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Correlation between cementite precipitation and Portevin-Le Chatelier effect in a hot-rolled medium Mn steel
Materials Letters 258 (2020) 126796
Contents lists available at ScienceDirect
Materials Letters journal homepage: www.elsevier.com/locate/mlblue
Correlation between cementite precipitation and Portevin-Le Chatelier effect in a hot-rolled medium Mn steel Yongjin Wang ⇑, Zetian Ma, Renbo Song ⇑, Shuai Zhao, Zherui Zhang, Weifeng Huo School of Materials Science and Engineering, University of Science and Technology Beijing, Beijing 100083, China
a r t i c l e
i n f o
Article history: Received 26 September 2019 Accepted 9 October 2019 Available online 10 October 2019 Keywords: Hot-rolled medium Mn steel Intercritical annealing Microstructure Cementite precipitation Portevin-Le Chatelier effect Deformation and fracture
a b s t r a c t The correlation between cementite precipitation and plastic instability phenomena especially Portevin-Le Chatelier (PLC) effect was investigated for a hot-rolled medium Mn steel. The microstructure demonstrated lamellar morphology of austenite and ferrite. Cementite precipitations could be retained by controlling certain intercritical annealing (IA) process. The initial high density of dislocations in ferrite contributed to the continuous yielding. Cementite precipitations lowered solute carbon content in ferrite and the dynamic interaction with dislocations was eliminated. A moderate product of strength and elongation (PSE) grades hot-rolled medium Mn steel without PLC effect during tensile test was obtained by controlling the IA process. Ó 2019 Elsevier B.V. All rights reserved.
1. Introduction
2. Experimental
Recently, medium Mn steels containing 5–12 wt% Mn have been considered as a promising candidate for the 3rd generation of advanced high strength steel [1]. High strength and high plasticity could be obtained due to the large fraction of retained austenite. However, some plastic instability phenomena such as Luders bands and Portevin-Le Chatelier (PLC) bands usually occur during the tensile test of medium Mn steels [2–5]. The high yield ratio (YR), discontinuous work hardening and serrated plastic flow deteriorate the formability and limit its practical application [6]. The multi-phase microstructure of medium-Mn steel makes it more complicated to analyze the PLC phenomenon. Microalloyed carbide particles have been attempted to improve the strength [7]. Carbide precipitation is reported to affect the microstructure characterization and failure behavior of ferritic steel [8]. However, the role of carbide precipitations during plastic deformation is uncertain. In the present study, the correlation between cementite precipitation and the serrated flow during plastic deformation is investigated. Besides, the microstructure evolution is clarified and its relationship with mechanical properties is studied.
The chemical composition of the investigated steel is 0.3 C, 9.3 Mn, 2.15 Al, 0.028Si, 0.008 S and balance Fe (wt%). A 25 kg ingot was prepared using a vacuum induction furnace, homogenized at 1200 °C for 2 h and hot rolled to ~ 3 mm thick plates, followed by laminar cooling to 450 °C and air cooling to ambient temperature. As shown in Fig. S1 in supplementary material, thermodynamic calculations suggest that the phase transformation of Ae3 and dissolution temperature of cementite are 762 °C and 645 °C, respectively. The samples were intercritically annealed (IA) at 640 ~ 720 °C for 30 min or 680 °C for 2 ~ 120 min followed by air cooling, respectively. For convenience, the specimens annealed at 640, 660, 680, 700 and 720 °C for 30 min are referred as IA64030, IA660-30, IA680-30, IA700-30 and IA720-30. The specimens annealed at 680 °C for 2, 10, 30, 60, 120 min are referred as IA680-2, IA680-10, IA680-30, IA680-60, IA680-120. Tensile tests were carried out using CMT5105 electronic universal experiment machine at ambient temperature with the strain rate of 10 3 s 1. Bone-like shape specimens were cut with the gauge length of 25 mm along rolling direction. Microstructure observations were conducted using transmission electron microscope (TEM, JEM2100) and electron backscattered diffraction (EBSD, EDAX OIM6). 3. Results and discussion
⇑ Corresponding authors. E-mail addresses: [email protected] (Y. Wang), [email protected]. cn (R. Song). https://doi.org/10.1016/j.matlet.2019.126796 0167-577X/Ó 2019 Elsevier B.V. All rights reserved.
