Characterization of Microstructures and Mechanical Properties of Cold ...

ISIJ International, Vol. 55 (2015), ISIJ International, No. 10 Vol. 55 (2015), No. 10, pp. 2229–2236

Characterization of Microstructures and Mechanical Properties of Cold-rolled Medium-Mn Steels with Different Annealing Processes Jun HU,1,2)* Wenquan CAO,2) Chongxiang HUANG,3) Cunyu WANG,2,4) Han DONG2) and Jian LI1) 1) Center for Fuel Cell Innovation, State Key Laboratory of Material Processing and Die & Mould Technology, School of Materials Science and Engineering, Huazhong University of Science and Technology, Wuhan, 430074 China. 2) Special Steel Institute, Central Iron and Steel Research Institute, No. 76 XueYuanNanLu, Beijing, 100081 China. 3) School of Aeronautics and Astronautics, Sichuan University, Chengdu, 610065 China. 4) Taiyuan Iron and Steel (group) Co Ltd, Taiyuan, 030003 China. (Received on March 31, 2015; accepted on June 15, 2015)

The cold rolled medium-Mn steel was processed by two different heat treatments (intercritical annealing and quenching + intercrictical annealing). The microstructure evolution was characterized by scanning electron microscopy, electron back scattered diffraction and transmission electron microscope. It was found that austenite reverted transformation took place during annealing processes. Blocky ultrafine ferrite/austenite duplex structure was formed after intercritical annealing, while lath ultrafine duplex structure was developed after quenching + intercritical annealing. Such ultrafine duplex structure has both high strength and high plasticity. The excellent properties and tensile behavior were related to the phase transformation from retained austenite to martensite, which was associated with the carbon content, the morphology of the retained austenite and the matrix microstructure in the steel with different processes. KEY WORDS: cold rolled medium-Mn steel; intercritical annealing; microstructure evolution; mechanical properties; TRIP effect.

1.

tion (ART) annealing.9–13) It has been reported13) that ARTannealing of Fe-0.2C-5Mn steel resulted in the microstructure evolution from fully martensite structure to an ultrafine lamellar ferrite and austenite duplex microstructure and the significant difference of Mn concentrations between ferrite lath and austenite lath indicated a strong Mn partitioning during ART-annealing process. It has been verified that the TRIP effect depended on the volume fraction and stability of retained austenite.14,15) The stability of retained austenite depends on a number of factors such as chemical composition, morphology and size of retained austenite. It has been reported that the stability increased with carbon content but decreased with austenite grain size and the film-like austenite is more stable than the blocky one.16–19) Previous studies have been done for medium-Mn steel as the forging state7,8,13,27,31) and an ultrafine lamellar ferrite and austenite duplex microstructure was developed after quenching + ART-annealing. Meanwhile an ultrafine granular ferrite and austenite duplex structure was developed in cold rolled medium-Mn steel after ART-annealing.29) In this paper the cold rolled medium-Mn steel was processed by two different ART-annealing treatments and the microstructure evolution during ART-annealing was examined by scanning electron microscopy (SEM), electron back scattered diffraction (EBSD) and transmission electron microscope (TEM). The major work is to study the microstructure evolution during different ART-annealing processes and

Introduction

Due to the increasing demand for light weighting and safety requirements in automotive industry, more and more efforts have been paid for the development of inexpensive advanced high strength steel.1–3) The typical example is transformation induced plasticity (TRIP) steels.4) The tensile strength of TRIP-assisted steels used in the automobile industry has tended to be in the range of 700–900 MPa, with an elongation at failure of between 15% and 35%.5) So it was supposed to be the promising candidates for the application of automotive bodies by reducing the weight without loss of crash-worthiness because of their good combinations of strength and plasticity6) than other high strength steel such as dual-phase steel (DP steel) and interstitial-free steels (IF steel). Recently, the medium manganese steel was proposed as a typical representative of the third generation of automotive steel, which had high tensile strength (1–1.5 GPa) combined with high total elongation (31–44%) with large fraction retained austenite (35–45%).7,8) The excellent mechanical properties were attributed to ultrafine grained (UFG) ferrite and austenite duplex structure. Such UFG structure was developed by the method of austenite reverted transforma* Corresponding author: E-mail: [email protected] DOI: http://dx.doi.org/10.2355/isijinternational.ISIJINT-2015-187

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further understand the effect of microstructure on mechanical properties of the steel. 2.

