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Hydrogen-induced delayed fracture of a Fe–22Mn–0.6C steel pre-strained at different strain rates
Available online at www.sciencedirect.com
Scripta Materialia 66 (2012) 947–950 www.elsevier.com/locate/scriptamat
Hydrogen-induced delayed fracture of a Fe–22Mn–0.6C steel pre-strained at different strain rates Motomichi Koyama,a,b,⇑ Eiji Akiyamab and Kaneaki Tsuzakia,b a
Doctoral Program in Materials Science and Engineering, University of Tsukuba, 1-2-1 Sengen, Tsukuba, Ibaraki 305-0047, Japan b National Institute for Materials Science, 1-2-1 Sengen, Tsukuba, Ibaraki 305-0047, Japan Received 22 January 2012; revised 24 February 2012; accepted 24 February 2012 Available online 3 March 2012
Hydrogen-induced delayed fracture under loading was investigated in a Fe–22Mn–0.6C twinning-induced plasticity steel that had been pre-deformed at various strain rates. Hydrogen-induced delayed fracture was suppressed by increasing the strain rate of the pre-deformation. In this study on the strain-rate effect, factors affecting the delayed fracture were found to be the negative strain-rate sensitivity of flow stress, stress drop caused by the relaxation phenomenon, and the increase in material strength due to strain aging. Ó 2012 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Austenitic steel; Tension test; Hydrogen embrittlement; Twinning; Strain aging
Twinning-induced plasticity (TWIP) steels belong to a new group of high-strength steels; these materials show an exceptional combination of elongation and tensile strength [1–4]. TWIP steels have recently begun to be used in automobile production [5,6]. The optimal chemical composition has been found to be Fe–22Mn– 0.6C (wt.%) [3,7,8] to obtain superior elongation and tensile strength. However, hydrogen-induced delayed fracture (HDF) has been reported in TWIP steels [2,5,9–12], and this poses a critical problem to the use of TWIP steels as materials for automobile production. To solve the problem, the mechanism of HDF in TWIP steels needs to be clarified. HDF of TWIP steels has been investigated by many researchers [9–15]. In general, factors affecting HDF are diffusible hydrogen content, residual stress and material strength (or microstructure). In a previous study, we evaluated the first two factors by tensile tests [16], and reported that, as in ferritic high-strength steels, the effect of diffusible hydrogen content and fracture stress on HDF could be described by a power law in a Fe–18Mn–0.6C steel [16]. The quantitative relationship between diffusible hydrogen content and fracture stress would change depending on the type of material, which
⇑ Corresponding
author at: Doctoral Program in Materials Science and Engineering, University of Tsukuba, 1-2-1 Sengen, Tsukuba, Ibaraki 305-0047, Japan. Tel.: +81 29 859 2000; fax: +81 29 859 2101; e-mail: [email protected]
is the third factor. In TWIP steels, it has been claimed that HDF is influenced by reduction in cohesive energy at grain boundaries [16], martensite formation [11], twin–twin interaction [10], twin–grain boundary interaction [10] and residual stress [14]. As a characteristic phenomenon in Fe–Mn–C TWIP steels, static [17,18] and dynamic [19–21] strain aging is also known to affect material strength and dislocation microstructure due to an increase in dislocation density [22] and the formation of dislocation–carbon pairs [19,20,23,24]. Additionally, a typical effect of dynamic strain aging, which is called negative strain-rate sensitivity, is a decrease in flow stress with increasing strain rate [24–27]. However, the effect of strain aging on HDF has never been discussed in TWIP steels. In this paper, we report and discuss the effect of the strain aging on HDF at various strain rates. A TWIP steel with chemical composition Fe– 22.1Mn–0.61C (wt.%) was prepared by vacuum induction melting; this is the same steel as was prepared for the previous work [18]. The thickness of the steel was reduced from 60 to 2.6 mm by hot rolling at 1273 K. Subsequently, the thickness was reduced to 1.4 mm by cold rolling. It was then solution treated at 1073 K for 1 h under an argon atmosphere. All the specimens were cut with spark erosion. The microstructure had an average grain size of 3 lm, including some annealing twin boundaries [18]. The specimen thickness was further reduced by mechanical grinding to 0.3 mm in order to remove the layer affected by the solution treatment.
