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Low-temperature superplasticity of extruded Mg–Sn–Al–Zn alloy
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Scripta Materialia 65 (2011) 202–205 www.elsevier.com/locate/scriptamat
Low-temperature superplasticity of extruded Mg–Sn–Al–Zn alloy S.S. Parka,⇑ and B.S. Youb a
Ulsan National Institute of Science and Technology (UNIST), Ulsan 689-798, Republic of Korea b Korea Institute of Materials Science (KIMS), Changwon 642-831, Republic of Korea Received 1 January 2011; revised 20 February 2011; accepted 4 April 2011 Available online 9 April 2011
The tensile properties of extruded Mg–Sn–Al–Zn alloy at elevated temperature were investigated. Low-temperature superplasticity was found in the alloy, which exhibited tensile elongations of 410–950% at strain rates in the range 1 103–1 104 s1 at 473 K. The superplastic deformation behavior was attributed to the fine-grained microstructure, which contained thermally stable Mg2Sn precipitates. Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Magnesium alloy; Extrusion; Grain size; Precipitate; Superplasticity
Mg–Sn-based alloys have recently attracted strong interest because of their excellent creep resistance at elevated temperature due to the presence of thermally stable Mg2Sn particles in the microstructure [1–5]. More recently, it has been shown that Mg–Sn-based alloys are highly suited to plastic deformation processes such as extrusion [6–10], hence suggesting that they have potential for the development of wrought Mg products as well. Sasaki et al. [6] reported that an extruded Mg–Sn–Zn–Al alloy shows high-strength/low-yield asymmetry at room temperature due to the fine-grained microstructure and the presence of fine Mg2Sn precipitates. For broader application of Mg alloy extrusions, however, it is necessary for them to be readily formable in order to fabricate complicated shapes. Since most Mg alloys do not possess sufficient formability at room temperature, it is necessary to fabricate them at elevated temperatures, using methods such as superplastic forming [11]. In general, superplastic properties are exhibited in materials having stable, equiaxed and fine grains (<10 lm) [12,13]. It is thus believed that extruded Mg– Sn-based alloys should exhibit excellent superplasticity since they exhibit grains <10 lm and as well as thermally stable fine precipitates that can retard grain growth at high temperature. To date, however, there have been few reports on the elevated-temperature deformation behavior of wrought Mg–Sn-based alloys [10]. In the present study, therefore, the tensile properties of extruded Mg–8 wt.% Sn–1 wt.% Al–1 wt.% Zn
⇑ Corresponding author. E-mail: [email protected]
(TAZ811) alloy at elevated temperature were investigated. The analyzed composition of the TAZ811 alloy was Mg–7.69 wt.% Sn–0.94 wt.% Al–0.96 wt.% Zn. Details of the billet casting procedure are described elsewhere [14]. After casting, the alloy was homogenized at 773 K for 3 h and then water-quenched to induce a supersaturated solid solution. Billet dimensions were 80 mm in diameter and 200 mm in length. Indirect extrusion experiments were implemented at an initial billet temperature of 523 K, a ram speed of 0.65 mm s1 and an extrusion ratio of 50. Tensile tests were performed at temperatures of 448 and 473 K at initial strain rates in the range 1 104–1 102 s1. The flow stress was determined at a true strain of 0.1 since microstructural change at small strain was expected to be negligible. Load relaxation tests were performed at 473 K at a strain rate of 1 102 s1. Round tensile specimens with a 20 mm gage length and a 5 mm gage diameter were used for both tensile and load-relaxation tests; specimens were stabilized for 10 min at the selected temperatures prior to testing. The specimens for the load relaxation test were loaded first in tension to a strain of 1.5% before stopping the crosshead motion. The solid solubilities of Sn in the TAZ811 alloy were calculated with Pandat 7.0 software. Figure 1a shows an optical micrograph of the extruded TAZ811 alloy, which reveals fine recrystallized grains. The grain size is as small as 1.5 lm on average. The transmission electron microscopy (TEM) image shown in Figure 1b indicates that the microstructure
1359-6462/$ - see front matter Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2011.04.005
S. S. Park, B. S. You / Scripta Materialia 65 (2011) 202–205
Figure 1. (a) Optical and (b) TEM micrographs of the extruded TAZ811 alloy.