XRD was installed with a Co target and a graphite crystal monochromator, and was operated with a tube current of 30 mA and a voltage of 30 kV. Diffraction patterns were obtained from 45° to 115°. The scanning step was 0.02° and integration time was 0.4 s. The volume fraction of retained austenite was calculated by the integral intensities of ferrite (200) and (211) and austenitic (200), (220) and (311) peaks by the following equation:22)

Experimental Procedure

2.1. Material Preparation The chemical composition of medium-Mn steel is C 0.11, Mn 5, P 0.008, S 0.002, N 0.003 (wt%), and other trace elements. The designed steel was melted in 50 kg vacuum induction furnace and casted into an ingot. The ingot was firstly forged into square slab with a thickness of about 30 mm, and then homogenized at 1 200°C for 2 h. Thereafter, the slabs were hot rolled to a thin sheet with 4 mm thickness in the temperature range of 1 200–900°C and air-cooled to room temperature (RT). For further cold rolling, a soft annealing at 650°C with 6 hours was conducted, which results in ferrite and austenite duplex structure. Finally the soft annealed steel was cold rolled to thickness of 1.8 mm.

Vi =

where Vi is the volume fraction of retained austenite, G is the parameter of crystal surface in austenite and martensite, Iα and Iγ are the integral intensities of the peaks in martensite and austenite, respectively. The carbon content of the austenite was calculated by the austenite lattice constant by:23) a0 = 3.555 + 0.044 x, ......................... (2) where a0 is the lattice constant of austenite determined by the above XRD measurement, x is the mass percentage of carbon in the retained austenite.

2.2. Heat Treatment The cold rolled steel was processed by two heat treatments as represented schematically in Fig. 1. One was ART annealing at 650°C for different times (1 minute, 5 minutes, 10 minutes, 1 hour, 6 hours, and 12 hours) and air cooling to RT, as shown in Fig. 1(a). The other one was first austenized at 800°C for 10 minutes followed by water quenching, then ART annealing at 650°C for different times (1 minute, 5 minutes, 10 minutes, 1 hour, 6 hours, and 12 hours) and air cooling to RT, as shown in Fig. 1(b). Here, it was named Quenching + ART treatment.

2.4. Tensile Test Tensile specimens with a gauge dimension of 50 × 10 × 1.5 mm3 were machined from the annealed steel with tensile axis parallel to prior rolling direction. The uniaxial tensile test was carried out at a starting strain rate of 2.5 × 10 − 4 s − 1 on a WE300B tensile testing machine at RT. Tensile tests were repeated twice for each heat treatment condition. 3.

2.3. Microstructural Characterization The microstructures were characterized by S-4300 cold field emission SEM and EBSD in a LEO 1530 field emission SEM. The raw EBSD data was processed by orientation averaging using the VMAP software20) to reduce orientation noise. After noise reduction, the angular resolution was improved to 1° in this study.21) Further analysis of these EBSD data was carried out by HKL software (Channel 5). TEM microstructures were examined by HITACHI H800 microscope, which were operated at 120 kV. The samples were mechanically ground to a thickness of about 40 μm, and then were electro-polished by a twin-jet machine in a solution of 6% perchloric acid and 94% alcohol at about − 20°C. The volume fraction of residual austenite was determined by XRD in a PHILIPS APD-10 X-ray diffractometer. For XRD measurement, the samples were mechanical polished and then electrolyzed at 5 V for 50 s in a 10% chromic acid electrolyte to remove the surface stress. The

Fig. 1.

