 15 s at MH-5. The tensile tests were performed at room temperature using a Shimadzu AG-X (10 kN) machine at a strain rate of 0.001 s1. The tensile direction was parallel to the extrusion direction and rolling direction. The microstructures were observed using an FEI NOVA400 scanning electron microscope operating at 20 kV. Electron backscattered diffraction (EBSD) analysis was performed using a step size of 0.5 lm on an FEI Nova 400 equipped with an HKL-EBSD system. Thin foils for transmission electron microscopy (TEM) observations were prepared by low-energy ion beam thinning. The precipitate characterization was carried out using a LIBRA 200 transmission electron microscope operating at 200 kV. Figure 1a shows the hardness curves of the extruded and cold-rolled samples obtained during isothermal ageing at 200 °C. The hardness value is increased by about 15 HV after 10% cold work, but further increasing the amount of cold work from 10% to 27% leads to an increase in hardness of only about 5 HV. The age hardening response of the cold-rolled samples is enhanced compared with the as-extruded sample. The time to reach peak hardness is about 16 h for all cold-rolled samples, while it is about 36 h for the extruded samples. The maximum hardness achievable in the two processing conditions is about 125 HV for the cold-rolled samples and 122 HV for the extruded samples. Figure 1b shows the engineering stress–strain tensile curves obtained from the extruded and cold-rolled samples. The tensile properties obtained from these stress– strain curves are presented in Table 1. Sample E has a 0.2% proof strength of 190 MPa and an elongation to
Alloys
0.2% Proof stress (MPa)
Ultimate tensile strength (MPa)
Elongation (%)
E E + 10%R E + 19%R E + 27%R E+A E + 10%R + A E + 19%R + A E + 27%R + A
190 255 295 305 305 375 420 445
295 300 350 375 446 470 477 482
19.5 12.0 4.5 3.5 7.0 3.0 2.3 2.0
fracture of 19.5%. The cold rolling significantly increases the yield stress at the expenses of the elongation to fracture. The 0.2% proof strength values of 255, 295 and 305 MPa are obtained in the E + 10%R, E + 19%R and E + 27%R conditions, respectively. The tensile properties of the extruded, cold-rolled and aged samples are provided in Figure 1c and Table 1. The ageing treatments greatly enhance the strength. The increments in the 0.2% proof strength are in the range 115–140 MPa. Note that the 0.2% proof strength values of the cold-rolled samples are much higher than that of the samples without any cold rolling. The 27% coldrolled and aged sample exhibits a 0.2% proof strength of 445 MPa, which is substantially higher than those reported previously in the literature [6–10,20], e.g. 331 MPa in an extruded and aged Mg–10Gd–2Y– 0.5Zr alloy [8], 250 MPa in a hot-rolled and aged Mg– 9.3Gd–4.1Y–0.7Zr alloy [7] and 360 MPa in hot-rolled and aged Mg–12Gd–1.9Y–0.7Zr and Mg–16.9Gd– 0.5Zr alloys [7], Table 2. While the strength level obtained in the present work is significantly higher than the alloys of similar compositions, it is slightly lower than that (473 MPa) reported recently in an extruded and aged Mg–10Gd–5.7Y–1.62Zn–0.65Zr alloy [15]. The latter alloy contains a much higher content of alloying elements, and its grain size (about 1 lm) is substantially smaller than that (20 lm) obtained in the present alloy. According to the Hall–Petch relationship, which is valid for basal slip deformation, a further increase in strength by at least 57 MPa is expected if the grain size is refined from 20 to 1 lm, assuming that the k value is about 140 MPa lm1/2 for grain sizes of about 1 lm [21,22] and about 370 MPa lm1/2 for grain sizes of the order of 20 lm [23,24]. Therefore, a further reduction in grain size via the optimization of the extrusion, cold rolling and ageing processes might potentially lead to the development of an alloy with 0.2% proof strength exceeding 500 MPa. Figures 2 and 3 compare the shape, size and orientation of magnesium grains, together with the information on twins and dislocations inside individual magnesium grains, in the E, E + 10%R and E + 27%R samples. The as-extruded microstructure comprises equiaxed grains, implying that dynamic recrystallization has fully occurred during the extrusion process. The average grain size is about 20 lm (Fig. 2a). The colour of the grains in the IPF map (Fig. 3a) indicates that the texture distribution in the as-extruded sample is reasonably homogeneous and nearly random in the centre and edge