As shown in Fig. 1(a) and (b), all the engineering stress-strain curves demonstrate continuous yielding behavior. No obvious
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Y. Wang et al. / Materials Letters 258 (2020) 126796
yielding point or yielding elongation can be observed. As for the IA680-30 specimen, the ultimate tensile strength (UST), yield strength (YS) and total elongation (TE) are 1123.3 MPa, 726.5 MPa and 55.7%, respectively. The highest product of UTS TE (PSE) can be obtained with the value of 62.5 GPa·% and yield ratio (YR) is 0.61. The PSE can reach above 50 GPa·% for IA660-30, IA680-60, IA680-120 and IA700-30 specimens. However, obvious serrated flow phenomenon can be observed during the strain hardening stage for specimens with PSE over 50 GPa·%. The stress serrated flow is usually referred as the PLC effect [9]. As for the curves of IA640-30 and IA680-10 specimens, stable strain hardening behavior without serrated flow can be seen. The PSE value for IA640-30 and IA680-10 specimens falls in a moderate over 30 GPa·% level (Table 1). The continuous yielding, stable strain hardening, reasonable YR (0.77 ~ 0.78) and moderate PSE contribute to a good formability and safe performance. The stressstrain behavior, especially strain hardening behavior, is sensitive to the process and microstructure of medium Mn steel. The explanation about the plastic instability such as Luders bands or PLC bands is divided into several hypotheses [10–12]. Sun et al. [13]
concluded that the serrated flow was attributed to the discontinuous TRIP effect, but direct evidence was missing. Yang et al. [14] proposed that the PLC bands werre in relation to the interaction of C atoms/C-Mn pairs and dislocations. The mechanism of a smooth strain hardening behavior without serrated flow needs to be studied. Microstructure evolution of specimens under different IA conditions is shown in Fig. 2. The microstructure after IA640-30, IA66030 and IA680-30 mainly consists of lamellar retained austenite (cL) and ferrite (aL). The lamellar morphology of austenite and ferrite is a common microstructure for hot rolled medium-Mn steel. The volume fractions of austenite for IA640-30, IA660-30 and IA68030 specimens are 35.1%, 56.6%, 63.2%, respectively (Fig. S2). Many cementite precipitations can be seen for IA640-30 specimen, while none cementite can be observed for IA660-30 and IA680-30 specimens. As for the microstructure after different IA times, martensite lath and large size blocky austenite can be seen for IA680-2 specimen, which is similar with the hot rolled microstructure (Fig. S3a). The volume fraction of austenite is 21.4% for IA680-2 specimen. It seems that annealing for 2 min is not long enough
Fig. 1. Engineering stress-strain curves after different temperatures (a) and different soaking times (b); Strain hardening rate (dr/de) curves of IA640-30, IA680-10 and IA680-30 specimens (c).
Table 1 Mechanical data for the specimens under different IA conditions. Specimen
UTS/MPa
YS/MPa
TE/%
PSE/GPa·%
YR
Specimen
UTS/MPa
YS/MPa
TE/%
PSE/GPa·%
YR
640–30 660–30 680–30 700–30 720–30
1017.2 1111.8 1123.3 1120.6 1183.4
782.3 726.5 690.3 699.3 842.4
33.9 52.6 55.7 52.8 1.04
34.5 58.4 62.5 59.2 1.2
0.76 0.65 0.61 0.62 0.71
680–2 680–10 680–60 680–120 —
1182.3 994.8 969.3 1000.0 —
992.1 767.0 667.1 639.9 —
19.8 39.7 54.2 55.7 —
23.4 39.5 52.6 55.7 —
0.84 0.77 0.68 0.64 —
Y. Wang et al. / Materials Letters 258 (2020) 126796
3
Fig. 2. SEM micrographs of specimens under different conditions and EBSD phase image (yellow: FCC; grey: BCC): (a, b) IA640-30; (c) IA660-30; (d) IA680-30; (e) IA680-2; (f) IA680-10. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
for the austenite reversion and the high strength and low elongation is consistent with the microstructure. For IA680-10 specimen, lamellar morphology of austenite and ferrite can be seen and the volume fraction of austenite increases up to 41.5%. An interesting characterization is obvious that cementite particles are retained for IA680-10 specimen. The cementite particles may not fully dissolve into the matrix since the annealing time is relatively short.