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1 , .......................... (1) 1 + G( Iα / I γ )

Results

3.1. X-ray Results Figure 2(a) shows the volume fraction of retained austenite at different time, as measured by X-ray experiments. It can be seen that austenite was quickly developed in the cold rolled steel by ART annealing. Even after 1 minute annealing, the austenite volume fraction could be 8% and after 6 hours its volume fraction could be high up to ~16%. The amount of retained austenite increases with increasing annealing time. For short time annealing ( < 1 hr), the austenite volume fraction in samples treated by ART annealing is higher than those after Quenching + ART treatment. Both treatments have a peak value at annealing time of 6 hrs, 16.5% for ART annealing and 22.3% for another. Figure 2(b) shows the evolution of carbon content in retained austenite with annealing time. The carbon content was modified based on the results calculated from Eq. (2), as the Mn content must be taken into consideration.13) After a

Schematic representation of two heat treatments applied to cold-rolled sheet steels.

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fast increase within 5 minutes, the carbon content decreases slowly with increasing annealing time. The carbon contents in samples after ART annealing are always lower than that those treated by Quenching + ART treatment.

carbide precipitations in samples with a short time annealing, as seen the white dots in Figs. 3(a) and 3(c). These precipitations have been proved to be (Fe,Mn)3C in previous study.24) With increasing annealing time, the precipitation disappeared gradually and the grains coarsened slowly, as shown in Figs. 3(b) and 3(d). The typical EBSD microstructures are shown in Fig. 4. In these figures the retained austenite grains are red colored, while the ferrite/martensite are white colored. The high angle boundaries and low angle boundaries are depicted as black lines and green lines, respectively. The phase identification is closely related with the confidence index (i.e., difference between votes for the first and second solution divided by the total possible number of votes corresponding to the detected pattern).25) It can be seen with increasing annealing time the austenite fraction increasing gradually. It should be noted that the austenite amount measured by EBSD was a little different from that by XRD result, since

3.2. Microstructures Figure 3 shows the SEM microstructures for heat treated samples. It can be seen that the ART annealing results in equiaxed grain structure, while the Quenching + ART treatment leads to a lath-typed structure, as compared Figs. 3(a)– 3(b) with Figs. 3(c)–3(d). Both structures are composed of concave and convex grains. As reported in previous studies,9–13) the austenite reverted transformation took place during annealing in medium-Mn steel, which resulted in a mixture composed of slightly recovered ferrite grains and a few of austenite grains. Since the austenite was etched easily in comparison with ferrite, the concave grains were identified as austenite grains. It is also found that there are a few

Fig. 2. Variation of reverted austenite and its carbon content with annealing time: (a) volume fraction of retained austenite, (b) carbon content in austenite.

Fig. 3. SEM microstructures of samples treated by ART annealing at 650°C for 5 minutes (a) and 6 hour (b), and Quenching +ART annealing at 650°C for 5 minutes (c) and 6 hour (d).

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Fig. 4. The representative band contrast EBSD maps of the steel after ART annealing at 650°C for (a) 5 minutes and (b) 6 hours and after Quenching +ART annealing at 650°C for (c) 5 minutes and (d) 6 hours. In which low-angle boundaries (2°≤θ≤15°) and high-angle boundaries (θ >15°) are depicted by green lines and black lines, respectively. The red color areas represent the retained austenite.

some ultrafine austenite may not be indentified and reflected in the figures, as shown in Fig. 4(c). In the early stage of annealing, there were a lot of low angle boundaries in the grains. With increasing annealing time the low angle boundaries decreased and the grain size was coarsened gradually. Figure 5 shows the typical TEM microstructures for the samples treated with different time. For the ART annealing specimens, the microstructures reveal equiaxed grains, as shown in Figs. 5(a) and 5(b). The grains containing stacking faults and twins were austenite, as verified by the corresponding selected area diffraction (SAD) patterns in the insets of Figs. 5(a) and 5(b). With increasing annealing time, both austenite grains and ferrite grains coarsened slightly. The average austenite grain size is about 0.3 μm in short time annealing sample and about 1 μm with annealing time up to 6 hrs. For the Quenching + ART specimens, the microstructures are characterized by alternatively stacked dark/bright laths, as shown in Figs. 5(c) and 5(d). The bright laths are ferrite, while the dark are austenite, as confirmed by the SAD pattern in Fig. 5(c). The average thickness of ferrite and austenite were about 0.3 μm for the sample annealed for 6 hrs.