1359-6462/$ - see front matter Ó 2012 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2012.02.040
948
M. Koyama et al. / Scripta Materialia 66 (2012) 947–950
Tensile tests were conducted at ambient temperature at various strain rates. The gauge dimensions of the tensile specimen in the present study was 4.0 mm wide 0.3 mm thick 10 mm long with a grip section on both ends that could be fixed in an Instron-type machine. The strains were determined by dividing the displacements by the initial gauge length. Hydrogen was introduced into the specimens by electrochemical charging with a current density of 7 A m–2 in a 3% NaCl aqueous solution containing 3 g l–1 of NH4SCN at ambient temperature under tensile loading at a constant strain and with a fixed cross-head position. A platinum wire was used as the counter-electrode. The hydrogen content was measured by thermal desorption analysis (TDA) from room temperature to 550 K. The heating rate was 200 K h–1. Diffusible hydrogen content was determined by measuring the cumulative desorbed hydrogen from ambient temperature to 473 K. The diffusible hydrogen is defined as hydrogen that diffuses at room temperature. The diffusible hydrogen is reported to play a key role in HDF [28]. First, we show the strain-rate dependence of the stress–strain response and the strain aging effect. Figure 1a shows the results of relaxation tests at various initial strain rates. The specimens were deformed until 49% strain, and then held for 1000 s under loading at the fixed cross-head position. The specimens were deformed
again at the same strain rate as the initial strain rate. The serrations are known to stem from dynamic strain aging [19,20,24]. The flow stress decreased with increasing strain rate. As mentioned above, the decrease in flow stress is related to dynamic strain aging. In other words, suppression of dynamic strain aging by increasing strain rate decreased the flow stresses. Figure 1b shows the engineering stress–strain curves corresponding to the part outlined by broken lines in Figure 1a. The stress drop and the stress increase after holding arise from the relaxation phenomenon [17] and static strain aging [17,18], respectively. The magnitudes of stress drop and increase were defined as rr = r1-r0 and rs = r2r0, where r0, r1 and r2 are indicated in Figure 1b. As shown in Figure 1c, the stress drop arising from the relaxation increased with increasing strain rate from 1.7 10 5 to 1.7 10 4 s 1, and became constant in the strain rate range from 1.7 10 4 to 1.7 10 1 s 1. The stress increase arising from static strain aging linearly increased with increasing logarithmic strain rate. Figure 2a shows engineering stress–strain curves at initial strain rates of 1.7 10 5, 1.7 10 4, 1.7 10 3 and 1.7 10 2 s 1. The tensile deformations were stopped once at 69% strain, and then hydrogen was introduced under loading. If fracture caused by the hydrogen charging did not occur within 10 h, the specimen was deformed again until fracture at the same
Figure 1. (a) Engineering stress–strain curves with relaxation tests. (b) The curves outlined by broken lines in (a). (c) The relationships between strain rate and the stress drop due to relaxation, and between strain rate and the stress increase due to static strain aging.
Figure 2. (a) Tensile tests under hydrogen charging. The hydrogen was introduced at 69% tensile strain. (b) Engineering stresses plotted against hydrogen-charging time. HD indicates diffusible hydrogen content.