By comparing the stress-strain curves and microstructure, we can see that high fraction of austenite contributes to a high TE and PSE. However, all the specimens with PSE over 50 GPa·% show serrated flow phenomenon during the strain hardening stage, while IA640-30 and IA680-10 specimens show a stable strain hardening. For specimens with reasonable PSE (>30 GPa·%), the microstructures show a matrix of lamellar austenite and ferrite
Fig. 3. TEM results on different specimen: bright field images of (a) IA680-30 and (b) IA640-30 specimens, dark field images of (c) IA640-30 specimen; EBSD kernel average misorientation (KAM) mappings on IA 640-30 specimen before deformation (d) and after fracture (e).
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Y. Wang et al. / Materials Letters 258 (2020) 126796
and the main difference is the cementite precipitations for IA64030 and IA680-10 specimens. The representative strain hardening curves in Fig. 1(c) show a moving trend with sharp decreasing, plateau decreasing and finally fracture decreasing. But the IA680-30 specimen with serrated flow during tensile test shows violent vibration during the plateau decreasing stage. It can be inferred that the matrix microstructure of lamellar austenite and ferrite determines the profile of stress-strain curve and the serrated flow or violent vibration is somewhat related with the cementite precipitation. Fig. 3(a) shows the TEM image of IA680-30 specimen and the lamellar microstructure of austenite and ferrite is in good agreement with the SEM results in Fig. 2. The average width of the lath is about 300 nm. The dislocation substructure can be observed in a two-beam condition. The interface between ferrite and austenite shows a concentration of dislocations. Many dislocations can be observed inside the ferrite lath. For the IA640-30 specimen, we can see similar dislocation distribution in Fig. 3(b). High density of dislocations can be observed at both the interface and inner of the ferrite lath. The dark field (DF) image (Fig. 3(c)) shows many cementite particles precipitate inside the ferrite lath. The cementite demonstrates two types of morphologies: rod-like and spherical particles. Combining the two-beam and DF image, we can see that the dislocation tangles are located near the cementite particles. Fig. 3(d) shows the kernel average misorientation (KAM) mappings of IA640-30 specimen before deformation. KAM is somewhat in proportional with the density of geometry necessary dislocations (GNDs) [15]. The KAM values are higher at the interface between lamellar austenite and ferrite. As shown in Fig. 3(e), the KAM values increases obviously after fracture. The deformation occurs at both the ferrite and austenite. Secondary deformation mechanism like deformation twinning can be observed at the austenite and the fraction of austenite decreases lower than 10%. The good match of continuous yielding and stable strain hardening helps improve the formability of the medium Mn steel. The initial high density of GNDs for hot-rolled medium Mn steel is considered as the results of the inactive recovery [16–18]. The unpinned GNDs in ferrite contribute to the continuous yielding, which is similar with the DP steel [19]. The PLC effect is related with the dynamic strain aging (DSA). All the specimens demonstrate the lamellar morphology of austenite and ferrite. The stress serration at the tensile curves disappears when the cementite precipitations exist. The cementite particles for IA640-30 and IA68010 specimens mainly precipitate from the ferrite and it will lower the C content in ferrite. Therefore, the ferrite grains with less solute carbon atoms will weaken the dynamic interaction with dislocations during strain hardening. Besides, the grain size of prior austenite is about 35 lm (Fig. S3b) and the coarse grains are reported to help remove the serrated flow [20,21]. 