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R. G. Li et al. / Scripta Materialia 64 (2011) 950–953
Table 2. Comparative tensile property and process condition in magnesium alloys with similar compositions. Compositions
Process
Grain size (lm)
0.2% proof stress (MPa)
Ultimate tensile strength (MPa)
Elongation (%)
Mg–11.3Gd–3.8Y–0.7Zr [6] Mg–11.2Gd–3.8Y–0.7Zn–0.7Zr [6] Mg–11.1Gd–3.8Y–2.3Zn–0.7Zr [6] Mg–11.1Gd–4.1Y–1.69Zn–0.49Zr[20] Mg–11Gd–2Nd–0.5Zr [10] Mg–18.23Gd–1.88Ag–0.34Zr [9] Mg–10Gd–2Y–0.5Zr [8]
Cast + T6 Cast + T6 Cast + T6 Cast + T6 Cast + T6 Cast + T6 Cast + T6 Extruded + T5 Extruded + T5 Rolled + T5 Rolled + T5 Rolled + T5 Extruded + Rolled + T5
– – – –
300 300 260 231 230 293 239 331 473 360 360 290 445
330 310 310 361 350 414 362 397 542 400 400 310 482
2.0 <2.0 12.0 4.0 2.3 2.2 4.7 12.8 8 4.0 5.0 16.0 2.0
Mg–10Gd–5.66Y–1.62Zn–0.65Zr [15]. Mg–17Gd–0.51Zr [7] Mg–12Gd–1.9Y–0.69Zr [7] Mg–9.3Gd–4.1Y–0.7Zr [7] Mg–14Gd–0.5Zr (this work)
50 – 43 4 1.1 100 100 100 20
regions of the sample. Twins are not detected within the experimental error. Furthermore, there is little colour variation within individual Mg grains, implying the absence of misorientation caused by sub-grain boundaries and excessive density of dislocations. A weaker extrusion texture has also been observed in a Mg– 1.55 wt.%Gd alloy extruded at 450 °C [17] and other
Figure 3. The (0 0 0 2), (11–20) and (10–10) pole figure maps derived from EBSD of the centre position of the E sample at the plane parallel to extrusion direction and of the centre positions of E + 10%R, E + 19% and the E + 27%R samples at the RD–ND plane.
Figure 2. EBSD IPF maps and Kikuchi band contrast maps of the (a) E, (b) E + 10%R and (c) E + 27%R samples. The horizontal direction is the extrusion direction for sample E, and the rolling direction for the E + 10%R and E + 27%R samples, respectively.
extruded Mg–RE based alloys [25–27]. The random texture and the absence of shear bands in the present alloy are attributable to a higher Gd concentration and a higher extrusion temperature. Figure 2b shows the IPF and Kikuchi band contrast maps of the E + 10%R sample. The grain size and grain shape are nearly the same as those of the as-extruded sample. In the IPF map, the grain colour suggests that basal planes of individual Mg grains are close to being parallel to the rolling surface. Many twins, most of which are tensile twins, and double twins are also observed in the Kikuchi band contrast map, represented by the red and blue lines, respectively, in Figure 2b. Figure 3b shows an intensive basal texture in the RD–ND plane and the centre position of the plate. These observations are similar to those made from cold-rolled AZ31 [28], in which an intensive basal texture forms. The formation of a strong basal texture in the cold-rolled samples of the present alloy could be accounted for by the operation of hai-basal, hai-prismatic
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loy, which is remarkably higher than those achieved in alloys of similar compositions that are produced by a single thermomechanical process of extrusion or rolling. This project was financially supported by the National Basic Research Program of China (“973” Project) (Grant No. 2007CB613703) and the Natural Science Foundation of China (Grant No. 50890172). Figure 4. TEM images of b0 precipitates in matrix in the (a) E + A and (b) E + 27%R + A samples. The electron beam is parallel to the [0 0 0 1] direction of the Mg matrix.