by the extensometer with strain gauge of 25 mm and it was taken down after the stress reach the maximum value, the non-uniform elongation part could not be determined and thus was not shown in the curves. It is seen that the samples treated by ART annealing exhibit distinct yield point elongation (YPE), while continuous yielding behavior is found in the specimens after Quenching + ART treatment. Sakuma et al.26) reported that the YPE in TRIP steel was related to a low density of dislocations in ferrite grains. Therefore it was supposed that the dislocation density was not so high that Lüders band was required for initiation of yielding. On the other hand, the continuous yielding behavior was considered to be attributed mainly to a high density of dislocations caused by martensite transformation during the water quenching process. However, due to the very complex microstructure of the cold rolled steel, there exist many other factors that may affect the tensile behavior such as the amount of austenite in the steel. For short time annealing (≤10 minutes) samples treated by ART annealing, the strain hardening rate is considerably low even after 20% strain (seeing Fig. 6(a)), possibly due to low austenite fraction in the steels. When annealing time is increased up to 1 hr, apparent strain hardening can be seen after initial 10% plastic strain which is apparently resulted from TRIP effect due to an increased austenite fraction. This phenomenon can be also found in the sample treated by Quenching + ART

3.3. Mechanical Properties The engineering stress-strain curves of the steels are shown in Fig. 6. The stress-strain curves were determined © 2015 ISIJ

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Fig. 5. Microstructure evolution of the cold rolled steel after ART annealing at 650°C for (a) 5 minutes and (b) 6 hours and after Quenching +ART annealing at 650°C for (c) 5 minutes and (d) 6 hours characterized by TEM.

Fig. 6. The engineering stress-strain curves of the cold rolled steel after annealing with different annealing time (a) ART annealing at 650°C and (b) Quenching +ART annealing at 650°C.

Quenching + ART treatment. Figure 7(c) shows the variation of tensile elongation with annealing time. The maximum values of total elongation and uniform elongation were obtained in the specimen treated by ART annealing for 1 hour, 42% and 34%, respectively. For the Quenching + ART treatment, they are 44% and 34%, respectively, when annealing time reached 6 hours. The product of tensile strength to total elongation was applied to evaluate the combined mechanical properties of medium-Mn steels, which indicates the forming ability and deformation energy absorbed in collision.27) It is seen in Fig. 7(d) that with increasing annealing time the product of Rm*A increases gradually. The maximum value (35.3 GPa%) appears at 1 hour for the sample treated by ART annealing, while it is 35.2 GPa% for the sample after Quenching + ART treatment for 6 hours. Further increasing

annealing, as shown in Fig. 6(b). With increasing austenite fraction, TRIP effect is promoted, as seen the higher strain hardening in samples with long annealing time. The average values of yield strength (Rp0.2, measured at 0.2% offset), tensile strength (Rm), total elongation (A), uniform elongation (Agt), yield ratio (r) and the production of tensile strength to total elongation (Rm*A) were plotted against annealing time in Fig. 7. It can be seen that the tensile strength changes slightly within 10 minutes, but increases gradually with increasing annealing time, as shown in Fig. 7(a). However, the yield strength decreases continuously with increasing annealing time. Accordingly, the yield ratio decreases rapidly when the annealing time is longer than 10 minutes, as seen in Fig. 7(b). Both tensile strength and yield strength of the specimens treated by ART annealing are higher than those of the specimens after 2233

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Fig. 7.

Mechanical properties of the cold rolled steel with the annealing time after different heat treatment (a) Rm and Rp0.2, (b) Yield ratio, r, (c) A and Agt and (d) Rm*A.

annealing time leads to a drop of the product. 4.