4. Conclusions A moderate PSE grades hot-rolled medium Mn steel without plastic instability phenomena is obtained by controlling the IA process. The microstructure demonstrates lamellar morphology of
austenite and ferrite after IA process. Cementite precipitations exist for IA640-30 and IA680-10 specimens. The initial high density of GNDs in ferrite contributes to the continuous yielding. Cementite precipitations lower the solute carbon content in ferrite. Together with the coarse grains, the dynamic interaction with dislocations is eliminated. Overall, the plastic instability phenomena, especially PLC effect, can be overcome by controlling the cementite precipitation for hot-rolled medium Mn steel. Declaration of Competing Interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Acknowledgements This work was supported by the Fundamental Research Funds for the Central Universities (No. FRF-TP-18-039A1); the Project funded by China Postdoctoral Science Foundation (No. 2019M650482). Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi.org/10.1016/j.matlet.2019.126796. References [1] S.P. Chen, R. Rana, A. Haldar, R.K. Ray, Prog. Mater. Sci. 89 (2017) 345–391. [2] H.W. Luo, H. Dong, M.X. Huang, Mater. Des. 83 (2015) 42–48. [3] H. Kamoutsi, E. Gioti, N. Gregory, Z. Cai, H. Ding, Metall. Mater. Trans. A 46 (2015) 4841–4846. [4] R. Schwab, V. Ruff, Acta Mater. 61 (2013) 1798–1808. [5] E. Emadoddin, A. Akbarzadeh, G.H. Daneshi, Mater. Sci. Eng. A 447 (2007) 174– 179. [6] H. Aboulfadl, J. Deges, P. Choi, D. Raabe, Acta Mater. 86 (2015) 34–42. [7] H.J. Pan, H. Ding, M.H. Cai, D. Kibaroglu, Y. Man, W.W. Song, Mater. Sci. Eng. A 766 (2019) 138371. [8] B.B. Straumal, Y.Q. Kucheev, L.I. Efron, A.L. Petelin, J.D. Majumdar, I. Manna, J. Mater. Eng. Perform. 21 (2012) 667–670. [9] X.G. Wang, L. Wang, M.X. Huang, Acta Mater. 124 (2017) 17–29. [10] R.K. Rai, J.K. Sahu, Mater. Lett. 210 (2018) 298–300. [11] M. Callahan, O. Hubert, F. Hild, A. Perlade, J.H. Schmitt, Mater. Sci. Eng. A 704 (2017) 391–400. [12] M.H. Zhang, L.F. Li, J. Ding, Q.B. Wu, Y.D. Wang, J. Almer, F.M. Guo, Y. Ren, Acta Mater. 141 (2017) 294–303. [13] B.H. Sun, N. Vanderesse, F. Fazeli, C. Scott, J.Q. Chen, P. Bocher, M. Jahazi, S. Yue, Scr. Mater. 133 (2017) 9–13. [14] F. Yang, H.W. Luo, E.X. Pu, S.L. Zhang, H. Dong, Int. J. Plast. 103 (2018) 188–202. [15] B. Hu, H.W. Luo, Acta Mater. 176 (2019) 250–263. [16] B.B. He, M. Wang, M.X. Huang, Metall. Mater. Trans. A 50 (2019) 2971–2977. [17] E. Welsch, D. Ponge, H. Hafez, S. Sandlobes, P. Choi, M. Herbig, S. Zaefferer, D. Raabe, Acta Mater. 116 (2016) 188–199. [18] J.H. Han, S.J. Lee, J.G. Jung, Y.K. Lee, Acta Mater. 78 (2014) 369–377. [19] M. Jafari, N. Saeidi, S. Ziaei-Rad, M. Jamshidian, H.S. Kim, Mech. Mater. 134 (2019) 132–142. [20] W.H. Wang, D. Wu, R.S. Chen, X.N. Zhang, J. Mater. Sci. Technol. 34 (2018) 1236–1242. [21] D. Wu, R.S. Chen, E.H. Han, Mater. Sci. Eng. A 532 (2012) 267–274.