M. Koyama et al. / Scripta Materialia 66 (2012) 947–950
of 982 MPa. In Figure 4, the brittle fracture stress tested at the initial strain rate of 1.7 10 5 s 1 and the maximum stress tested at the initial strain rate of 1.7 10 3 s 1can be converted to true stresses of 1472 and 1828 MPa, respectively. Despite the fact that the brittle fracture stress was obviously lower than that of the ductile fracture stress, and diffusible hydrogen contents of them were comparable, HDF occurred at an initial strain rate of 1.7 10 5 s 1. This trend cannot be explained by the reduction in stress associated with the negative strain-rate sensitivity and the relaxation. In 1200
Strain rate: 1.7×10-5 s-1 Hydrogen was introduced at 49% HD: 6.7 wt.ppm
1000
Engineering stress (MPa)
strain rate with the initial strain rate. The holding time is limited to 10 h by the ability of the tensile test machine. Figure 2b shows stress–time curves corresponding to Figure 2a. HDF was observed at strain rates of 1.7 10 5 and 1.7 10 4 s 1. The fracture mode was intergranular as shown in Figure 3a and b. In contrast, HDF was not observed at strain rates of 1.7 10 3 and 1.7 10 2 s 1. The fracture mode observed after the tensile tests subsequent to constant strain holding was totally ductile as shown in Figure 3c and d. From the viewpoint of time to fracture, the HDF was suppressed by increasing strain rate as shown in Figure 2b. This is because the engineering flow stress at 69% strain decreased from 1053 to 999MPawith increasing strain rate from 1.7 10 5 to 1.7 10 2 s 1as shown in Figure 2a. Additionally, the stress drop associated with the relaxation [17] was observed, which was enhanced by increasing the strain rate from 1.7 10 5 to 1.7 10 4 s 1as shown in Figure 1c. These factors decreased the stress at the constant strain with increasing strain rate, suppressing the HDF. Another important trend of HDF is shown in Figure 4. The curve obtained at 1.7 10 3 s 1 is the same as that in Figure 2a. The other engineering stress–strain curve was obtained at an initial strain rate of 1.7 10 5 s 1. The tensile deformation was stopped at 49% strain, and the strain was subsequently held for 10 h under loading at constant strain with hydrogen charging. Then, the specimen was deformed again until fracture. The diffusible hydrogen contents of the specimen, which were measured after fracture, were 6.9 and 6.7 ppm, respectively. The test at the initial strain rate of 1.7 10 5 s 1 showed HDF at the engineering stress
949
800
Hydrogen charging 600
Strain rate: 1.7×10-3 s-1 Hydrogen was introduced at 69% HD: 6.9 wt.ppm
400
200
0 0
10
20 30 40 50 60 70 Engineering plastic strain (%)
80
90
Figure 4. The result of a tensile test including relaxation for 10 h at 49% strain at an initial strain rate of 1.7 10 5 s 1. The tensile curve obtained at the strain rate of 1.7 10 3 s 1 is the same as that of Figure 2a.
Figure 3. Fractographs showing the intergranular fracture surface at (a) low and (b) high magnifications that were provided by the deformation at 1.7 10 5 s 1 in Figure 2a, and showing the ductile fracture surface at (c) low and (d) high magnifications that were provided by the deformation at 1.7 10 2 s 1 in Figure 2a.
950
M. Koyama et al. / Scripta Materialia 66 (2012) 947–950
the previous studies [17,18] and Figure 1, static strain aging under loading was reported to significantly increase the strength of the material due to formation of carbon–dislocation pairs. The susceptibility to HDF is generally reported to become higher with increasing material strength level in the same alloy group [29–31]. The hardening mechanism of dynamic strain aging is similar to that of static strain aging, indicating that static strain aging as well as the promotion of dynamic strain aging by decreasing strain rate must be considered to deteriorate the resistance to HDF. The occurrence of HDF at relatively low stress in Figure 4 is related to the hardening caused by dynamic strain aging. If the specimen deformed at 1.7 10 3 s 1was exposed under loading for a long time, i.e. more than 10 h, it also would show HDF due to strengthening by static strain aging. The effect of the increase in material strength due to static and dynamic strain aging would be the third factor affecting the present HDF property. In conclusion, due to the influence of strain aging under loading, there are three factors affecting HDF: the negative strain-rate sensitivity of flow stress, the drop in stress caused by the relaxation phenomenon, and the increase in material strength due to strain aging. All three factors were strongly dependent on strain rate; the negative strain-rate sensitivity [2,24] and static strain aging [17,18] are widely accepted to arise from carbon diffusion. Hence, we note the following two points in Fe–Mn–C TWIP steels: 1. HDF is sensitive to the strain rate of predeformation. 2. It is important to consider the diffusion of both hydrogen and carbon diffusion when discussing HDF. M.K. acknowledges a Research Fellowship of NIMS Junior Researcher (2009-2010) and the Japan Society for the Promotion of Science for Young Scientists (2011). We would also like to acknowledge Dr. Takahiro Sawaguchi for taking part in the discussions during the experiments. POSCO supported this work by providing the samples and funding. [1] O. Gra¨ssel, G. Frommeyer, Mater. Sci. Technol. 14 (1998) 1213. [2] B.C. De Cooman, K.-G. Chin, J. Kim, High MnTWIP steels for automotive applications, in: M. Chiaberge (Ed.), New Trends and Developments in Automotive System Engineering, InTech., 2011, ISBN: 978-953-307-517-4. [3] O. Bouaziz, S. Allain, C.P. Scott, P. Cugy, D. Barbier, Curr. Opin. Solid State Mater. Sci. 15 (2011) 141.
