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Static strain aging behavior of an X100 pipeline steel
Materials Science and Engineering A 550 (2012) 418–422
Contents lists available at SciVerse ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Static strain aging behavior of an X100 pipeline steel Wengui Zhao, Meng Chen, Shaohui Chen, Jinbo Qu ∗ Institute of Research of Iron and Steel (IRIS), Sha-steel, Jiangsu 215625, China
a r t i c l e
i n f o
Article history: Received 25 March 2012 Received in revised form 26 April 2012 Accepted 26 April 2012 Available online 4 May 2012 Keywords: Pipeline steel X100 Static strain aging Mechanical properties
a b s t r a c t The effect of static strain aging on the microstructure and mechanical properties of an industrially produced X100 pipeline steel was investigated. It was found that strain aging plays an important role in the tensile property and features of stress–strain curves, with the increase of either pre-strain or aging temperature, the strength increases while elongation decreases. The impact toughness was considerably affected by strain aging only for a pre-strain of 3%, which has a very low fraction of high-angle grain boundaries. © 2012 Elsevier B.V. All rights reserved.
1. Introduction Increasing demand for oil and natural gas calls for large amounts of steel for the construction of many pipelines in China [1]. Upgrading of pipeline steel from X60, X70 to X80 has been successfully achieved in the past two decades in order to increase transportation capacity and reduce cost [2–5]. At present, one of the major goals for steel industry is to further develop ultrahigh strength pipeline steels, e.g. X100, with good toughness and weldability. During the pipe making process, cold forming (with a strain of about 1%), followed by anti-corrosion coating at around 523 K, is essentially a static strain aging step, which will generally increase strength and decrease ductility. Because of its engineering significance, strain aging behavior has drawn increasing attention from the steel plate and pipe producers [6–12], particularly for the strainbased design applications. Yoo et al. [11] studied the microstructure and mechanical properties of an X80 pipeline steel and found that fine and well-distributed M/A (martensite/austenite) constituents were favorable for good mechanical properties and strain aging resistance. Gao et al. [12] conducted a strain aging test, i.e. 1% pre-strain plus 483 K aging, on a 22 mm thickness X80 plate and measured a 6% increase in both yield strength and yield ratio, which was close to that (9.5% in yield strength and 6.9% in yield ratio) after an industrial piping process. For ultrahigh strength grade pipeline steels, however, strain aging and its impact on microstructure and mechanical properties has not been adequately investigated.
∗ Corresponding author. Fax: +86 51258953902. E-mail address: [email protected] (J. Qu). 0921-5093/$ – see front matter © 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.04.095
This study was undertaken to determine the effect of static strain aging on an industrially produced X100 pipeline steel. Different strain aging parameters, i.e. pre-strains and aging temperatures, were selected to simulate the real pipe making process. The pre-strained and aged specimens were examined in terms of microstructural characteristics, tensile and impact mechanical properties. 2. Experimental The steel investigated was made at Hongfa steelmaking and casting plant of Sha-steel; the chemical composition is characterized by a low carbon content with additions of manganese, chromium, molybdenum, nickel and microalloying elements (niobium and titanium), as shown in Table 1. The 220 mm thickness slabs were hot rolled into 14.7 mm thickness plates by a strict thermo-mechanical controlled process (TMCP) at Shajing heavy plate plant of Sha-steel. The hot rolled plate used in this study has a yield strength of 695 MPa, ultimate tensile strength of 860 MPa (with a yield ratio of 0.81), total elongation of 19.0% at room temperature and Charpy V-notch (CVN) impact toughness of 256 J at −253 K. Rectangular samples (with a dimension of 300 by 30 mm) were cut from the plate with tensile pre-strains of 1%, 2% and 3% applied, respectively, perpendicular to the rolling direction. After the prestrains, the samples were aged for 30 min at 423, 473 and 523 K, respectively. Standard tensile and impact test specimens were machined from the pre-strained and aged samples. Round tensile specimens with a gauge diameter of 8 mm and length of 40 mm were tested at room temperature at a crosshead speed of 4 mm/min on an Instron universal testing machine (model:
W. Zhao et al. / Materials Science and Engineering A 550 (2012) 418–422
419
Table 1 Chemical composition of the X100 steel (wt.%). C
Si
Mn
Cr + Ni + Mo
Nb + Ti
Al
P
S
0.07
0.25
1.94
0.72
0.095
0.04
0.010
0.003
Instron 5585H, Canton, MA). Standard CVN impact specimens (10 mm × 10 mm × 55 mm) were tested at −253 K on an Instron impact testing machine (model: Instron SI-1M, Canton, MA). Microstructural analysis was carried out using a field emission scanning electron microscope (FESEM, model: JSM-7001F, Japan) equipped with an EBSD (model: Channel 5.0, HKL) and a high resolution transmission electron microscope (HRTEM, model: JEM2100F, Japan). Fractographs of the impact fractured specimens were also observed on the FESEM. 3. Results and discussion Fig. 1 shows TEM micrographs of the as-rolled and pre-strained (1%, 2% and 3%) and 473 K aged specimens. As shown in Fig. 1(a), the as-rolled microstructure is mainly composed of non-equiaxed lath ferrite, in the form of parallel platelets with a width of 250–400 nm, containing a large number of substructures and high density dislocations. With the increase of pre-strain from 1% to 3%, although there is no change in grain size, the density of dislocations in the parallel platelets increases gradually, as shown in Fig. 1(b)–(d). In the case of 3% pre-strain, in particular, sub-laths with a width of 30–50 nm, which are not entirely parallel but intersected with each other, clearly appear in the platelets, as illustrated in Fig. 1(d). The grain boundary misorientations of the sub-laths are mostly less than 15◦ , indicating low-angle grain boundaries (LAGBs). EBSD analysis was conducted on the specimens to determine the fraction of high-angle grain boundaries (HAGBs), as shown by
the grain boundary misorientation distributions in Fig. 2. It can be calculated from Fig. 2 that the fractions of HAGBs of the as-rolled, 2% and 3% pre-strained specimens (473 K aged) are 0.653, 0.633 and 0.338, respectively. The reason for the very low HAGB fraction of the 3% pre-strained specimen may be due to its large quantity of low misorientation sub-lath boundaries. For the specimens with the smallest pre-strain (1%) or lowest aging temperature (423 K), HAGB fractions are measured very close to that of the as-rolled specimen. Fig. 3 shows the stress–strain curves and Table 2 lists the mechanical properties of all the specimens. With the increase of either pre-strain or aging temperature, the strength increases while elongation decreases. The as-rolled specimen exhibits a continuous yielding behavior during tensile test, which is mainly attributed to the mobile dislocations formed in the matrix during TMCP. Many dislocation sources come into action at low strain and plastic flow begins simultaneously through the specimen, suppressing discontinuous yielding [13]. In contrast, the pre-strained and aged specimens show discontinuous yielding and yield plateaus, which is a characteristic of strain aging. Aging of pre-strained material allows the interstitial solute atoms to diffuse to the existing dislocations, forming Cottrell atmospheres, which prevent existing dislocation movement. It is believed that new dislocations must be generated to cause plastic strain. The applied stress must increase as observed to generate and propagate new dislocations, and results in the observed increase in strength. This is in good agreement with the TEM observations shown in Fig. 1.
Fig. 1. TEM micrographs of the as-rolled (a), 1% (b), 2% (c) and 3% (d) pre-strained and 473 K aged specimens.
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W. Zhao et al. / Materials Science and Engineering A 550 (2012) 418–422
Fig. 2. Grain boundary misorientation distributions of as-rolled (a), 2% (b) and 3% (c) pre-strained and 473 K aged specimens.
The 3% pre-strained specimens, in spite of aging temperature, exhibit a sharp yield point followed by a small plateau and then continuous drop of stress till fracture in their stress–strain curves. It is indicated that a 3% pre-strain introduces severe strain aging effect so that strain hardening does not occur. Impact toughness of the specimens is also given in Table 2. It can be seen that for the pre-strains of 1% and 2%, the aging temperature has virtually no influence on toughness. However, the 3% pre-strained and 473 and 523 K aged specimens have relatively low values, namely 109 J and 108 J. Fig. 4 shows SEM fractographs of some impact fractured specimens. The as-rolled and 2% pre-strained specimens experience a ductile fracture, featuring numerous fine and uniform dimples, while the 3% pre-strained specimen shows a cleavage fracture surface, to some extent.