Contents lists available at ScienceDirect
Materials Letters journal homepage: www.elsevier.com/locate/mlblue
Correlation between cementite precipitation and Portevin-Le Chatelier effect in a hot-rolled medium Mn steel Yongjin Wang ⇑, Zetian Ma, Renbo Song ⇑, Shuai Zhao, Zherui Zhang, Weifeng Huo School of Materials Science and Engineering, University of Science and Technology Beijing, Beijing 100083, China
a r t i c l e
i n f o
Article history: Received 26 September 2019 Accepted 9 October 2019 Available online 10 October 2019 Keywords: Hot-rolled medium Mn steel Intercritical annealing Microstructure Cementite precipitation Portevin-Le Chatelier effect Deformation and fracture
a b s t r a c t The correlation between cementite precipitation and plastic instability phenomena especially Portevin-Le Chatelier (PLC) effect was investigated for a hot-rolled medium Mn steel. The microstructure demonstrated lamellar morphology of austenite and ferrite. Cementite precipitations could be retained by controlling certain intercritical annealing (IA) process. The initial high density of dislocations in ferrite contributed to the continuous yielding. Cementite precipitations lowered solute carbon content in ferrite and the dynamic interaction with dislocations was eliminated. A moderate product of strength and elongation (PSE) grades hot-rolled medium Mn steel without PLC effect during tensile test was obtained by controlling the IA process. Ó 2019 Elsevier B.V. All rights reserved.
1. Introduction
2. Experimental
Recently, medium Mn steels containing 5–12 wt% Mn have been considered as a promising candidate for the 3rd generation of advanced high strength steel [1]. High strength and high plasticity could be obtained due to the large fraction of retained austenite. However, some plastic instability phenomena such as Luders bands and Portevin-Le Chatelier (PLC) bands usually occur during the tensile test of medium Mn steels [2–5]. The high yield ratio (YR), discontinuous work hardening and serrated plastic flow deteriorate the formability and limit its practical application [6]. The multi-phase microstructure of medium-Mn steel makes it more complicated to analyze the PLC phenomenon. Microalloyed carbide particles have been attempted to improve the strength [7]. Carbide precipitation is reported to affect the microstructure characterization and failure behavior of ferritic steel [8]. However, the role of carbide precipitations during plastic deformation is uncertain. In the present study, the correlation between cementite precipitation and the serrated flow during plastic deformation is investigated. Besides, the microstructure evolution is clarified and its relationship with mechanical properties is studied.
The chemical composition of the investigated steel is 0.3 C, 9.3 Mn, 2.15 Al, 0.028Si, 0.008 S and balance Fe (wt%). A 25 kg ingot was prepared using a vacuum induction furnace, homogenized at 1200 °C for 2 h and hot rolled to ~ 3 mm thick plates, followed by laminar cooling to 450 °C and air cooling to ambient temperature. As shown in Fig. S1 in supplementary material, thermodynamic calculations suggest that the phase transformation of Ae3 and dissolution temperature of cementite are 762 °C and 645 °C, respectively. The samples were intercritically annealed (IA) at 640 ~ 720 °C for 30 min or 680 °C for 2 ~ 120 min followed by air cooling, respectively. For convenience, the specimens annealed at 640, 660, 680, 700 and 720 °C for 30 min are referred as IA64030, IA660-30, IA680-30, IA700-30 and IA720-30. The specimens annealed at 680 °C for 2, 10, 30, 60, 120 min are referred as IA680-2, IA680-10, IA680-30, IA680-60, IA680-120. Tensile tests were carried out using CMT5105 electronic universal experiment machine at ambient temperature with the strain rate of 10 3 s 1. Bone-like shape specimens were cut with the gauge length of 25 mm along rolling direction. Microstructure observations were conducted using transmission electron microscope (TEM, JEM2100) and electron backscattered diffraction (EBSD, EDAX OIM6). 3. Results and discussion
⇑ Corresponding authors. E-mail addresses: [email protected] (Y. Wang), [email protected]. cn (R. Song). https://doi.org/10.1016/j.matlet.2019.126796 0167-577X/Ó 2019 Elsevier B.V. All rights reserved.