and hc + ai-pyramidal slip systems and the tensile twinning system [28] during the rolling process.Figure 2c shows IPF and Kikuchi band contrast of the E + 27%R sample. The grain shape and the number density of twins are similar to those of the E + 10%R sample, and most of {11–20} and {10–10} planes of Mg tend to be parallel to the normal direction. The texture intensity, shown in Figure 3d, is markedly increased compared to the E + 10%R sample. However, the texture intensity is very similar to that of the E + 19%R sample (Fig. 3c). The observation here is consistent with the previous work on cold-rolled pure magnesium, Mg– 0.2 wt.% Ce and AZ31 alloys, where the cold rolling texture does not change much when the rolling reduction is beyond 10% [29]. A higher density of dislocations in the a-Mg matrix is expected in the E + 27%R sample than that in the E + 10%R sample due to more dislocation slip occurring during rolling. Figure 4 shows TEM images of the E + A and E + 27%R + A samples. A denser distribution of b0 precipitates is noted in the E + 27%R + A sample. The b0 precipitates are about 13.8 nm in length and 7.2 nm in width in the E + A sample, and 8.9 nm in length and 6.8 nm in width in the E + 27%R + A sample, when viewed along the [0 0 0 1] direction. If the equilibrium solid solubility of Gd in Mg (0.6 at.% or 4.0 wt.% at 200 °C) is assumed to be reached in the peak-aged condition in each of the two samples, then the precipitate volume fractions are similar in these two samples and the reduction in precipitate size would lead to an approximately twofold increase in precipitate number density. The increased number density of precipitates in the E + 27%R + A sample is due to the higher dislocation density in the matrix, which facilitates the nucleation of precipitates during the ageing process. The increased number density of precipitates increases the 0.2% proof strength from 305 MPa in the E + A sample to 445 MPa in the E + 27%R + A sample. Meanwhile, the ageing time required to reach peak hardness is decreased from 36 h for the extruded sample to 16 h for the cold-rolled samples. In summary, the combined use of extrusion, cold rolling and ageing processes can significantly enhance the strength of the Mg–14Gd–0.5Zr alloy. The cold rolling after extrusion and before ageing can lead to the formation of stronger basal texture and higher density of precipitates in the Mg matrix, which in turn leads to an enhancement of the yield strength of the alloy. A yield stress of 445 MPa is achieved in the Mg–14Gd–0.5Zr al-
[1] T. Honma, T. Ohkubo, K. Hono, S. Kamado, Mater. Sci. Eng., A 395 (2005) 301. [2] J.F. Nie, X. Gao, S.M. Zhu, Scripta Mater. 53 (2005) 1049. [3] X. Gao, S.M. He, X.Q. Zeng, L.M. Peng, W.J. Ding, J.F. Nie, Mater. Sci. Eng., A 431 (2006) 322. [4] M. Nishijima, K. Hiraga, M. Yamasaki, Y. Kawamura, Mater. Trans. 47 (2006) 2109. [5] M. Nishijima, K. Hiraga, Mater. Trans. 48 (2007) 10. [6] T. Honma, T. Ohkubo, S. Kamado, K. Hono, Acta Mater. 55 (2007) 4137. [7] I.A. Anyanwu, S. Kamado, Y. Kojima, Mater. Trans. 42 (2001) 1206. [8] S.M. He, X.Q. Zeng, L.M. Peng, X. Gao, J.F. Nie, W.J. Ding, J. Alloys Compd. 427 (2007) 316. [9] K. Yamada, H. Hoshikawa, S. Maki, T. Ozaki, Y. Kuroki, S. Kamado, Y. Kojima, Scripta Mater. 61 (2009) 636. [10] K.Y. Zheng, J. Dong, X.Q. Zeng, W.J. Ding, Mater. Sci. Eng., A 489 (2008) 44. [11] X.Y. Fang, D.Q. Yi, J.F. Nie, Metall. Mater. Trans. A 40A (2009) 2761. [12] M. Yamasaki, T. Anan, S. Yoshimoto, Y. Kawamura, Scripta Mater. 53 (2005) 799. [13] M. Nishijima, K. Hiraga, M. Yamasaki, Y. Kawamura, Mater. Trans. 49 (2008) 227. [14] X. Gao, J.F. Nie, Scripta Mater. 