[4] M. Koyama, T. Sawaguchi, K. Tsuzaki, Tetsu-to-Hagane´, 98, in press. [5] O. Kwon, Proc. HMnS (2011) CD-ROM. [6] C. Scott, S. Allain, M. Faral, N. Guelton, Rev. Metal. Cah.D’inf. Tech. 103 (2006) p. 293. [7] I. Gutierrez-Urrutia, S. Zaefferer, D. Raabe, Mater. Sci. Eng. A 527 (2010) p. 3552. [8] I. Gutierrez-Urrutia, D. Raabe, Acta Mater. 59 (2011) 6449. [9] M. Koyama, E. Akiyama, K. Tsuzaki, Corr. Sci. 54 (2012) 1. [10] M. Koyama, E. Akiyama, T. Sawaguchi, D. Raabe, K. Tsuzaki, Scripta Mater., in press. [11] Y.S. Chun, J.S. Kim, K.-T. Park, Y.-K. Lee, C.S. Lee, Mater. Sci. Eng. A 533 (2012) 87. [12] J.A. Ronevich, S.K. Kim, J.G. Speer, D.K. Matlock, Scripta Mater., in press, doi:10.1016/j.scriptamat.2011. 12.012. [13] K.H. So, J.S. Kim, Y.S. Chun, K.-T. Park, Y.-K. Lee, C.S. Lee, ISIJ Int. 49 (2009) 1952. [14] K-G. Chin, C-Y. Kang, S.Y. Sin, S. Hong, S. Lee, H-S. Kim, K-H. Kim, N.J. Kim, Mater. Sci. Eng. A 528 (2011) p. 2922. [15] R.T. van Tol, L. Zhao, J. Sietsma, Proc. HMnS (2011) CD-ROM. [16] M. Koyama, E. Akiyama, K. Tsuzaki, Corr. Sci., in press. [17] S. Allain, O. Bouaziz, T. Lebedkina, M. Lebyodkin, Scripta Mater. 64 (2011) 741. [18] M. Koyama, E. Akiyama, K. Tsuzaki, Mater. Charact., submitted for publication. [19] L. Chen, H.-S. Kim, S.-K. Kim, B.C. De Cooman, ISIJ Int. 47 (2007) 1804. [20] S.-J. Lee, J. Kim, S.N. Kane, B.C. De Cooman, Acta Mater. 59 (2011) 6809. [21] M. Koyama, T. Sawaguchi, T. Lee, C.S. Lee, K. Tsuzaki, Mater. Sci. Eng. A 528 (2011) 7310. [22] B. Hutchinson, N. Ridley, Scripta Mater. 55 (2006) 299. [23] M. Koyama, T. Sawaguchi, K. Tsuzaki, Mater. Trans. 53 (2012) 546. [24] Y.N. Dastur, W.C. Leslie, Metall. Trans. A 12A (1981) 749. [25] C. Canadinc, C. Efstathiou, H. Sehitoglu, Scripta Mater. 59 (2008) 1103. [26] L. Durrenberger, A. Molinari, A. Rusinek, Mater. Sci. Eng. A 478 (2008) p. 297. [27] S. Allain, P. Cugy, C. Scott, J.-P. Chateau, A. Rusinek, A. Deschamps, Int. J. Mater. Res. 99 (2008) 7. [28] K. Takai, R. Watanuki, ISIJ Int. 43 (2003) 520. [29] M. Wang, E. Akiyama, K. Tsuzaki, Scripta Mater. 52 (2005) 403. [30] E. Akiyama, S. Li, Z. Zhang, M. Wang, K. Matsukado, K. Tsuzaki, Proc. Int. Hydrol. Conf. (2008) 54. [31] T. Fujita, Y. Yamada, Physical metallurgy and SCC in high strength steels, in: R.W. Staehie, J. Hochmann, R.D. McCright, J.E. Slater (Eds.), Stress Corrosion Cracking and Hydrogen Embrittlement of Iron Base Alloys, NACE-5, NACE, 1977, pp. 736–746.