The impact toughness can be related to the misorientations between the cracked grains and their neighboring grains. The change in operating slip system, which is necessary for a crack to cross a grain boundary, results in the grain boundaries becoming barriers for dislocation motion. The crack propagation direction will change after crossing a grain boundary, and hence the larger the misorientation between two grains, the more pronounced will be the deceleration in crack-growth rate at the grain boundary [14]. For HAGB, the formation of a dislocation free zone (DFZ) in a neighboring grain will depend on the operation of its own slip system, which will lead to a significant difference of direction and crack length between the two sides of the grain boundary [15]. Once a microcrack forms and propagates into the surrounding matrix, high misorientation boundary will act as an effective
Table 2 Mechanical properties of the pre-strained and aged specimens. YS (MPa)
UTS (MPa)
EL%
Yield ratio
−253 K Kv (J)
423
811 872 920
870 885 –
18.9 17.4 14.1
0.93 0.99 –
259 257 258
1% 2% 3%
473
828 900 942
877 905 –
17.8 15.3 12.9
0.94 0.99 –
219 232 109
1% 2% 3%
523
843 920 964
885 928 –
16.9 15.5 13.4
0.95 0.99 –
237 245 108
Pre-strain 1% 2% 3%
Aging temperature (K)
YS: yield strength, UTS: ultimate tensile strength, and EL: elongation.
W. Zhao et al. / Materials Science and Engineering A 550 (2012) 418–422
Fig. 3. Stress–strain curves of the as-rolled, pre-strained and 423 K (a), 473 K (b) and 523 K (c) aged specimens.
Fig. 4. SEM fractographs of the impact specimens: (a) as-rolled, (b) 2% pre-strained and 473 K aged and (c) 3% pre-strained and 473 K aged.
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obstacle [7,16–18]. In comparison, for LAGB, the small difference in orientation between neighboring grains will provide a relatively small obstacle. The above discussion is well supported by the measured impact toughness and HAGB fractions in this study. 4. Summary Static strain aging tests, with different pre-strains and aging temperatures, were performed on an industrially produced X100 steel. The following summary can be drawn: 1. The strength increased while elongation decreased with the increase of either pre-strain or aging temperature. After strain aging, the tensile stress–strain curves show discontinuous yielding behavior with characteristic yield plateaus. 2. The impact toughness was found to have no change for the pre-strains of 1% and 2%, but decrease for 3%, which can be explained by the fraction of high-angle grain boundaries. The 3% pre-strained and 473 K and 523 K aged specimens have low fractions of high-angle grain boundaries, leading to poor impact toughness and cleavage fracture. Acknowledgements This work has been financially supported by Sha-steel. Professor T. Emi is thanked due to constructive suggestions about the project work and continuing encouragements.