58 (2008) 619. [15] T. Homma, N. Kunito, S. Kamado, Scripta Mater. 61 (2009) 644. [16] N. Stanford, D. Atwell, M.R. Barnett, Acta Mater. 58 (2010) 6773. [17] N. Stanford, M.R. Barnett, Mater. Sci. Eng., A 496 (2008) 399. [18] L. Gao, R.S. Chen, E.H. Han, J. Alloys Compd. 481 (2009) 379. [19] A.A. Nayeb-Hashemi, J.B. Clark, Phase Diagrams of Binary Magnesium Alloys, ASM International, Metals Park, OH, 1988. [20] K. Yamada, Y. Ohkubo, M. Shiono, H. Watanabe, S. Kamado, Y. Kojima, Mater. Trans. 47 (2006) 1066. [21] M. Mabuchi, Y. Chino, H. Iwasaki, T. Aizawa, K. Higashi, Mater. Trans. 42 (2001) 1182. [22] M.R. Barnett, Z. Keshavarz, A.G. Beer, D. Atwell, Acta Mater. 52 (2004) 5093. [23] C.R. Hutchinson, J.F. Nie, S. Gorsse, Metall. Mater. Trans. A 36 (2005) 2093. [24] C.H. Caceres, C.J. Davidson, J.R. Griffiths, C.L. Newton, Mater. Sci. Eng., A 235 (2002) 344. [25] N. Stanford, Mater. Sci. Eng., A 527 (2010) 2669. [26] J. Bohlen, S.B. Yi, D. Letzig, K.U. Kainer, Mater. Sci. Eng., A 527 (2010) 7092. [27] K. Hantzsche, J. Bohlen, J. Wendt, K.U. Kainer, S.B. Yi, D. Letzig, Scripta Mater. 63 (2010) 725. [28] A. Styczynski, Ch. Hartig, J. Bohlen, D. Letzig, Scripta Mater. 50 (2004) 943. [29] M.R. Barnett, M.D. Nave, C.J. Bettles, Mater. Sci. Eng., A 386 (2004) 205.
Scripta Materialia 64 (2011) 950–953 www.elsevier.com/locate/scriptamat
Development of high-strength magnesium alloys via combined processes of extrusion, rolling and ageing R.G. Li,a J.F. Nie,a,b G.J. Huang,a Y.C. Xina and Q. Liua,⇑ a
School of Materials Science and Engineering, Chongqing University, Chongqing 400045, People’s Republic of China b Department of Materials Engineering, Monash University, Victoria 3800, Australia Received 20 November 2010; revised 26 January 2011; accepted 27 January 2011 Available online 17 February 2011
A remarkably high tensile yield strength of 445 MPa is obtained in a Mg–14Gd–0.5Zr (wt.%) alloy when this alloy is prepared by the combined processes of hot extrusion, cold rolling and ageing. This significant improvement in strength is associated with a denser distribution of precipitates, an intensive basal texture and a refined grain size. Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Mg–Gd–Zr alloy; Cold rolling; Precipitate; Ageing
Magnesium alloys containing Gd have received much attention in recent years [1–18]. The equilibrium solid solubility of Gd in Mg is relatively high [19], and the Mg–Gd system is therefore regarded as an ideal one for developing high-strength alloys via precipitation hardening [2]. However, the 0.2% yield strength that has been achieved in peak-aged samples of the Mg–Gd binary alloys only reach 300 MPa [6], which is significantly lower than that of the corresponding Al alloys. Therefore, considerable efforts have been made to further improve the strength of Mg–Gd–based alloys. One approach for improving strength is to add suitable microalloying elements to enhance the precipitation hardening response during the isothermal ageing treatment [14]. Another approach involves the combined use of thermomechanical processing and ageing. It has been reported [7] that a 0.2% proof strength of about 360 MPa can be achieved in an Mg–16.9Gd–0.51Zr (wt.%) alloy when this alloy is hot rolled and aged. This alloy contains an equiaxed grain of above 100 lm, and it is strengthened predominantly by fine precipitates. A tensile yield strength of 311 MPa has also been achieved in an Mg–10Gd–2Y–0.5Zr (wt.%) alloy that was processed by hot extrusion and ageing [8]. In this case, the alloy contains equiaxed grains of 4–7 lm, and the grain boundary strengthening and precipitation hardening contribute to the alloy’s strength.