of the extruded TAZ811 alloy can be further characterized by the presence of dynamically precipitated globular Mg2Sn particles with sizes of 50–150 nm at the grain boundaries as well as within the a-Mg matrix, as was similarly observed elsewhere [6–10]. The extruded TAZ811 alloy with fine grains and Mg2Sn precipitates showed excellent tensile properties at room temperature, exhibiting yield and ultimate strengths of 277 and 334 MPa, respectively, with an elongation of 23%. The tensile properties of the extruded TAZ811 alloy at elevated temperatures of 448 and 473 K are shown in Figure 2a and b. As shown in Figure 2a, elongation increases as the strain rate decreases at both testing temperatures. The largest tensile elongations obtained at a strain rate of 1 104 s1 are 540% at 448 K and 950% at 473 K. This clearly shows that low-temperature superplasticity with elongations exceeding 500% is achievable in the TAZ811 alloy, even though this had merely been subjected to an ordinary extrusion process. It is interesting to note that the TAZ811 alloy exhibits a considerable tensile elongation, as large as 410%, at a 10 times faster strain rate of 1 103 s1 at 473 K. When compared to the situation at 473 K and at 1 103 s1, the tensile elongation attained in the extruded TAZ811 alloy turned out to be superior to those of fine-grained Mg alloys with grain sizes of 0.4–1.4 lm, which were produced via severe plastic deformation processes such as equal-channel angular extrusion or different-speed rolling [15–18]. The strain-rate sensitivity exponent m, defined as @logr=@log_e, where r is the flow stress and e_ is the strain rate, was calculated as shown in Figure 2b. It can be seen that m values tend to increase with decreasing strain rate, suggesting that consideration of threshold stress is not necessary at the temperatures and strain rates investigated in this study [15]. An m value of 0.37 was obtained at low strain rates in the range 1 104– 5 104 s1 at 448 K and 1 104–2 103 s1 at
473 K. However, an m value of 0.18 was observed at relatively high strain rates. The m values in the present study are apparently lower than the value of 0.5, which is typically observed in superplastic alloys whose dominant deformation mechanism is grain boundary sliding (GBS) [12]. This suggests that the deformation of the extruded TAZ811 alloy during tensile testing does not occur principally via GBS, although the alloy exhibits superplastic deformation behavior in the relatively low-strain-rate regime. The activation energy for plastic flow was also calculated from the Arrhenius plot of ln e_ vs. 1/T at a constant flow stress of 60 MPa. The evaluated activation energy is 95 kJ mol1, which is very close to the activation energy for grain boundary diffusion of Mg (92 kJ mol–1) but considerably lower than that for lattice self–diffusion of Mg (135 kJ mol–1) [19], suggesting that the deformation of the extruded TAZ811 alloy at 448–473 K is mainly controlled by grain boundary diffusion. Figure 3a and b show the electron back-scatter diffraction (EBSD) orientation maps and pole figures from the grip and gage sections of a tensile sample after a tensile elongation of 900% at 473 K and at a strain rate of 1 104 s1, where m = 0.37. The grip section without tensile deformation reveals a grain size and texture similar to those of the as-extruded alloy, although it was exposed to the elevated temperature for 25 h, demonstrating the unique thermal stability of the TAZ811 alloy. The pole figure from the grip section reveals a type of fiber texture in which basal poles are preferentially perpendicular to the extrusion direction (ED), which is typical of extruded Mg alloys [6,8–10,14]. On the other hand, the gage section subjected to tensile deformation shows larger grains than those of the as-extruded alloy, indicating that dynamic grain growth occurred during tensile deformation, as is usually reported for the superplastically deformed alloys [20]. In addition, grain elongation along the tensile direction and texture randomization are found to occur, suggesting that both grain matrix slip deformation (GMD) and GBS operate together during tensile deformation [21], and are responsible for the large tensile elongation observed under these test conditions. As shown in Figure 3c, however,
103
1200
(a)
(b)
448 K
1000
448 K
473 K
473 K
800
σ (MPa)
Elongation (%)
203
600
102 m = 0.18
400 m = 0.37
200 10
0
10-5
10-4
10-3
(s-1)
10-2
10-1
10-5
10-4
10-3
10-2
10-1
(s-1)
Figure 2. Variations in (a) elongation and (b) flow stress of the extruded TAZ811 alloy as a function of strain rate.
Figure 3. EBSD orientation maps and pole figures of the extruded TAZ811 alloys obtained (a,b) after a tensile elongation of 900% at 473 K and at 1 104 s1 and (c) after a tensile elongation of 190% at 473 K and at 5 103 s1, respectively; (a) grip and (b and c) gage.