As shown in Fig. 1(a) and (b), all the engineering stress-strain curves demonstrate continuous yielding behavior. No obvious
2
Y. Wang et al. / Materials Letters 258 (2020) 126796
yielding point or yielding elongation can be observed. As for the IA680-30 specimen, the ultimate tensile strength (UST), yield strength (YS) and total elongation (TE) are 1123.3 MPa, 726.5 MPa and 55.7%, respectively. The highest product of UTS TE (PSE) can be obtained with the value of 62.5 GPa·% and yield ratio (YR) is 0.61. The PSE can reach above 50 GPa·% for IA660-30, IA680-60, IA680-120 and IA700-30 specimens. However, obvious serrated flow phenomenon can be observed during the strain hardening stage for specimens with PSE over 50 GPa·%. The stress serrated flow is usually referred as the PLC effect [9]. As for the curves of IA640-30 and IA680-10 specimens, stable strain hardening behavior without serrated flow can be seen. The PSE value for IA640-30 and IA680-10 specimens falls in a moderate over 30 GPa·% level (Table 1). The continuous yielding, stable strain hardening, reasonable YR (0.77 ~ 0.78) and moderate PSE contribute to a good formability and safe performance. The stressstrain behavior, especially strain hardening behavior, is sensitive to the process and microstructure of medium Mn steel. The explanation about the plastic instability such as Luders bands or PLC bands is divided into several hypotheses [10–12]. Sun et al. [13]
concluded that the serrated flow was attributed to the discontinuous TRIP effect, but direct evidence was missing. Yang et al. [14] proposed that the PLC bands werre in relation to the interaction of C atoms/C-Mn pairs and dislocations. The mechanism of a smooth strain hardening behavior without serrated flow needs to be studied. Microstructure evolution of specimens under different IA conditions is shown in Fig. 2. The microstructure after IA640-30, IA66030 and IA680-30 mainly consists of lamellar retained austenite (cL) and ferrite (aL). The lamellar morphology of austenite and ferrite is a common microstructure for hot rolled medium-Mn steel. The volume fractions of austenite for IA640-30, IA660-30 and IA68030 specimens are 35.1%, 56.6%, 63.2%, respectively (Fig. S2). Many cementite precipitations can be seen for IA640-30 specimen, while none cementite can be observed for IA660-30 and IA680-30 specimens. As for the microstructure after different IA times, martensite lath and large size blocky austenite can be seen for IA680-2 specimen, which is similar with the hot rolled microstructure (Fig. S3a). The volume fraction of austenite is 21.4% for IA680-2 specimen. It seems that annealing for 2 min is not long enough
Fig. 1. Engineering stress-strain curves after different temperatures (a) and different soaking times (b); Strain hardening rate (dr/de) curves of IA640-30, IA680-10 and IA680-30 specimens (c).
Table 1 Mechanical data for the specimens under different IA conditions. Specimen
UTS/MPa
YS/MPa
TE/%
PSE/GPa·%
YR
Specimen
UTS/MPa
YS/MPa
TE/%
PSE/GPa·%
YR
640–30 660–30 680–30 700–30 720–30
1017.2 1111.8 1123.3 1120.6 1183.4
782.3 726.5 690.3 699.3 842.4
33.9 52.6 55.7 52.8 1.04
34.5 58.4 62.5 59.2 1.2
0.76 0.65 0.61 0.62 0.71
680–2 680–10 680–60 680–120 —
1182.3 994.8 969.3 1000.0 —
992.1 767.0 667.1 639.9 —
19.8 39.7 54.2 55.7 —
23.4 39.5 52.6 55.7 —
0.84 0.77 0.68 0.64 —
Y. Wang et al. / Materials Letters 258 (2020) 126796
3
Fig. 2. SEM micrographs of specimens under different conditions and EBSD phase image (yellow: FCC; grey: BCC): (a, b) IA640-30; (c) IA660-30; (d) IA680-30; (e) IA680-2; (f) IA680-10. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
for the austenite reversion and the high strength and low elongation is consistent with the microstructure. For IA680-10 specimen, lamellar morphology of austenite and ferrite can be seen and the volume fraction of austenite increases up to 41.5%. An interesting characterization is obvious that cementite particles are retained for IA680-10 specimen. The cementite particles may not fully dissolve into the matrix since the annealing time is relatively short.