Discussions

4.1. Formation of Ultrafine Austenite Grains by Different Heat Treatments It reveals two distinct morphologies developed in the steels for two type treatments. One is ultrafine granular shaped austenite developed on the boundaries of the cold rolled ferrite subgrains, and another is lath typed austenite formed between martensitic laths. It can be concluded that the different grain morphologies were mainly controlled by their mother microstructure, such as the lath-typed austenite grains developed from the original martensitic lath but the granular austenite grains inherited from their granular cold rolled structure. For the ART annealing steel, the segregations of Mn and C onto the boundaries can account for the nucleation and growth of austenite phase, which are constrained by the diffusion of Mn onto the boundaries.24) The formation of ultrafine grains was attributed to the high density of dislocation and other defects introduced into the steel during cold rolling process.28) In other words, there was a large amount of deformation energy stored in the cold rolled steel and it would accelerate the phase transformation from ferrite to austenite during subsequent annealing process at the intercritical region. At the same time the ultrafine ferrite remained even after a long time ART-annealing since no recrystallization of the ferrite took place during the annealing process. For the Quenching + ART treated steel, it has been © 2015 ISIJ

Fig. 8. The total carbon content in the austenite with annealing time.

reported that austenite nucleated mainly between martensitic laths,13) and the growth behavior of austenite is carried out by the thickening of austenite lath, i.e., one dimension growth along thickness direction. Therefore, the thickness of both austenite lath and ferrite lath remains small ( < 0.5 μm) even after long time annealing at 650°C. The total carbon content in austenite which is the product of the volume fraction of the austenite (fA) to the carbon content in austenite (CA) as function of annealing time was plotted in Fig. 8. It can be seen that the product changes slightly with increasing annealing time for ART annealing sample, indicating a quick diffusion of carbon atom into austenite in the early stage of annealing. This is due to the high bulk diffusion coefficient of carbon atoms in steel. The 2234


Fig. 9.

(a) Relationship between elongation (A and Agt) and CA*fA, (b) Dependence of the product of Rm*A on austenite volume fraction.

sluggish coarsening of grains could be ascribed to the high Gibbs free energy in ferrite phase with high Mn content and the slow diffusion of Mn in both ferrite and austenite.29) However, for Quenching + ART treatment the carbon content in austenite increases linearly with increasing annealing time. It has been reported that the thickening of austenite lamellar was controlled by Mn diffusion in austenite.13) Due to the slow diffusion of Mn in austenite, the volume fraction of austenite increased gradually, which resulted in an increasing in total carbon content in austenite.

the steels by different heat treatments. There is obvious Lüders deformation at the begging of plastic deformation for ART samples (Fig. 6(a)). In comparison, it provides higher and continuous work hardening in Quenching + ART samples. Besides, serrated and/or jerk flow feature were observed from the stress-strain curves (Fig. 6(b)), indicating dynamic strain aging during tension test, which may be related to the ultrafine grain size in the ART samples. According to Edmonds et al.,35) for the TRIP steel film-like retained austenite is more effective than blocky retained austenite. Thus the martensite/austenite island should be avoided. In this study the blocky retained austenite was obtained in the ART steel but lath-like/film-like retained austenite in the Quenching + ART steel. Therefore, it can be inferred that the large blocky retained austenite was not as stable as lath-like retained austenite to facilitate phase transformation at small strain, providing early and strong TRIP effect and hence rapidly increase in initial strain hardening. The lath-like retained austenite can be transformed at higher strain to sustain TRIP effect.36)

4.2.