⇑ Corresponding author. Fax: +86 23 65111295; e-mail: qingliu@cqu. edu.cn
More recently, a value as high as 473 MPa has been reported for the 0.2% proof strength in a Mg–10Gd– 5.7Y–1.62Zn–0.65Zr (wt.%) alloy that was processed by hot extrusion and ageing [15]. The grain size achieved in this alloy is about 1.1 lm, which is considerably smaller than those typically obtained in magnesium extrusion alloys. In addition, the dense distribution of precipitates inside the individual grains and along the grain boundaries also contributes to the ultrahigh strength achieved in this alloy [15]. Cold rolling of magnesium alloys can introduce a high density of dislocations, which may facilitate precipitation and therefore enhance the age hardening response. Therefore, the application of cold rolling after hot extrusion and before artificial ageing is a useful process in the development of high-strength magnesium alloys. Here we report that an ultrahigh 0.2% proof strength of over 440 MPa, together with an ultimate strength of 482 MPa, can be achieved in an Mg– 14Gd–0.5Zr (wt.%) alloy that is processed by the combined use of hot extrusion, cold rolling and ageing. The total concentration of alloying elements in this alloy is much lower than that in the Mg–10Gd–5.7Y–1.62Zn– 0.65Zr (wt.%) alloy [15]. An alloy of nominal composition Mg–14Gd–0.5Zr (wt.%)/Mg–2.67Gd–0.15Zr (at.%) was prepared from high-purity (99.9%) Mg and Gd and an Mg–33Zr (wt.%) master alloy by induction melting in a mild steel crucible at 775 °C and casting into a steel mold under mixed SF6 and CO2. The cast billet was 230 mm in
1359-6462/$ - see front matter Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2011.01.042
R. G. Li et al. / Scripta Materialia 64 (2011) 950–953
951
Table 1. Tensile properties of the extruded, cold-rolled and aged samples.
Figure 1. (a) Ageing curves of the extruded and cold-rolled samples. Engineering stress–strain tensile curves obtained from (b) the extruded and cold-rolled samples and (c) the extruded, cold-rolled and aged samples.
diameter and 630 mm in height. It was machined into cylindrical-shaped samples with a height of 100 mm and a diameter of 80 mm. The billet was homogenized at 505 °C for 24 h and subsequently extruded at 505 °C, with an extrusion ratio of 28:1 and an extrusion rate of 0.6 mm s1, followed by quenching in water of room temperature. The extruded cylindrical bars have a diameter of 16 mm and a length of 250 mm, and they were denoted as E samples. Some extruded samples were also machined into plates 7 mm wide and 2 mm thick, which were subsequently cold, rolled with total thickness reductions of 10%, 19% and 27% and 10% reduction in each pass. They were denoted as E + 10%R, E + 19%R and E + 27%R, respectively. Cracks started to appear at the edges of samples when the amount of cold work was over 27%. Some of the extruded and cold-rolled samples were also aged and letter A was used to represent aged. Vickers hardness testing was carried out using a 1 kg load and a holding time of 15 s at MH-5. The tensile tests were performed at room temperature using a Shimadzu AG-X (10 kN) machine at a strain rate of 0.001 s1. The tensile direction was parallel to the extrusion direction and rolling direction. The microstructures were observed using an FEI NOVA400 scanning electron microscope operating at 20 kV. Electron backscattered diffraction (EBSD) analysis was performed using a step size of 0.5 lm on an FEI Nova 400 equipped with an HKL-EBSD system. Thin foils for transmission electron microscopy (TEM) observations were prepared by low-energy ion beam thinning. The precipitate characterization was carried out using a LIBRA 200 transmission electron microscope operating at 200 kV. Figure 1a shows the hardness curves of the extruded and cold-rolled samples obtained during isothermal ageing at 200 °C. The hardness value is increased by about 15 HV after 10% cold work, but further increasing the amount of cold work from 10% to 27% leads to an increase in hardness of only about 5 HV. The age hardening response of the cold-rolled samples is enhanced compared with the as-extruded sample. The time to reach peak hardness is about 16 h for all cold-rolled samples, while it is about 36 h for the extruded samples. The maximum hardness achievable in the two processing conditions is about 125 HV for the cold-rolled samples and 122 HV for the extruded samples. Figure 1b shows the engineering stress–strain tensile curves obtained from the extruded and cold-rolled samples. The tensile properties obtained from these stress– strain curves are presented in Table 1. Sample E has a 0.2% proof strength of 190 MPa and an elongation to