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the gage section after a tensile elongation of 190% at 473 K and at a strain rate of 5 103 s1, where m = 0.18, is seen to retain a strong fiber texture, suggesting that GBS does not actively operate during the tensile deformation. To analyze the deformation mechanisms that operate during deformation in more detail, the TAZ811 alloy was subjected to a load relaxation test at 473 K within the framework of an internal variable theory [22]. In this theory, GMD and GBS can be considered to compete against each other at high temperature. Details of the internal variable theory and load relaxation test are described elsewhere [22–24]. Figure 4 shows the flow curve of the extruded TAZ811 alloy obtained from the loadrelaxation test. As plotted, the experimental flow curve was successfully fitted by considering both GMD and GBS, indicating that these two mechanisms operate together during the deformation of the alloy. This result is in good agreement with the EBSD data from the superplastically deformed alloy shown in Figure 3. However, the strain rate by GMD (_eGMD ) was found to be larger than the strain rate by GBS (_eGBS ) at all strain rates, showing that GMD is the dominant deformation mechanism at 473 K regardless of the strain rate. The values of e_ GMD and e_ GBS for each total strain rate (_e ¼ e_ GMD þ e_ GBS ) are listed in Table 1. Since e_ GMD is faster than e_ GBS , the ratios of e_ GBS =_eGMD are lower than unity. However, it should be noted that the ratio increases markedly with decreasing strain rate for testing strain rates of 1 104–1 102 s1, exhibiting a maximum value of 0.54 at the slowest strain rate of 1 104 s1. This means that the activation of GBS is still significant enough to attain low-temperature superplasticity in the extruded TAZ811 alloy even though GBS does not contribute to the deformation as a dominant mechanism. It is noteworthy that the value of e_ GBS =_eGMD is as small as 0.12 at a strain rate of 1 103 s1, where a superplastic elongation of 410% is attainable in the tensile test. The present study shows that fine-grained TAZ811 alloy subjected to extrusion can be superplastic even at low temperatures of 448–473 K. It is generally known that grain refinement is beneficial for enhancing hightemperature plasticity, thereby improving formability, which is due to the activation of GBS [12]. However, 2.5
log σ (MPa)
2
1.5 Experimental
1
GMD GBS GMD + GBS
0.5 -6
-5
-4
-3
log
(s-1)
-2
-1
Figure 4. Experimental flow stress–strain rate curve obtained from the load relaxation test at 473 K and calculated flow stress–strain rate curves for GMD and GBS.
Table 1. Comparison between tensile and load-relaxation tests at 473 K. Tensile test 1
Load-relaxation test
e_ (s )
Elongation (%)
e_ GMD (s1)
e_ GBS (s1)
e_ GBS/_eGMD
1 104 2 104 5 104 1 103 2 103 5 103 1 102
950 900 540 410 300 190 140
6.48 105 1.44 104 4.12 104 8.89 104 1.87 103 4.84 103 9.79 103
3.52 105 5.62 105 8.84 105 1.11 104 1.30 104 1.58 104 2.07 104
0.54 0.39 0.21 0.12 0.07 0.03 0.02
grain coarsening during high-temperature deformation is often problematic, resulting in the deterioration of the enhanced plasticity expected from grain refinement [25]. It has previously been reported that thermally stable dispersoid particles at grain boundaries can retard the growth of fine grains during high-temperature deformation, thereby resulting in superplasticity [26]. Similarly, the presence of Mg2Sn precipitates in the finegrained TAZ811 alloy is considered to contribute decisively to the low-temperature superplasticity. In general, for a second phase to have high thermal stability, alloying elements associated with the second phase should have low diffusivity or low solid solubility in the matrix [27]. Otherwise, the second phase should have low interfacial energy with the matrix [27]. As reported previously [6], globular Mg2Sn particles formed via dynamic precipitation during the extrusion process do not have a specific orientation relationship with the Mg matrix, indicating that they are incoherent with the Mg matrix. In such a case, it is not expected that the Mg2Sn particles will have low interfacial energies. However, the equilibrium solid solubilities of Sn in the TAZ811 alloy are only 0.079 and 0.25 wt.% at 448 and 473 K, respectively, although the maximum solid solubility of Sn in the alloy is as large as 10.1 wt.