References [1] W.G. Zhao, W. Wang, S.H. Chen, J.B. Qu, Mater. Sci. Eng. A 528 (2011) 7417–7422. [2] X. Wang, F.R. Xiao, Y.H. Fu, X.W. Chen, B. Liao, Mater. Sci. Eng. A 530 (2011) 539–547. [3] B. Hwang, Y.M. Kim, S. Lee, N.J. Kim, S.S. Ahn, Metall. Mater. Trans. A (2005) 725–739. [4] S.Y. Han, S.Y. Shin, S.G. Lee, N.J. Kim, J.H. Bae, K. Kim, Metall. Mater. Trans. A (2010) 329–340. [5] Q.R. Xiong, Y.R. Feng, C.Y. Huo, W.W. Li, Mater. Mech. Eng. 12 (2005) 21–25. [6] A.J. Abdollah, L.R.O. Hein, M.S. Pereira, T.M. Hashimoto, Mater. Sci. Technol. 15 (1999) 1167–1170. [7] L. Himmel, K. Goodman, W.L. Haworth, in: R.A. Kot, B.L. Bramfitt (Eds.), Fundamental of Dual Phase Steels, AIME, New York, 1981, pp. 305–313. [8] S. Gündüz, Mater. Sci. Eng. A 486 (2008) 63–71. [9] R.G. Davies, in: R.A. Kot, B.L. Bramfitt (Eds.), Fundamental of Dual Phase Steels, AIME, New York, 1981, pp. 265–277. [10] M.S. Rashid, Metall. Trans. A 6 (1975) 1265–1272. [11] J.Y. Yoo, S.H. Chon, D.H. Seo, Pipeline Technology Conference, Ostend, 2009, pp. 12–14. [12] J.Z. Gao, Q.R. Ma, C.A. Wang, Y.L. Li, J.K. Li, Y. Xie, Mater. Mech. Eng. 34 (1) (2010) 5–8. [13] R.G. Davies, C.L. Magee, in: R.A. Kott, J.W. Morris (Eds.), Structure and Properties of Dual Phase Steels, AIME, New York, 1979, pp. 1–20. [14] Y. Zhong, F.R. Xiao, J.W. Zhang, Y.Y. Shan, W. Wang, K. Yang, Acta Mater. 54 (2006) 435–443. [15] D. Chen, M.E. Sixta, X.F. Zhang, L.C. De Jonghe, R.O. Ritchie, Acta Mater. 48 (2000) 4599–4608. [16] R.P. Krupitzer, Proceedings of the 22nd Mechanical Working Conference, ISSAIME, Toronto, October 29–30, 1980, pp. 1–21. [17] G.R. Speich, in: R.A. Kot, B.L. Bramfitt (Eds.), Fundamental of Dual Phase Steels, AIME, New York, 1981, pp. 3–45. [18] C. Lee, B.K. Zuidema, in: R. Pradhan (Ed.), Proceedings of the Symposium on High Strength Sheet Steels for the Automotive Industry, Iron and Steel Society, Warrendale, PA, 1994, pp. 103–110.
Contents lists available at SciVerse ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Static strain aging behavior of an X100 pipeline steel Wengui Zhao, Meng Chen, Shaohui Chen, Jinbo Qu ∗ Institute of Research of Iron and Steel (IRIS), Sha-steel, Jiangsu 215625, China
a r t i c l e
i n f o
Article history: Received 25 March 2012 Received in revised form 26 April 2012 Accepted 26 April 2012 Available online 4 May 2012 Keywords: Pipeline steel X100 Static strain aging Mechanical properties
a b s t r a c t The effect of static strain aging on the microstructure and mechanical properties of an industrially produced X100 pipeline steel was investigated. It was found that strain aging plays an important role in the tensile property and features of stress–strain curves, with the increase of either pre-strain or aging temperature, the strength increases while elongation decreases. The impact toughness was considerably affected by strain aging only for a pre-strain of 3%, which has a very low fraction of high-angle grain boundaries. © 2012 Elsevier B.V. All rights reserved.
1. Introduction Increasing demand for oil and natural gas calls for large amounts of steel for the construction of many pipelines in China [1]. Upgrading of pipeline steel from X60, X70 to X80 has been successfully achieved in the past two decades in order to increase transportation capacity and reduce cost [2–5]. At present, one of the major goals for steel industry is to further develop ultrahigh strength pipeline steels, e.g. X100, with good toughness and weldability. During the pipe making process, cold forming (with a strain of about 1%), followed by anti-corrosion coating at around 523 K, is essentially a static strain aging step, which will generally increase strength and decrease ductility. Because of its engineering significance, strain aging behavior has drawn increasing attention from the steel plate and pipe producers [6–12], particularly for the strainbased design applications. Yoo et al. [11] studied the microstructure and mechanical properties of an X80 pipeline steel and found that fine and well-distributed M/A (martensite/austenite) constituents were favorable for good mechanical properties and strain aging resistance. Gao et al. [12] conducted a strain aging test, i.e. 1% pre-strain plus 483 K aging, on a 22 mm thickness X80 plate and measured a 6% increase in both yield strength and yield ratio, which was close to that (9.5% in yield strength and 6.9% in yield ratio) after an industrial piping process. For ultrahigh strength grade pipeline steels, however, strain aging and its impact on microstructure and mechanical properties has not been adequately investigated.