Scripta Materialia 66 (2012) 947–950 www.elsevier.com/locate/scriptamat
Hydrogen-induced delayed fracture of a Fe–22Mn–0.6C steel pre-strained at different strain rates Motomichi Koyama,a,b,⇑ Eiji Akiyamab and Kaneaki Tsuzakia,b a
Doctoral Program in Materials Science and Engineering, University of Tsukuba, 1-2-1 Sengen, Tsukuba, Ibaraki 305-0047, Japan b National Institute for Materials Science, 1-2-1 Sengen, Tsukuba, Ibaraki 305-0047, Japan Received 22 January 2012; revised 24 February 2012; accepted 24 February 2012 Available online 3 March 2012
Hydrogen-induced delayed fracture under loading was investigated in a Fe–22Mn–0.6C twinning-induced plasticity steel that had been pre-deformed at various strain rates. Hydrogen-induced delayed fracture was suppressed by increasing the strain rate of the pre-deformation. In this study on the strain-rate effect, factors affecting the delayed fracture were found to be the negative strain-rate sensitivity of flow stress, stress drop caused by the relaxation phenomenon, and the increase in material strength due to strain aging. Ó 2012 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Austenitic steel; Tension test; Hydrogen embrittlement; Twinning; Strain aging
Twinning-induced plasticity (TWIP) steels belong to a new group of high-strength steels; these materials show an exceptional combination of elongation and tensile strength [1–4]. TWIP steels have recently begun to be used in automobile production [5,6]. The optimal chemical composition has been found to be Fe–22Mn– 0.6C (wt.%) [3,7,8] to obtain superior elongation and tensile strength. However, hydrogen-induced delayed fracture (HDF) has been reported in TWIP steels [2,5,9–12], and this poses a critical problem to the use of TWIP steels as materials for automobile production. To solve the problem, the mechanism of HDF in TWIP steels needs to be clarified. HDF of TWIP steels has been investigated by many researchers [9–15]. In general, factors affecting HDF are diffusible hydrogen content, residual stress and material strength (or microstructure). In a previous study, we evaluated the first two factors by tensile tests [16], and reported that, as in ferritic high-strength steels, the effect of diffusible hydrogen content and fracture stress on HDF could be described by a power law in a Fe–18Mn–0.6C steel [16]. The quantitative relationship between diffusible hydrogen content and fracture stress would change depending on the type of material, which
⇑ Corresponding
author at: Doctoral Program in Materials Science and Engineering, University of Tsukuba, 1-2-1 Sengen, Tsukuba, Ibaraki 305-0047, Japan. Tel.: +81 29 859 2000; fax: +81 29 859 2101; e-mail: [email protected]
is the third factor. In TWIP steels, it has been claimed that HDF is influenced by reduction in cohesive energy at grain boundaries [16], martensite formation [11], twin–twin interaction [10], twin–grain boundary interaction [10] and residual stress [14]. As a characteristic phenomenon in Fe–Mn–C TWIP steels, static [17,18] and dynamic [19–21] strain aging is also known to affect material strength and dislocation microstructure due to an increase in dislocation density [22] and the formation of dislocation–carbon pairs [19,20,23,24]. Additionally, a typical effect of dynamic strain aging, which is called negative strain-rate sensitivity, is a decrease in flow stress with increasing strain rate [24–27]. However, the effect of strain aging on HDF has never been discussed in TWIP steels. In this paper, we report and discuss the effect of the strain aging on HDF at various strain rates. A TWIP steel with chemical composition Fe– 22.1Mn–0.61C (wt.%) was prepared by vacuum induction melting; this is the same steel as was prepared for the previous work [18]. The thickness of the steel was reduced from 60 to 2.6 mm by hot rolling at 1273 K. Subsequently, the thickness was reduced to 1.4 mm by cold rolling. It was then solution treated at 1073 K for 1 h under an argon atmosphere. All the specimens were cut with spark erosion. The microstructure had an average grain size of 3 lm, including some annealing twin boundaries [18]. The specimen thickness was further reduced by mechanical grinding to 0.3 mm in order to remove the layer affected by the solution treatment.