% at 793 K. Moreover, the diffusivity value of Sn in Mg is known to be much smaller than that of other alloying elements, e.g. Zn [28]. This suggests that the TAZ811 alloy containing Mg2Sn precipitates is better suited to obtaining high thermal stability than conventional Mg–Zn-based alloys with intermetallic phases composed of Mg and Zn. Moreover, the presence of second-phase particles influences the activation of deformation mechanisms. In coarse-grained alloys in which limited GBS is expected, Mg2Sn precipitates finely dispersed within the a-Mg matrix can impede GMD as a dominant deformation mechanism, resulting in the enhancement of creep resistance at elevated temperatures [1–4]. As the grain size becomes smaller, and hence GBS becomes more active, the presence of particles at grain boundaries would become more important in determining the creep resistance. However, the Mg2Sn precipitates in the present study, with sizes of 50–150 nm, are not expected to be effective in preventing GBS, as their size is small enough to avoid the pile-ups of dislocations around the precipitates, as suggested by Dunand and Jansen [29]. Instead, GBS would continue to occur in the presence of the precipitates since grain growth during high-temperature deformation can be
S. S. Park, B. S. You / Scripta Materialia 65 (2011) 202–205
retarded, as mentioned above. Furthermore, the internal strengthening of the a-Mg matrix by the Mg2Sn precipitates is expected to restrict GMD, leading to the facilitation of GBS relative to GMD [26]. This implies that the presence of finely and homogeneously dispersed Mg2Sn precipitates in the fine-grained structure can be quite beneficial to attaining superplasticity by means of effectively activating GBS during high-temperature deformation, enabling the extruded TAZ811 alloy to exhibit low-temperature superplasticity. In summary, the low-temperature superplasticity of the extruded TAZ811 alloy was investigated. The alloy exhibited tensile elongations of 410–950% at strain rates in the range 1 103–1 104 s1 at 473 K. The superplastic deformation behavior can be attributed to the presence of finely dispersed Mg2Sn precipitates in the fine-grained structure, which contributes to activating GBS by means of preventing grain growth as well as restricting GMD. This work was supported by a grant from The World Premier Materials Program funded by The Ministry of Knowledge Economy, Republic of Korea. [1] D.H. Kang, S.S. Park, N.J. Kim, Mater. Sci. Eng. A 413– 414 (2005) 555. [2] D.H. Kang, S.S. Park, Y.S. Oh, N.J. Kim, Mater. Sci. Eng. A 449–451 (2007) 318. [3] S. Wei, Y. Chen, Y. Tang, H. Liu, S. Xiao, G. Niu, X. Zhang, Y. Zhao, Mater. Sci. Eng. A 492 (2008) 20. [4] B.H. Kim, S.W. Lee, Y.H. Park, I.M. Park, J. Alloy. Compd. 493 (2010) 502. [5] M.A. Gibson, X. Fang, C.J. Bettles, C.R. Hutchinson, Scripta Mater. 63 (2010) 899. [6] T.T. Sasaki, K. Yamamoto, T. Honma, S. Kamado, K. Hono, Scripta Mater. 59 (2008) 1111. [7] T.T. Sasaki, J.D. Ju, K. Hono, K.S. Shin, Scripta Mater. 61 (2009) 80.
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[8] S.S. Park, W.N. Tang, B.S. You, Mater. Lett. 64 (2010) 31. [9] W.L. Cheng, S.S. Park, B.S. You, B.H. Koo, Mater. Sci. Eng. A 527 (2010) 4650. [10] S.S. Park, Y.J. Kim, W.L. Cheng, Y.M. Kim, B.S. You, Phil. Mag. Lett. 91 (2011) 37. [11] S.X. Song, J.A. Horton, N.J. Kim, T.G. Nieh, Scripta Mater. 56 (2007) 393. [12] O.D. Sherby, J. Wadsworth, Prog. Mater. Sci. 33 (1989) 169. [13] A.H. Chokshi, A.K. Mukherjee, T.G. Langdon, Mater. Sci. Eng. R10 (1993) 237. [14] S.S. Park, B.S. You, D.J. Yoon, J. Mater. Process. Tech. 209 (2009) 5940. [15] M. Mabuchi, H. Iwasaki, K. Yanase, K. Higashi, Scripta Mater. 36 (1997) 681. [16] M. Mabuchi, K. Ameyama, H. Iwasaki, K. Higashi, Acta Mater. 47 (1999) 2047. [17] H. Watanabe, T. Mukai, K. Ishikawa, K. Higashi, Scripta Mater. 46 (2002) 851. [18] W.J. Kim, J.D. Park, J.Y. Wang, W.S. Yoon, Scripta Mater. 57 (2007) 755. [19] H.J. Frost, M.F. Ashby, Deformation–Mechanism Maps, Pergamon Press, Oxford, 1982. [20] M.A. Clark, T.H. Alden, Acta Metall. 21 (1973) 1195. [21] J.A. del Valle, O.A. Ruano, Acta Mater. 55 (2007) 455. [22] D. Lee, E.W. Hart, Metall. Trans. 2A (1971) 1245. [23] T.K. Ha, Y.W. Chang, Acta Mater. 46 (1998) 2741. [24] S.S. Park, H. Garmestani, G.T. Bae, N.J. Kim, P.E. Krajewski, S. Kim, E.W. Lee, Mater. Sci. Eng. A 435–436 (2006) 687. [25] W.J. Kim, B.H. Lee, J. Alloy. Compd. 482 (2009) 106. [26] S.S. Park, G.T. Bae, D.H. Kang, B.S. You, N.J. Kim, Scripta Mater. 61 (2009) 223. [27] N.J. Kim, Mater. Sci. Eng. A 449–451 (2007) 51. [28] W.F. Gale, T.C. Totemeier (Eds.), Smithells Metals Reference Book, 8th ed., Elsevier, Boston, MA, 2004. [29] D.C. Dunand, A.M. Jansen, Acta Mater. 45 (1997) 4569.