∗ Corresponding author. Fax: +86 51258953902. E-mail address: [email protected] (J. Qu). 0921-5093/$ – see front matter © 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.04.095
This study was undertaken to determine the effect of static strain aging on an industrially produced X100 pipeline steel. Different strain aging parameters, i.e. pre-strains and aging temperatures, were selected to simulate the real pipe making process. The pre-strained and aged specimens were examined in terms of microstructural characteristics, tensile and impact mechanical properties. 2. Experimental The steel investigated was made at Hongfa steelmaking and casting plant of Sha-steel; the chemical composition is characterized by a low carbon content with additions of manganese, chromium, molybdenum, nickel and microalloying elements (niobium and titanium), as shown in Table 1. The 220 mm thickness slabs were hot rolled into 14.7 mm thickness plates by a strict thermo-mechanical controlled process (TMCP) at Shajing heavy plate plant of Sha-steel. The hot rolled plate used in this study has a yield strength of 695 MPa, ultimate tensile strength of 860 MPa (with a yield ratio of 0.81), total elongation of 19.0% at room temperature and Charpy V-notch (CVN) impact toughness of 256 J at −253 K. Rectangular samples (with a dimension of 300 by 30 mm) were cut from the plate with tensile pre-strains of 1%, 2% and 3% applied, respectively, perpendicular to the rolling direction. After the prestrains, the samples were aged for 30 min at 423, 473 and 523 K, respectively. Standard tensile and impact test specimens were machined from the pre-strained and aged samples. Round tensile specimens with a gauge diameter of 8 mm and length of 40 mm were tested at room temperature at a crosshead speed of 4 mm/min on an Instron universal testing machine (model:
W. Zhao et al. / Materials Science and Engineering A 550 (2012) 418–422
419
Table 1 Chemical composition of the X100 steel (wt.%). C
Si
Mn
Cr + Ni + Mo
Nb + Ti
Al
P
S
0.07
0.25
1.94
0.72
0.095
0.04
0.010
0.003
Instron 5585H, Canton, MA). Standard CVN impact specimens (10 mm × 10 mm × 55 mm) were tested at −253 K on an Instron impact testing machine (model: Instron SI-1M, Canton, MA). Microstructural analysis was carried out using a field emission scanning electron microscope (FESEM, model: JSM-7001F, Japan) equipped with an EBSD (model: Channel 5.0, HKL) and a high resolution transmission electron microscope (HRTEM, model: JEM2100F, Japan). Fractographs of the impact fractured specimens were also observed on the FESEM. 3. Results and discussion Fig. 1 shows TEM micrographs of the as-rolled and pre-strained (1%, 2% and 3%) and 473 K aged specimens. As shown in Fig. 1(a), the as-rolled microstructure is mainly composed of non-equiaxed lath ferrite, in the form of parallel platelets with a width of 250–400 nm, containing a large number of substructures and high density dislocations. With the increase of pre-strain from 1% to 3%, although there is no change in grain size, the density of dislocations in the parallel platelets increases gradually, as shown in Fig. 1(b)–(d). In the case of 3% pre-strain, in particular, sub-laths with a width of 30–50 nm, which are not entirely parallel but intersected with each other, clearly appear in the platelets, as illustrated in Fig. 1(d). The grain boundary misorientations of the sub-laths are mostly less than 15◦ , indicating low-angle grain boundaries (LAGBs). EBSD analysis was conducted on the specimens to determine the fraction of high-angle grain boundaries (HAGBs), as shown by
the grain boundary misorientation distributions in Fig. 2. It can be calculated from Fig. 2 that the fractions of HAGBs of the as-rolled, 2% and 3% pre-strained specimens (473 K aged) are 0.653, 0.633 and 0.338, respectively. The reason for the very low HAGB fraction of the 3% pre-strained specimen may be due to its large quantity of low misorientation sub-lath boundaries. For the specimens with the smallest pre-strain (1%) or lowest aging temperature (423 K), HAGB fractions are measured very close to that of the as-rolled specimen. Fig. 3 shows the stress–strain curves and Table 2 lists the mechanical properties of all the specimens. With the increase of either pre-strain or aging temperature, the strength increases while elongation decreases. The as-rolled specimen exhibits a continuous yielding behavior during tensile test, which is mainly attributed to the mobile dislocations formed in the matrix during TMCP. Many dislocation sources come into action at low strain and plastic flow begins simultaneously through the specimen, suppressing discontinuous yielding [13]. In contrast, the pre-strained and aged specimens show discontinuous yielding and yield plateaus, which is a characteristic of strain aging. Aging of pre-strained material allows the interstitial solute atoms to diffuse to the existing dislocations, forming Cottrell atmospheres, which prevent existing dislocation movement. It is believed that new dislocations must be generated to cause plastic strain. The applied stress must increase as observed to generate and propagate new dislocations, and results in the observed increase in strength. This is in good agreement with the TEM observations shown in Fig. 1.