1359-6462/$ - see front matter Ó 2012 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2012.02.040
948
M. Koyama et al. / Scripta Materialia 66 (2012) 947–950
Tensile tests were conducted at ambient temperature at various strain rates. The gauge dimensions of the tensile specimen in the present study was 4.0 mm wide 0.3 mm thick 10 mm long with a grip section on both ends that could be fixed in an Instron-type machine. The strains were determined by dividing the displacements by the initial gauge length. Hydrogen was introduced into the specimens by electrochemical charging with a current density of 7 A m–2 in a 3% NaCl aqueous solution containing 3 g l–1 of NH4SCN at ambient temperature under tensile loading at a constant strain and with a fixed cross-head position. A platinum wire was used as the counter-electrode. The hydrogen content was measured by thermal desorption analysis (TDA) from room temperature to 550 K. The heating rate was 200 K h–1. Diffusible hydrogen content was determined by measuring the cumulative desorbed hydrogen from ambient temperature to 473 K. The diffusible hydrogen is defined as hydrogen that diffuses at room temperature. The diffusible hydrogen is reported to play a key role in HDF [28]. First, we show the strain-rate dependence of the stress–strain response and the strain aging effect. Figure 1a shows the results of relaxation tests at various initial strain rates. The specimens were deformed until 49% strain, and then held for 1000 s under loading at the fixed cross-head position. The specimens were deformed
again at the same strain rate as the initial strain rate. The serrations are known to stem from dynamic strain aging [19,20,24]. The flow stress decreased with increasing strain rate. As mentioned above, the decrease in flow stress is related to dynamic strain aging. In other words, suppression of dynamic strain aging by increasing strain rate decreased the flow stresses. Figure 1b shows the engineering stress–strain curves corresponding to the part outlined by broken lines in Figure 1a. The stress drop and the stress increase after holding arise from the relaxation phenomenon [17] and static strain aging [17,18], respectively. The magnitudes of stress drop and increase were defined as rr = r1-r0 and rs = r2r0, where r0, r1 and r2 are indicated in Figure 1b. As shown in Figure 1c, the stress drop arising from the relaxation increased with increasing strain rate from 1.7 10 5 to 1.7 10 4 s 1, and became constant in the strain rate range from 1.7 10 4 to 1.7 10 1 s 1. The stress increase arising from static strain aging linearly increased with increasing logarithmic strain rate. Figure 2a shows engineering stress–strain curves at initial strain rates of 1.7 10 5, 1.7 10 4, 1.7 10 3 and 1.7 10 2 s 1. The tensile deformations were stopped once at 69% strain, and then hydrogen was introduced under loading. If fracture caused by the hydrogen charging did not occur within 10 h, the specimen was deformed again until fracture at the same
Figure 1. (a) Engineering stress–strain curves with relaxation tests. (b) The curves outlined by broken lines in (a). (c) The relationships between strain rate and the stress drop due to relaxation, and between strain rate and the stress increase due to static strain aging.
Figure 2. (a) Tensile tests under hydrogen charging. The hydrogen was introduced at 69% tensile strain. (b) Engineering stresses plotted against hydrogen-charging time. HD indicates diffusible hydrogen content.
M. Koyama et al. / Scripta Materialia 66 (2012) 947–950
of 982 MPa. In Figure 4, the brittle fracture stress tested at the initial strain rate of 1.7 10 5 s 1 and the maximum stress tested at the initial strain rate of 1.7 10 3 s 1can be converted to true stresses of 1472 and 1828 MPa, respectively. Despite the fact that the brittle fracture stress was obviously lower than that of the ductile fracture stress, and diffusible hydrogen contents of them were comparable, HDF occurred at an initial strain rate of 1.7 10 5 s 1. This trend cannot be explained by the reduction in stress associated with the negative strain-rate sensitivity and the relaxation. In 1200
Strain rate: 1.7×10-5 s-1 Hydrogen was introduced at 49% HD: 6.7 wt.ppm
1000
Engineering stress (MPa)