Scripta Materialia 65 (2011) 202–205 www.elsevier.com/locate/scriptamat
Low-temperature superplasticity of extruded Mg–Sn–Al–Zn alloy S.S. Parka,⇑ and B.S. Youb a
Ulsan National Institute of Science and Technology (UNIST), Ulsan 689-798, Republic of Korea b Korea Institute of Materials Science (KIMS), Changwon 642-831, Republic of Korea Received 1 January 2011; revised 20 February 2011; accepted 4 April 2011 Available online 9 April 2011
The tensile properties of extruded Mg–Sn–Al–Zn alloy at elevated temperature were investigated. Low-temperature superplasticity was found in the alloy, which exhibited tensile elongations of 410–950% at strain rates in the range 1 103–1 104 s1 at 473 K. The superplastic deformation behavior was attributed to the fine-grained microstructure, which contained thermally stable Mg2Sn precipitates. Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Magnesium alloy; Extrusion; Grain size; Precipitate; Superplasticity
Mg–Sn-based alloys have recently attracted strong interest because of their excellent creep resistance at elevated temperature due to the presence of thermally stable Mg2Sn particles in the microstructure [1–5]. More recently, it has been shown that Mg–Sn-based alloys are highly suited to plastic deformation processes such as extrusion [6–10], hence suggesting that they have potential for the development of wrought Mg products as well. Sasaki et al. [6] reported that an extruded Mg–Sn–Zn–Al alloy shows high-strength/low-yield asymmetry at room temperature due to the fine-grained microstructure and the presence of fine Mg2Sn precipitates. For broader application of Mg alloy extrusions, however, it is necessary for them to be readily formable in order to fabricate complicated shapes. Since most Mg alloys do not possess sufficient formability at room temperature, it is necessary to fabricate them at elevated temperatures, using methods such as superplastic forming [11]. In general, superplastic properties are exhibited in materials having stable, equiaxed and fine grains (<10 lm) [12,13]. It is thus believed that extruded Mg– Sn-based alloys should exhibit excellent superplasticity since they exhibit grains <10 lm and as well as thermally stable fine precipitates that can retard grain growth at high temperature. To date, however, there have been few reports on the elevated-temperature deformation behavior of wrought Mg–Sn-based alloys [10]. In the present study, therefore, the tensile properties of extruded Mg–8 wt.% Sn–1 wt.% Al–1 wt.% Zn
⇑ Corresponding author. E-mail: [email protected]
(TAZ811) alloy at elevated temperature were investigated. The analyzed composition of the TAZ811 alloy was Mg–7.69 wt.% Sn–0.94 wt.% Al–0.96 wt.% Zn. Details of the billet casting procedure are described elsewhere [14]. After casting, the alloy was homogenized at 773 K for 3 h and then water-quenched to induce a supersaturated solid solution. Billet dimensions were 80 mm in diameter and 200 mm in length. Indirect extrusion experiments were implemented at an initial billet temperature of 523 K, a ram speed of 0.65 mm s1 and an extrusion ratio of 50. Tensile tests were performed at temperatures of 448 and 473 K at initial strain rates in the range 1 104–1 102 s1. The flow stress was determined at a true strain of 0.1 since microstructural change at small strain was expected to be negligible. Load relaxation tests were performed at 473 K at a strain rate of 1 102 s1. Round tensile specimens with a 20 mm gage length and a 5 mm gage diameter were used for both tensile and load-relaxation tests; specimens were stabilized for 10 min at the selected temperatures prior to testing. The specimens for the load relaxation test were loaded first in tension to a strain of 1.5% before stopping the crosshead motion. The solid solubilities of Sn in the TAZ811 alloy were calculated with Pandat 7.0 software. Figure 1a shows an optical micrograph of the extruded TAZ811 alloy, which reveals fine recrystallized grains. The grain size is as small as 1.5 lm on average. The transmission electron microscopy (TEM) image shown in Figure 1b indicates that the microstructure
1359-6462/$ - see front matter Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2011.04.005
S. S. Park, B. S. You / Scripta Materialia 65 (2011) 202–205
Figure 1. (a) Optical and (b) TEM micrographs of the extruded TAZ811 alloy.