Alloys
0.2% Proof stress (MPa)
Ultimate tensile strength (MPa)
Elongation (%)
E E + 10%R E + 19%R E + 27%R E+A E + 10%R + A E + 19%R + A E + 27%R + A
190 255 295 305 305 375 420 445
295 300 350 375 446 470 477 482
19.5 12.0 4.5 3.5 7.0 3.0 2.3 2.0
fracture of 19.5%. The cold rolling significantly increases the yield stress at the expenses of the elongation to fracture. The 0.2% proof strength values of 255, 295 and 305 MPa are obtained in the E + 10%R, E + 19%R and E + 27%R conditions, respectively. The tensile properties of the extruded, cold-rolled and aged samples are provided in Figure 1c and Table 1. The ageing treatments greatly enhance the strength. The increments in the 0.2% proof strength are in the range 115–140 MPa. Note that the 0.2% proof strength values of the cold-rolled samples are much higher than that of the samples without any cold rolling. The 27% coldrolled and aged sample exhibits a 0.2% proof strength of 445 MPa, which is substantially higher than those reported previously in the literature [6–10,20], e.g. 331 MPa in an extruded and aged Mg–10Gd–2Y– 0.5Zr alloy [8], 250 MPa in a hot-rolled and aged Mg– 9.3Gd–4.1Y–0.7Zr alloy [7] and 360 MPa in hot-rolled and aged Mg–12Gd–1.9Y–0.7Zr and Mg–16.9Gd– 0.5Zr alloys [7], Table 2. While the strength level obtained in the present work is significantly higher than the alloys of similar compositions, it is slightly lower than that (473 MPa) reported recently in an extruded and aged Mg–10Gd–5.7Y–1.62Zn–0.65Zr alloy [15]. The latter alloy contains a much higher content of alloying elements, and its grain size (about 1 lm) is substantially smaller than that (20 lm) obtained in the present alloy. According to the Hall–Petch relationship, which is valid for basal slip deformation, a further increase in strength by at least 57 MPa is expected if the grain size is refined from 20 to 1 lm, assuming that the k value is about 140 MPa lm1/2 for grain sizes of about 1 lm [21,22] and about 370 MPa lm1/2 for grain sizes of the order of 20 lm [23,24]. Therefore, a further reduction in grain size via the optimization of the extrusion, cold rolling and ageing processes might potentially lead to the development of an alloy with 0.2% proof strength exceeding 500 MPa. Figures 2 and 3 compare the shape, size and orientation of magnesium grains, together with the information on twins and dislocations inside individual magnesium grains, in the E, E + 10%R and E + 27%R samples. The as-extruded microstructure comprises equiaxed grains, implying that dynamic recrystallization has fully occurred during the extrusion process. The average grain size is about 20 lm (Fig. 2a). The colour of the grains in the IPF map (Fig. 3a) indicates that the texture distribution in the as-extruded sample is reasonably homogeneous and nearly random in the centre and edge
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R. G. Li et al. / Scripta Materialia 64 (2011) 950–953
Table 2. Comparative tensile property and process condition in magnesium alloys with similar compositions. Compositions
Process
Grain size (lm)
0.2% proof stress (MPa)
Ultimate tensile strength (MPa)
Elongation (%)
Mg–11.3Gd–3.8Y–0.7Zr [6] Mg–11.2Gd–3.8Y–0.7Zn–0.7Zr [6] Mg–11.1Gd–3.8Y–2.3Zn–0.7Zr [6] Mg–11.1Gd–4.1Y–1.69Zn–0.49Zr[20] Mg–11Gd–2Nd–0.5Zr [10] Mg–18.23Gd–1.88Ag–0.34Zr [9] Mg–10Gd–2Y–0.5Zr [8]
Cast + T6 Cast + T6 Cast + T6 Cast + T6 Cast + T6 Cast + T6 Cast + T6 Extruded + T5 Extruded + T5 Rolled + T5 Rolled + T5 Rolled + T5 Extruded + Rolled + T5
– – – –
300 300 260 231 230 293 239 331 473 360 360 290 445
330 310 310 361 350 414 362 397 542 400 400 310 482
2.0 <2.0 12.0 4.0 2.3 2.2 4.7 12.8 8 4.0 5.0 16.0 2.0
Mg–10Gd–5.66Y–1.62Zn–0.65Zr [15]. Mg–17Gd–0.51Zr [7] Mg–12Gd–1.9Y–0.69Zr [7] Mg–9.3Gd–4.1Y–0.7Zr [7] Mg–14Gd–0.5Zr (this work)
50 – 43 4 1.1 100 100 100 20
regions of the sample. Twins are not detected within the experimental error. Furthermore, there is little colour variation within individual Mg grains, implying the absence of misorientation caused by sub-grain boundaries and excessive density of dislocations. A weaker extrusion texture has also been observed in a Mg– 1.55 wt.%Gd alloy extruded at 450 °C [17] and other
Figure 3. The (0 0 0 2), (11–20) and (10–10) pole figure maps derived from EBSD of the centre position of the E sample at the plane parallel to extrusion direction and of the centre positions of E + 10%R, E + 19% and the E + 27%R samples at the RD–ND plane.