By comparing the stress-strain curves and microstructure, we can see that high fraction of austenite contributes to a high TE and PSE. However, all the specimens with PSE over 50 GPa·% show serrated flow phenomenon during the strain hardening stage, while IA640-30 and IA680-10 specimens show a stable strain hardening. For specimens with reasonable PSE (>30 GPa·%), the microstructures show a matrix of lamellar austenite and ferrite
Fig. 3. TEM results on different specimen: bright field images of (a) IA680-30 and (b) IA640-30 specimens, dark field images of (c) IA640-30 specimen; EBSD kernel average misorientation (KAM) mappings on IA 640-30 specimen before deformation (d) and after fracture (e).
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Y. Wang et al. / Materials Letters 258 (2020) 126796
and the main difference is the cementite precipitations for IA64030 and IA680-10 specimens. The representative strain hardening curves in Fig. 1(c) show a moving trend with sharp decreasing, plateau decreasing and finally fracture decreasing. But the IA680-30 specimen with serrated flow during tensile test shows violent vibration during the plateau decreasing stage. It can be inferred that the matrix microstructure of lamellar austenite and ferrite determines the profile of stress-strain curve and the serrated flow or violent vibration is somewhat related with the cementite precipitation. Fig. 3(a) shows the TEM image of IA680-30 specimen and the lamellar microstructure of austenite and ferrite is in good agreement with the SEM results in Fig. 2. The average width of the lath is about 300 nm. The dislocation substructure can be observed in a two-beam condition. The interface between ferrite and austenite shows a concentration of dislocations. Many dislocations can be observed inside the ferrite lath. For the IA640-30 specimen, we can see similar dislocation distribution in Fig. 3(b). High density of dislocations can be observed at both the interface and inner of the ferrite lath. The dark field (DF) image (Fig. 3(c)) shows many cementite particles precipitate inside the ferrite lath. The cementite demonstrates two types of morphologies: rod-like and spherical particles. Combining the two-beam and DF image, we can see that the dislocation tangles are located near the cementite particles. Fig. 3(d) shows the kernel average misorientation (KAM) mappings of IA640-30 specimen before deformation. KAM is somewhat in proportional with the density of geometry necessary dislocations (GNDs) [15]. The KAM values are higher at the interface between lamellar austenite and ferrite. As shown in Fig. 3(e), the KAM values increases obviously after fracture. The deformation occurs at both the ferrite and austenite. Secondary deformation mechanism like deformation twinning can be observed at the austenite and the fraction of austenite decreases lower than 10%. The good match of continuous yielding and stable strain hardening helps improve the formability of the medium Mn steel. The initial high density of GNDs for hot-rolled medium Mn steel is considered as the results of the inactive recovery [16–18]. The unpinned GNDs in ferrite contribute to the continuous yielding, which is similar with the DP steel [19]. The PLC effect is related with the dynamic strain aging (DSA). All the specimens demonstrate the lamellar morphology of austenite and ferrite. The stress serration at the tensile curves disappears when the cementite precipitations exist. The cementite particles for IA640-30 and IA68010 specimens mainly precipitate from the ferrite and it will lower the C content in ferrite. Therefore, the ferrite grains with less solute carbon atoms will weaken the dynamic interaction with dislocations during strain hardening. Besides, the grain size of prior austenite is about 35 lm (Fig. S3b) and the coarse grains are reported to help remove the serrated flow [20,21]. 