of the extruded TAZ811 alloy can be further characterized by the presence of dynamically precipitated globular Mg2Sn particles with sizes of 50–150 nm at the grain boundaries as well as within the a-Mg matrix, as was similarly observed elsewhere [6–10]. The extruded TAZ811 alloy with fine grains and Mg2Sn precipitates showed excellent tensile properties at room temperature, exhibiting yield and ultimate strengths of 277 and 334 MPa, respectively, with an elongation of 23%. The tensile properties of the extruded TAZ811 alloy at elevated temperatures of 448 and 473 K are shown in Figure 2a and b. As shown in Figure 2a, elongation increases as the strain rate decreases at both testing temperatures. The largest tensile elongations obtained at a strain rate of 1 104 s1 are 540% at 448 K and 950% at 473 K. This clearly shows that low-temperature superplasticity with elongations exceeding 500% is achievable in the TAZ811 alloy, even though this had merely been subjected to an ordinary extrusion process. It is interesting to note that the TAZ811 alloy exhibits a considerable tensile elongation, as large as 410%, at a 10 times faster strain rate of 1 103 s1 at 473 K. When compared to the situation at 473 K and at 1 103 s1, the tensile elongation attained in the extruded TAZ811 alloy turned out to be superior to those of fine-grained Mg alloys with grain sizes of 0.4–1.4 lm, which were produced via severe plastic deformation processes such as equal-channel angular extrusion or different-speed rolling [15–18]. The strain-rate sensitivity exponent m, defined as @logr=@log_e, where r is the flow stress and e_ is the strain rate, was calculated as shown in Figure 2b. It can be seen that m values tend to increase with decreasing strain rate, suggesting that consideration of threshold stress is not necessary at the temperatures and strain rates investigated in this study [15]. An m value of 0.37 was obtained at low strain rates in the range 1 104– 5 104 s1 at 448 K and 1 104–2 103 s1 at
473 K. However, an m value of 0.18 was observed at relatively high strain rates. The m values in the present study are apparently lower than the value of 0.5, which is typically observed in superplastic alloys whose dominant deformation mechanism is grain boundary sliding (GBS) [12]. This suggests that the deformation of the extruded TAZ811 alloy during tensile testing does not occur principally via GBS, although the alloy exhibits superplastic deformation behavior in the relatively low-strain-rate regime. The activation energy for plastic flow was also calculated from the Arrhenius plot of ln e_ vs. 1/T at a constant flow stress of 60 MPa. The evaluated activation energy is 95 kJ mol1, which is very close to the activation energy for grain boundary diffusion of Mg (92 kJ mol–1) but considerably lower than that for lattice self–diffusion of Mg (135 kJ mol–1) [19], suggesting that the deformation of the extruded TAZ811 alloy at 448–473 K is mainly controlled by grain boundary diffusion. Figure 3a and b show the electron back-scatter diffraction (EBSD) orientation maps and pole figures from the grip and gage sections of a tensile sample after a tensile elongation of 900% at 473 K and at a strain rate of 1 104 s1, where m = 0.37. The grip section without tensile deformation reveals a grain size and texture similar to those of the as-extruded alloy, although it was exposed to the elevated temperature for 25 h, demonstrating the unique thermal stability of the TAZ811 alloy. The pole figure from the grip section reveals a type of fiber texture in which basal poles are preferentially perpendicular to the extrusion direction (ED), which is typical of extruded Mg alloys [6,8–10,14]. On the other hand, the gage section subjected to tensile deformation shows larger grains than those of the as-extruded alloy, indicating that dynamic grain growth occurred during tensile deformation, as is usually reported for the superplastically deformed alloys [20]. In addition, grain elongation along the tensile direction and texture randomization are found to occur, suggesting that both grain matrix slip deformation (GMD) and GBS operate together during tensile deformation [21], and are responsible for the large tensile elongation observed under these test conditions. As shown in Figure 3c, however,
103
1200
(a)
(b)
448 K
1000
448 K
473 K
473 K
800
σ (MPa)
Elongation (%)
203
600
102 m = 0.18
400 m = 0.37
200 10
0
10-5
10-4
10-3
(s-1)
10-2
10-1
10-5
10-4
10-3
10-2
10-1
(s-1)
Figure 2. Variations in (a) elongation and (b) flow stress of the extruded TAZ811 alloy as a function of strain rate.