Figure 2. EBSD IPF maps and Kikuchi band contrast maps of the (a) E, (b) E + 10%R and (c) E + 27%R samples. The horizontal direction is the extrusion direction for sample E, and the rolling direction for the E + 10%R and E + 27%R samples, respectively.
extruded Mg–RE based alloys [25–27]. The random texture and the absence of shear bands in the present alloy are attributable to a higher Gd concentration and a higher extrusion temperature. Figure 2b shows the IPF and Kikuchi band contrast maps of the E + 10%R sample. The grain size and grain shape are nearly the same as those of the as-extruded sample. In the IPF map, the grain colour suggests that basal planes of individual Mg grains are close to being parallel to the rolling surface. Many twins, most of which are tensile twins, and double twins are also observed in the Kikuchi band contrast map, represented by the red and blue lines, respectively, in Figure 2b. Figure 3b shows an intensive basal texture in the RD–ND plane and the centre position of the plate. These observations are similar to those made from cold-rolled AZ31 [28], in which an intensive basal texture forms. The formation of a strong basal texture in the cold-rolled samples of the present alloy could be accounted for by the operation of hai-basal, hai-prismatic
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loy, which is remarkably higher than those achieved in alloys of similar compositions that are produced by a single thermomechanical process of extrusion or rolling. This project was financially supported by the National Basic Research Program of China (“973” Project) (Grant No. 2007CB613703) and the Natural Science Foundation of China (Grant No. 50890172). Figure 4. TEM images of b0 precipitates in matrix in the (a) E + A and (b) E + 27%R + A samples. The electron beam is parallel to the [0 0 0 1] direction of the Mg matrix.
and hc + ai-pyramidal slip systems and the tensile twinning system [28] during the rolling process.Figure 2c shows IPF and Kikuchi band contrast of the E + 27%R sample. The grain shape and the number density of twins are similar to those of the E + 10%R sample, and most of {11–20} and {10–10} planes of Mg tend to be parallel to the normal direction. The texture intensity, shown in Figure 3d, is markedly increased compared to the E + 10%R sample. However, the texture intensity is very similar to that of the E + 19%R sample (Fig. 3c). The observation here is consistent with the previous work on cold-rolled pure magnesium, Mg– 0.2 wt.% Ce and AZ31 alloys, where the cold rolling texture does not change much when the rolling reduction is beyond 10% [29]. A higher density of dislocations in the a-Mg matrix is expected in the E + 27%R sample than that in the E + 10%R sample due to more dislocation slip occurring during rolling. Figure 4 shows TEM images of the E + A and E + 27%R + A samples. A denser distribution of b0 precipitates is noted in the E + 27%R + A sample. The b0 precipitates are about 13.8 nm in length and 7.2 nm in width in the E + A sample, and 8.9 nm in length and 6.8 nm in width in the E + 27%R + A sample, when viewed along the [0 0 0 1] direction. If the equilibrium solid solubility of Gd in Mg (0.6 at.% or 4.0 wt.