strain rate with the initial strain rate. The holding time is limited to 10 h by the ability of the tensile test machine. Figure 2b shows stress–time curves corresponding to Figure 2a. HDF was observed at strain rates of 1.7 10 5 and 1.7 10 4 s 1. The fracture mode was intergranular as shown in Figure 3a and b. In contrast, HDF was not observed at strain rates of 1.7 10 3 and 1.7 10 2 s 1. The fracture mode observed after the tensile tests subsequent to constant strain holding was totally ductile as shown in Figure 3c and d. From the viewpoint of time to fracture, the HDF was suppressed by increasing strain rate as shown in Figure 2b. This is because the engineering flow stress at 69% strain decreased from 1053 to 999MPawith increasing strain rate from 1.7 10 5 to 1.7 10 2 s 1as shown in Figure 2a. Additionally, the stress drop associated with the relaxation [17] was observed, which was enhanced by increasing the strain rate from 1.7 10 5 to 1.7 10 4 s 1as shown in Figure 1c. These factors decreased the stress at the constant strain with increasing strain rate, suppressing the HDF. Another important trend of HDF is shown in Figure 4. The curve obtained at 1.7 10 3 s 1 is the same as that in Figure 2a. The other engineering stress–strain curve was obtained at an initial strain rate of 1.7 10 5 s 1. The tensile deformation was stopped at 49% strain, and the strain was subsequently held for 10 h under loading at constant strain with hydrogen charging. Then, the specimen was deformed again until fracture. The diffusible hydrogen contents of the specimen, which were measured after fracture, were 6.9 and 6.7 ppm, respectively. The test at the initial strain rate of 1.7 10 5 s 1 showed HDF at the engineering stress
949
800
Hydrogen charging 600
Strain rate: 1.7×10-3 s-1 Hydrogen was introduced at 69% HD: 6.9 wt.ppm
400
200
0 0
10
20 30 40 50 60 70 Engineering plastic strain (%)
80
90
Figure 4. The result of a tensile test including relaxation for 10 h at 49% strain at an initial strain rate of 1.7 10 5 s 1. The tensile curve obtained at the strain rate of 1.7 10 3 s 1 is the same as that of Figure 2a.
Figure 3. Fractographs showing the intergranular fracture surface at (a) low and (b) high magnifications that were provided by the deformation at 1.7 10 5 s 1 in Figure 2a, and showing the ductile fracture surface at (c) low and (d) high magnifications that were provided by the deformation at 1.7 10 2 s 1 in Figure 2a.
950
M. Koyama et al. / Scripta Materialia 66 (2012) 947–950
the previous studies [17,18] and Figure 1, static strain aging under loading was reported to significantly increase the strength of the material due to formation of carbon–dislocation pairs. The susceptibility to HDF is generally reported to become higher with increasing material strength level in the same alloy group [29–31]. The hardening mechanism of dynamic strain aging is similar to that of static strain aging, indicating that static strain aging as well as the promotion of dynamic strain aging by decreasing strain rate must be considered to deteriorate the resistance to HDF. The occurrence of HDF at relatively low stress in Figure 4 is related to the hardening caused by dynamic strain aging. If the specimen deformed at 1.7 10 3 s 1was exposed under loading for a long time, i.e. more than 10 h, it also would show HDF due to strengthening by static strain aging. The effect of the increase in material strength due to static and dynamic strain aging would be the third factor affecting the present HDF property. In conclusion, due to the influence of strain aging under loading, there are three factors affecting HDF: the negative strain-rate sensitivity of flow stress, the drop in stress caused by the relaxation phenomenon, and the increase in material strength due to strain aging. All three factors were strongly dependent on strain rate; the negative strain-rate sensitivity [2,24] and static strain aging [17,18] are widely accepted to arise from carbon diffusion. Hence, we note the following two points in Fe–Mn–C TWIP steels: 1. HDF is sensitive to the strain rate of predeformation. 2. It is important to consider the diffusion of both hydrogen and carbon diffusion when discussing HDF. M.K. acknowledges a Research Fellowship of NIMS Junior Researcher (2009-2010) and the Japan Society for the Promotion of Science for Young Scientists (2011). We would also like to acknowledge Dr. Takahiro Sawaguchi for taking part in the discussions during the experiments. POSCO supported this work by providing the samples and funding. [1] O. Gra¨ssel, G. Frommeyer, Mater. Sci. Technol. 14 (1998) 1213. [2] B.C. De Cooman, K.-G. Chin, J. Kim, High MnTWIP steels for automotive applications, in: M. Chiaberge (Ed.), New Trends and Developments in Automotive System Engineering, InTech., 2011, ISBN: 978-953-307-517-4. [3] O. Bouaziz, S. Allain, C.P. Scott, P. Cugy, D. Barbier, Curr. Opin. Solid State Mater. Sci. 15 (2011) 141.
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