Figure 3. EBSD orientation maps and pole figures of the extruded TAZ811 alloys obtained (a,b) after a tensile elongation of 900% at 473 K and at 1 104 s1 and (c) after a tensile elongation of 190% at 473 K and at 5 103 s1, respectively; (a) grip and (b and c) gage.
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S. S. Park, B. S. You / Scripta Materialia 65 (2011) 202–205
the gage section after a tensile elongation of 190% at 473 K and at a strain rate of 5 103 s1, where m = 0.18, is seen to retain a strong fiber texture, suggesting that GBS does not actively operate during the tensile deformation. To analyze the deformation mechanisms that operate during deformation in more detail, the TAZ811 alloy was subjected to a load relaxation test at 473 K within the framework of an internal variable theory [22]. In this theory, GMD and GBS can be considered to compete against each other at high temperature. Details of the internal variable theory and load relaxation test are described elsewhere [22–24]. Figure 4 shows the flow curve of the extruded TAZ811 alloy obtained from the loadrelaxation test. As plotted, the experimental flow curve was successfully fitted by considering both GMD and GBS, indicating that these two mechanisms operate together during the deformation of the alloy. This result is in good agreement with the EBSD data from the superplastically deformed alloy shown in Figure 3. However, the strain rate by GMD (_eGMD ) was found to be larger than the strain rate by GBS (_eGBS ) at all strain rates, showing that GMD is the dominant deformation mechanism at 473 K regardless of the strain rate. The values of e_ GMD and e_ GBS for each total strain rate (_e ¼ e_ GMD þ e_ GBS ) are listed in Table 1. Since e_ GMD is faster than e_ GBS , the ratios of e_ GBS =_eGMD are lower than unity. However, it should be noted that the ratio increases markedly with decreasing strain rate for testing strain rates of 1 104–1 102 s1, exhibiting a maximum value of 0.54 at the slowest strain rate of 1 104 s1. This means that the activation of GBS is still significant enough to attain low-temperature superplasticity in the extruded TAZ811 alloy even though GBS does not contribute to the deformation as a dominant mechanism. It is noteworthy that the value of e_ GBS =_eGMD is as small as 0.12 at a strain rate of 1 103 s1, where a superplastic elongation of 410% is attainable in the tensile test. The present study shows that fine-grained TAZ811 alloy subjected to extrusion can be superplastic even at low temperatures of 448–473 K. It is generally known that grain refinement is beneficial for enhancing hightemperature plasticity, thereby improving formability, which is due to the activation of GBS [12]. However, 2.5
log σ (MPa)
2
1.5 Experimental
1
GMD GBS GMD + GBS
0.5 -6
-5
-4
-3
log
(s-1)
-2
-1
Figure 4. Experimental flow stress–strain rate curve obtained from the load relaxation test at 473 K and calculated flow stress–strain rate curves for GMD and GBS.