4. Conclusions A moderate PSE grades hot-rolled medium Mn steel without plastic instability phenomena is obtained by controlling the IA process. The microstructure demonstrates lamellar morphology of
austenite and ferrite after IA process. Cementite precipitations exist for IA640-30 and IA680-10 specimens. The initial high density of GNDs in ferrite contributes to the continuous yielding. Cementite precipitations lower the solute carbon content in ferrite. Together with the coarse grains, the dynamic interaction with dislocations is eliminated. Overall, the plastic instability phenomena, especially PLC effect, can be overcome by controlling the cementite precipitation for hot-rolled medium Mn steel. Declaration of Competing Interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Acknowledgements This work was supported by the Fundamental Research Funds for the Central Universities (No. FRF-TP-18-039A1); the Project funded by China Postdoctoral Science Foundation (No. 2019M650482). Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi.org/10.1016/j.matlet.2019.126796. References [1] S.P. Chen, R. Rana, A. Haldar, R.K. Ray, Prog. Mater. Sci. 89 (2017) 345–391. [2] H.W. Luo, H. Dong, M.X. Huang, Mater. Des. 83 (2015) 42–48. [3] H. Kamoutsi, E. Gioti, N. Gregory, Z. Cai, H. Ding, Metall. Mater. Trans. A 46 (2015) 4841–4846. [4] R. Schwab, V. Ruff, Acta Mater. 61 (2013) 1798–1808. [5] E. Emadoddin, A. Akbarzadeh, G.H. Daneshi, Mater. Sci. Eng. A 447 (2007) 174– 179. [6] H. Aboulfadl, J. Deges, P. Choi, D. Raabe, Acta Mater. 86 (2015) 34–42. [7] H.J. Pan, H. Ding, M.H. Cai, D. Kibaroglu, Y. Man, W.W. Song, Mater. Sci. Eng. A 766 (2019) 138371. [8] B.B. Straumal, Y.Q. Kucheev, L.I. Efron, A.L. Petelin, J.D. Majumdar, I. Manna, J. Mater. Eng. Perform. 21 (2012) 667–670. [9] X.G. Wang, L. Wang, M.X. Huang, Acta Mater. 124 (2017) 17–29. [10] R.K. Rai, J.K. Sahu, Mater. Lett. 210 (2018) 298–300. [11] M. Callahan, O. Hubert, F. Hild, A. Perlade, J.H. Schmitt, Mater. Sci. Eng. A 704 (2017) 391–400. [12] M.H. Zhang, L.F. Li, J. Ding, Q.B. Wu, Y.D. Wang, J. Almer, F.M. Guo, Y. Ren, Acta Mater. 141 (2017) 294–303. [13] B.H. Sun, N. Vanderesse, F. Fazeli, C. Scott, J.Q. Chen, P. Bocher, M. Jahazi, S. Yue, Scr. Mater. 133 (2017) 9–13. [14] F. Yang, H.W. Luo, E.X. Pu, S.L. Zhang, H. Dong, Int. J. Plast. 103 (2018) 188–202. [15] B. Hu, H.W. Luo, Acta Mater. 176 (2019) 250–263. [16] B.B. He, M. Wang, M.X. Huang, Metall. Mater. Trans. A 50 (2019) 2971–2977. [17] E. Welsch, D. Ponge, H. Hafez, S. Sandlobes, P. Choi, M. Herbig, S. Zaefferer, D. Raabe, Acta Mater. 116 (2016) 188–199. [18] J.H. Han, S.J. Lee, J.G. Jung, Y.K. Lee, Acta Mater. 78 (2014) 369–377. [19] M. Jafari, N. Saeidi, S. Ziaei-Rad, M. Jamshidian, H.S. Kim, Mech. Mater. 134 (2019) 132–142. [20] W.H. Wang, D. Wu, R.S. Chen, X.N. Zhang, J. Mater. Sci. Technol. 34 (2018) 1236–1242. [21] D. Wu, R.S. Chen, E.H. Han, Mater. Sci. Eng. A 532 (2012) 267–274.