% at 200 °C) is assumed to be reached in the peak-aged condition in each of the two samples, then the precipitate volume fractions are similar in these two samples and the reduction in precipitate size would lead to an approximately twofold increase in precipitate number density. The increased number density of precipitates in the E + 27%R + A sample is due to the higher dislocation density in the matrix, which facilitates the nucleation of precipitates during the ageing process. The increased number density of precipitates increases the 0.2% proof strength from 305 MPa in the E + A sample to 445 MPa in the E + 27%R + A sample. Meanwhile, the ageing time required to reach peak hardness is decreased from 36 h for the extruded sample to 16 h for the cold-rolled samples. In summary, the combined use of extrusion, cold rolling and ageing processes can significantly enhance the strength of the Mg–14Gd–0.5Zr alloy. The cold rolling after extrusion and before ageing can lead to the formation of stronger basal texture and higher density of precipitates in the Mg matrix, which in turn leads to an enhancement of the yield strength of the alloy. A yield stress of 445 MPa is achieved in the Mg–14Gd–0.5Zr al-
[1] T. Honma, T. Ohkubo, K. Hono, S. Kamado, Mater. Sci. Eng., A 395 (2005) 301. [2] J.F. Nie, X. Gao, S.M. Zhu, Scripta Mater. 53 (2005) 1049. [3] X. Gao, S.M. He, X.Q. Zeng, L.M. Peng, W.J. Ding, J.F. Nie, Mater. Sci. Eng., A 431 (2006) 322. [4] M. Nishijima, K. Hiraga, M. Yamasaki, Y. Kawamura, Mater. Trans. 47 (2006) 2109. [5] M. Nishijima, K. Hiraga, Mater. Trans. 48 (2007) 10. [6] T. Honma, T. Ohkubo, S. Kamado, K. Hono, Acta Mater. 55 (2007) 4137. [7] I.A. Anyanwu, S. Kamado, Y. Kojima, Mater. Trans. 42 (2001) 1206. [8] S.M. He, X.Q. Zeng, L.M. Peng, X. Gao, J.F. Nie, W.J. Ding, J. Alloys Compd. 427 (2007) 316. [9] K. Yamada, H. Hoshikawa, S. Maki, T. Ozaki, Y. Kuroki, S. Kamado, Y. Kojima, Scripta Mater. 61 (2009) 636. [10] K.Y. Zheng, J. Dong, X.Q. Zeng, W.J. Ding, Mater. Sci. Eng., A 489 (2008) 44. [11] X.Y. Fang, D.Q. Yi, J.F. Nie, Metall. Mater. Trans. A 40A (2009) 2761. [12] M. Yamasaki, T. Anan, S. Yoshimoto, Y. Kawamura, Scripta Mater. 53 (2005) 799. [13] M. Nishijima, K. Hiraga, M. Yamasaki, Y. Kawamura, Mater. Trans. 49 (2008) 227. [14] X. Gao, J.F. Nie, Scripta Mater. 58 (2008) 619. [15] T. Homma, N. Kunito, S. Kamado, Scripta Mater. 61 (2009) 644. [16] N. Stanford, D. Atwell, M.R. Barnett, Acta Mater. 58 (2010) 6773. [17] N. Stanford, M.R. Barnett, Mater. Sci. Eng., A 496 (2008) 399. [18] L. Gao, R.S. Chen, E.H. Han, J. Alloys Compd. 481 (2009) 379. [19] A.A. Nayeb-Hashemi, J.B. Clark, Phase Diagrams of Binary Magnesium Alloys, ASM International, Metals Park, OH, 1988. [20] K. Yamada, Y. Ohkubo, M. Shiono, H. Watanabe, S. Kamado, Y. Kojima, Mater. Trans. 47 (2006) 1066. [21] M. Mabuchi, Y. Chino, H. Iwasaki, T. Aizawa, K. Higashi, Mater. Trans. 42 (2001) 1182. [22] M.R. Barnett, Z. Keshavarz, A.G. Beer, D. Atwell, Acta Mater. 52 (2004) 5093. [23] C.R. Hutchinson, J.F. Nie, S. Gorsse, Metall. Mater. Trans. A 36 (2005) 2093. [24] C.H. Caceres, C.J. Davidson, J.R. Griffiths, C.L. Newton, Mater. Sci. Eng., A 235 (2002) 344. [25] N. Stanford, Mater. Sci. Eng., A 527 (2010) 2669. [26] J. Bohlen, S.B. Yi, D. Letzig, K.U. Kainer, Mater. Sci. Eng., A 527 (2010) 7092. [27] K. Hantzsche, J. Bohlen, J. Wendt, K.U. Kainer, S.B. Yi, D. Letzig, Scripta Mater. 63 (2010) 725. [28] A. Styczynski, Ch. Hartig, J. Bohlen, D. Letzig, Scripta Mater. 50 (2004) 943. [29] M.R. Barnett, M.D. Nave, C.J. Bettles, Mater. Sci. Eng., A 386 (2004) 205.