Fig. 1. TEM micrographs of the as-rolled (a), 1% (b), 2% (c) and 3% (d) pre-strained and 473 K aged specimens.
420
W. Zhao et al. / Materials Science and Engineering A 550 (2012) 418–422
Fig. 2. Grain boundary misorientation distributions of as-rolled (a), 2% (b) and 3% (c) pre-strained and 473 K aged specimens.
The 3% pre-strained specimens, in spite of aging temperature, exhibit a sharp yield point followed by a small plateau and then continuous drop of stress till fracture in their stress–strain curves. It is indicated that a 3% pre-strain introduces severe strain aging effect so that strain hardening does not occur. Impact toughness of the specimens is also given in Table 2. It can be seen that for the pre-strains of 1% and 2%, the aging temperature has virtually no influence on toughness. However, the 3% pre-strained and 473 and 523 K aged specimens have relatively low values, namely 109 J and 108 J. Fig. 4 shows SEM fractographs of some impact fractured specimens. The as-rolled and 2% pre-strained specimens experience a ductile fracture, featuring numerous fine and uniform dimples, while the 3% pre-strained specimen shows a cleavage fracture surface, to some extent.
The impact toughness can be related to the misorientations between the cracked grains and their neighboring grains. The change in operating slip system, which is necessary for a crack to cross a grain boundary, results in the grain boundaries becoming barriers for dislocation motion. The crack propagation direction will change after crossing a grain boundary, and hence the larger the misorientation between two grains, the more pronounced will be the deceleration in crack-growth rate at the grain boundary [14]. For HAGB, the formation of a dislocation free zone (DFZ) in a neighboring grain will depend on the operation of its own slip system, which will lead to a significant difference of direction and crack length between the two sides of the grain boundary [15]. Once a microcrack forms and propagates into the surrounding matrix, high misorientation boundary will act as an effective
Table 2 Mechanical properties of the pre-strained and aged specimens. YS (MPa)
UTS (MPa)
EL%
Yield ratio
−253 K Kv (J)
423
811 872 920
870 885 –
18.9 17.4 14.1
0.93 0.99 –
259 257 258
1% 2% 3%
473
828 900 942
877 905 –
17.8 15.3 12.9
0.94 0.99 –
219 232 109
1% 2% 3%
523
843 920 964
885 928 –
16.9 15.5 13.4
0.95 0.99 –
237 245 108
Pre-strain 1% 2% 3%
Aging temperature (K)
YS: yield strength, UTS: ultimate tensile strength, and EL: elongation.
W. Zhao et al. / Materials Science and Engineering A 550 (2012) 418–422
Fig. 3. Stress–strain curves of the as-rolled, pre-strained and 423 K (a), 473 K (b) and 523 K (c) aged specimens.
Fig. 4. SEM fractographs of the impact specimens: (a) as-rolled, (b) 2% pre-strained and 473 K aged and (c) 3% pre-strained and 473 K aged.