Table 1. Comparison between tensile and load-relaxation tests at 473 K. Tensile test 1
Load-relaxation test
e_ (s )
Elongation (%)
e_ GMD (s1)
e_ GBS (s1)
e_ GBS/_eGMD
1 104 2 104 5 104 1 103 2 103 5 103 1 102
950 900 540 410 300 190 140
6.48 105 1.44 104 4.12 104 8.89 104 1.87 103 4.84 103 9.79 103
3.52 105 5.62 105 8.84 105 1.11 104 1.30 104 1.58 104 2.07 104
0.54 0.39 0.21 0.12 0.07 0.03 0.02
grain coarsening during high-temperature deformation is often problematic, resulting in the deterioration of the enhanced plasticity expected from grain refinement [25]. It has previously been reported that thermally stable dispersoid particles at grain boundaries can retard the growth of fine grains during high-temperature deformation, thereby resulting in superplasticity [26]. Similarly, the presence of Mg2Sn precipitates in the finegrained TAZ811 alloy is considered to contribute decisively to the low-temperature superplasticity. In general, for a second phase to have high thermal stability, alloying elements associated with the second phase should have low diffusivity or low solid solubility in the matrix [27]. Otherwise, the second phase should have low interfacial energy with the matrix [27]. As reported previously [6], globular Mg2Sn particles formed via dynamic precipitation during the extrusion process do not have a specific orientation relationship with the Mg matrix, indicating that they are incoherent with the Mg matrix. In such a case, it is not expected that the Mg2Sn particles will have low interfacial energies. However, the equilibrium solid solubilities of Sn in the TAZ811 alloy are only 0.079 and 0.25 wt.% at 448 and 473 K, respectively, although the maximum solid solubility of Sn in the alloy is as large as 10.1 wt.% at 793 K. Moreover, the diffusivity value of Sn in Mg is known to be much smaller than that of other alloying elements, e.g. Zn [28]. This suggests that the TAZ811 alloy containing Mg2Sn precipitates is better suited to obtaining high thermal stability than conventional Mg–Zn-based alloys with intermetallic phases composed of Mg and Zn. Moreover, the presence of second-phase particles influences the activation of deformation mechanisms. In coarse-grained alloys in which limited GBS is expected, Mg2Sn precipitates finely dispersed within the a-Mg matrix can impede GMD as a dominant deformation mechanism, resulting in the enhancement of creep resistance at elevated temperatures [1–4]. As the grain size becomes smaller, and hence GBS becomes more active, the presence of particles at grain boundaries would become more important in determining the creep resistance. However, the Mg2Sn precipitates in the present study, with sizes of 50–150 nm, are not expected to be effective in preventing GBS, as their size is small enough to avoid the pile-ups of dislocations around the precipitates, as suggested by Dunand and Jansen [29]. Instead, GBS would continue to occur in the presence of the precipitates since grain growth during high-temperature deformation can be
S. S. Park, B. S. You / Scripta Materialia 65 (2011) 202–205
retarded, as mentioned above. Furthermore, the internal strengthening of the a-Mg matrix by the Mg2Sn precipitates is expected to restrict GMD, leading to the facilitation of GBS relative to GMD [26]. This implies that the presence of finely and homogeneously dispersed Mg2Sn precipitates in the fine-grained structure can be quite beneficial to attaining superplasticity by means of effectively activating GBS during high-temperature deformation, enabling the extruded TAZ811 alloy to exhibit low-temperature superplasticity. In summary, the low-temperature superplasticity of the extruded TAZ811 alloy was investigated. The alloy exhibited tensile elongations of 410–950% at strain rates in the range 1 103–1 104 s1 at 473 K. The superplastic deformation behavior can be attributed to the presence of finely dispersed Mg2Sn precipitates in the fine-grained structure, which contributes to activating GBS by means of preventing grain growth as well as restricting GMD. This work was supported by a grant from The World Premier Materials Program funded by The Ministry of Knowledge Economy, Republic of Korea. [1] D.H. Kang, S.S. Park, N.J. Kim, Mater. Sci. Eng. A 413– 414 (2005) 555. [2] D.H. Kang, S.S. Park, Y.S. Oh, N.J. Kim, Mater. Sci. Eng. A 449–451 (2007) 318. [3] S. Wei, Y. Chen, Y. Tang, H. Liu, S. Xiao, G. Niu, X. Zhang, Y. Zhao, Mater. Sci. Eng. A 492 (2008) 20. [4] B.H. Kim, S.W. Lee, Y.H. Park, I.M. Park, J. Alloy. Compd. 493 (2010) 502. [5] M.A. Gibson, X. Fang, C.J. Bettles, C.R. Hutchinson, Scripta Mater. 63 (2010) 899. [6] T.T. Sasaki, K. Yamamoto, T. Honma, S. Kamado, K. Hono, Scripta Mater. 59 (2008) 1111. [7] T.T. Sasaki, J.D. Ju, K. Hono, K.S. Shin, Scripta Mater. 61 (2009) 80.
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