Dependence of Mechanical Properties on Microstructure It is well known that the TRIP effect is resulted from the strain-induced martensite transformation, which provides high work hardening rate and delays necking at higher strain. The austenite phase played an important role in phase transformation, depending on both the amount and stability of retained austenite. It is known that the stability of austenite was reduced by the decrease of carbon content in austenite and the increase of its grain size.27,30) Figure 9(a) shows the uniform and total elongation as a function of the product of austenite volume fraction (fA) and carbon content in austenite (CA). It can be seen that both uniform and total elongation increases with the product approximately linearly. Actually the uniform elongation shows clear dependence with the product value for the ART samples and Quenching + ART samples, respectively. The uniform elongation is not only related to austenite fraction and stability, but the morphology of the retained austenite. That’s why we see the changing trend of uniform elongation as function of the product of austenite fraction and stability is different between ART samples and the Quenching + ART samples as shown in the figure. Figure 9(b) shows the product of tensile strength to total elongation as a function of austenite fraction, which is compared with the typical 1st and 2nd generation automobile steels.31) It can be seen that all the Rm*A values nearly linearly increase with increasing austenite fraction. During tensile process, the austenite phase transformed into martensite, which gives relative high uniform elongation due to the TRIP effects.32–34) The higher austenite volume fraction, the larger the uniform elongation and the total elongation could be obtained. Interestingly, it reveals different stress-strain curves for

5.

Conclusions

In this study the cold-rolled medium-Mn steel was processed by two heat treatments and the microstructures and mechanical properties were investigated. The main conclusions are summarized as follows. (1) The ultrafine blocky ferrite with about 20% retained austenite was obtained in the cold rolled steel after ART annealing, while ultrafine lath ferrite/austenite duplex structure was developed during the annealing process after water quenching. The carbides precipitated from the martensite matrix at the earlier stage of annealing and gradually dissolved into the new developed austenite. (2) With increasing annealing time the tensile strength increased but the yield strength decreased, and both the elongation and the product of tensile strength to total elongation first increased then decreased for both steels. The maximum values (35.3 GPa% and 35.2 GPa%) were obtained at the annealing time of 1 hour for ART annealed steel and 6 hours for Quenching + ART treated steel, respectively. (3) The retained austenite played an important role in improving the mechanical properties due to the TRIP effect. The amount of retained austenite increased with the anneal2235

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ing time. And the stability of retained austenite was related to the carbon content and the morphology which resulted in the different tensile behaviors of the steels. The lathlike retained austenite was supposed be more stable which enabled sustained TRIP effects at subsequent higher strain.

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Acknowledgment This research is supported by the National Natural Science Foundation of China (51371057, 51101036, and 11172187) and National Basic Research Program of China (973 Program, No. 2010CB630803). REFERENCES 1) A. Grajcar, R. Kuziak and W. Zalecki: Arch. Civ. Mech. Eng., 2 (2012), 334. 2) R. Kuziak, R. Kawalla and S. Waengler: Arch. Civ. Mech. Eng., 8 (2008), 103. 3) C. Wang, H. Ding, M. H. Cai and B. Rolfe: Mater. Sci. Eng. A, 610 (2014), 436. 4) Z. Zhang, K. Manabe, Y. Li and F. X. Zhu: Steel Res. Int., 83 (2012), 645. 5) R. Zhu, S. Li, I. Karaman, R. Arroyave, T. Niendorf and H. J. Maier: Acta Mater., 60 (2012), 3022. 6) R. Zhu, S. Li, M. Song, I. Karaman and R. Arroyave: Mater. Sci. Eng. A, 569 (2013), 137. 7) J. Shi, X. J. Sun, M. Q. Wang, W. J. Hui, H. Dong and W. Q. Cao: Scr. Mater., 63 (2010), 815. 8) J. Shi, W. Q. Cao and H. Dong: Mater. Sci. Forum, 654–656 (2010), 238. 9) R. L. Miller: Metall. Trans., 3 (1972), 905. 10) K. T. Park, E. G. Lee and C. S. Lee: ISIJ Int., 47 (2007), 294. 11) N. Nakada, T. Tsuchiyama, S. Takaki and S. Hashizume: ISIJ Int., 47 (2007), 1527. 12) T. Hara, N. Maruyama, Y. Shinohara, H. Asahi, G. Shigesato, M. Sugiyama and T. Koseki: ISIJ Int., 49 (2009), 1792. 13) C. Wang, J. Shi, C. Y. Wang, W. J. Hui, M. Q. Wang, H. Dong and W. Q. Cao: ISIJ Int., 51 (2011), 651. 14) M. DeMeyer, D. Vanderschueren and B. C. De Cooman: ISIJ Int., 39

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