421
422
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obstacle [7,16–18]. In comparison, for LAGB, the small difference in orientation between neighboring grains will provide a relatively small obstacle. The above discussion is well supported by the measured impact toughness and HAGB fractions in this study. 4. Summary Static strain aging tests, with different pre-strains and aging temperatures, were performed on an industrially produced X100 steel. The following summary can be drawn: 1. The strength increased while elongation decreased with the increase of either pre-strain or aging temperature. After strain aging, the tensile stress–strain curves show discontinuous yielding behavior with characteristic yield plateaus. 2. The impact toughness was found to have no change for the pre-strains of 1% and 2%, but decrease for 3%, which can be explained by the fraction of high-angle grain boundaries. The 3% pre-strained and 473 K and 523 K aged specimens have low fractions of high-angle grain boundaries, leading to poor impact toughness and cleavage fracture. Acknowledgements This work has been financially supported by Sha-steel. Professor T. Emi is thanked due to constructive suggestions about the project work and continuing encouragements.
References [1] W.G. Zhao, W. Wang, S.H. Chen, J.B. Qu, Mater. Sci. Eng. A 528 (2011) 7417–7422. [2] X. Wang, F.R. Xiao, Y.H. Fu, X.W. Chen, B. Liao, Mater. Sci. Eng. A 530 (2011) 539–547. [3] B. Hwang, Y.M. Kim, S. Lee, N.J. Kim, S.S. Ahn, Metall. Mater. Trans. A (2005) 725–739. [4] S.Y. Han, S.Y. Shin, S.G. Lee, N.J. Kim, J.H. Bae, K. Kim, Metall. Mater. Trans. A (2010) 329–340. [5] Q.R. Xiong, Y.R. Feng, C.Y. Huo, W.W. Li, Mater. Mech. Eng. 12 (2005) 21–25. [6] A.J. Abdollah, L.R.O. Hein, M.S. Pereira, T.M. Hashimoto, Mater. Sci. Technol. 15 (1999) 1167–1170. [7] L. Himmel, K. Goodman, W.L. Haworth, in: R.A. Kot, B.L. Bramfitt (Eds.), Fundamental of Dual Phase Steels, AIME, New York, 1981, pp. 305–313. [8] S. Gündüz, Mater. Sci. Eng. A 486 (2008) 63–71. [9] R.G. Davies, in: R.A. Kot, B.L. Bramfitt (Eds.), Fundamental of Dual Phase Steels, AIME, New York, 1981, pp. 265–277. [10] M.S. Rashid, Metall. Trans. A 6 (1975) 1265–1272. [11] J.Y. Yoo, S.H. Chon, D.H. Seo, Pipeline Technology Conference, Ostend, 2009, pp. 12–14. [12] J.Z. Gao, Q.R. Ma, C.A. Wang, Y.L. Li, J.K. Li, Y. Xie, Mater. Mech. Eng. 34 (1) (2010) 5–8. [13] R.G. Davies, C.L. Magee, in: R.A. Kott, J.W. Morris (Eds.), Structure and Properties of Dual Phase Steels, AIME, New York, 1979, pp. 1–20. [14] Y. Zhong, F.R. Xiao, J.W. Zhang, Y.Y. Shan, W. Wang, K. Yang, Acta Mater. 54 (2006) 435–443. [15] D. Chen, M.E. Sixta, X.F. Zhang, L.C. De Jonghe, R.O. Ritchie, Acta Mater. 48 (2000) 4599–4608. [16] R.P. Krupitzer, Proceedings of the 22nd Mechanical Working Conference, ISSAIME, Toronto, October 29–30, 1980, pp. 1–21. [17] G.R. Speich, in: R.A. Kot, B.L. Bramfitt (Eds.), Fundamental of Dual Phase Steels, AIME, New York, 1981, pp. 3–45. [18] C. Lee, B.K. Zuidema, in: R. Pradhan (Ed.), Proceedings of the Symposium on High Strength Sheet Steels for the Automotive Industry, Iron and Steel Society, Warrendale, PA, 1994, pp. 103–110.