Journal of Materials Science & Technology  2019 , 35 (8): 1779-1786 https://doi.org/10.1016/j.jmst.2019.04.007

Orginal Article

Deformation-induced martensitic transformation kinetics and correlative micromechanical behavior of medium-Mn transformation-induced plasticity steel

Minghe Zhang, Haiyang Chen, Youkang Wang, Shengjie Wang, Runguang Li, Shilei Li, Yan-Dong Wang*

Beijing Advanced Innovation Center for Materials Genome Engineering, State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing, 100083, China

Corresponding authors:   *Corresponding author.E-mail address: ydwang@mail.neu.edu.cn (Y.-D. Wang).

Received: 2018-09-18

Revised:  2018-11-13

Accepted:  2019-01-22

Online:  2019-08-05

Copyright:  2019 Editorial board of Journal of Materials Science & Technology Copyright reserved, Editorial board of Journal of Materials Science & Technology

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Abstract

An in situ high-energy X-ray diffraction (HE-XRD) technique was mainly used to investigate the micromechanical behavior of medium-Mn Fe-0.12C-10.16Mn-1.87Al (in wt%) transformation-induced plasticity (TRIP) steel subjected to intercritical annealing at 625 °C, 650 °C, 675 °C and 700 °C for 1 h. As the intercritical annealing temperature increased, the volume fraction of retained austenite (RA) and ultimate tensile stress (UTS) increased, while the Lüders strain and yield stress (YS) decreased. The incremental work-hardening exponent of experimental steel increased with increasing intercritical annealing temperature. The overall trend of the transformation kinetics of the RA with respect to the true strain followed the sigmoidal shape predicted by the Olson and Cohen (OC) model. Load partitioning occurred among the ferrite, austenite and martensite immediately after entering the yielding stage. Because the stability of the RA decreased with increasing intercritical annealing temperature, the load undertaken by the martensite increased. The moderate transformation kinetics of the RA and effective load partitioning among constituent phases were found to contribute to a favorable combination of strength and ductility for this medium-Mn TRIP steel.

Keywords: Medium-Mn TRIP steel ; High-energy X-ray diffraction ; Transformation kinetics ; Load partitioning

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Minghe Zhang, Haiyang Chen, Youkang Wang, Shengjie Wang, Runguang Li, Shilei Li, Yan-Dong Wang. Deformation-induced martensitic transformation kinetics and correlative micromechanical behavior of medium-Mn transformation-induced plasticity steel[J]. Journal of Materials Science & Technology, 2019, 35(8): 1779-1786 https://doi.org/10.1016/j.jmst.2019.04.007

1. Introduction

The automotive industry is under increasingly stringent vehicle-emissions standards from governmental regulations and high consumer demand for passenger safety. The development of 3rd generation of advanced high-strength steel (AHSS) with an excellent combination of strength and ductility, in addition to improvements in safety and crashworthiness qualities, has attracted great interest from academic and industrial institutions [[1], [2], [3]]. One of the most promising candidates is the medium-Mn transformation-induced plasticity (TRIP) steel with a Mn content between 5 and 10 wt%. This type of steel comprises of an ultra-fine-grained (UFG) mixture of ferrite (α), austenite (γ), and in some cases martensite (α') [[4], [5], [6]]. Such microstructures of medium-Mn TRIP steel are typically obtained from intercritical annealing performed after cold rolling. Areas previously examined in detail include the microstructural evolution and development of the mechanical properties as functions of the chemical composition, e.g., C, Mn and Al, as well as the intercritical annealing temperature and the intercritical annealing time [[6], [7], [8], [9], [10], [11], [12]]. Through optimization of the intercritical annealing conditions, e.g., the isothermal holding temperature and time, the amount and stability of retained austenite (RA) can be tailored to enhance the TRIP effect, allowing for a good combination of ultimate tensile strength (UTS), in the range of 800-1100 MPa, and a total elongation (TE) of about 25%-50% [[7], [8], [9], [10]].

The excellent combination of strength and ductility of medium-Mn TRIP steels is mainly ascribed to two effects, namely, their ultra-fine-grained (UFG) duplex microstructure and the activation of the TRIP effect during tensile deformation [[10], [11], [12], [13], [14], [15]]. The metastable RA with suitable stability transforms to martensite during plastic deformation, resulting in complex interactions among ferrite, austenite and martensite [13,14]. An in-depth understanding of the stress/strain partitioning among the constituent phases facilitates the development of the desirable combination of UTS and TE in medium-Mn TRIP steels [14]. The strain partitioning among constituent phases can be studied by in situ microscopic digital image correlation (DIC) [15], in situ tensile testing with electron back-scattered diffraction (EBSD) [16], and in situ high-energy X-ray diffraction (HE-XRD). In particular, the in situ HE-XRD technique can not only elucidate the stress/strain partitioning among constituent phases but also reveal the microstructure evolution during deformation [[17], [18], [19]]. However, the micromechanical behavior of medium-Mn TRIP steels with various austenite stability, particularly the correlation between the transformation kinetics of RA to martensite and the load partitioning among constituent phases, has not been systematically investigated.

In this study, the in situ synchrotron-based HE-XRD technique was used to monitor the microstructural evolution under the tensile deformation for medium-Mn TRIP steel with different microstructures. The transformation kinetics of the RA to martensite and the evolution of the lattice strains of the constituent phases were revealed as a function of strain. Our investigation focuses mainly on the relationship between the kinetics of deformation-induced martensitic transformation and load partitioning among constituent phases in a medium-Mn TRIP steel.

2. Materials and experimental procedures

2.1. Materials processing and microscopy observations

A medium-Mn steel with a chemical composition of Fe-0.12C-10.16Mn-1.87Al (in wt%) was fabricated in a vacuum furnace. After homogenization at 1200 °C for 1 h, 20 mm thick plates were hot-rolled between 1100 and 900 °C down to 4 mm thickness and subsequently cold-rolled to 1.5 mm thickness. The cold-rolled sheets were then intercritically annealed at temperatures of 625 °C, 650 °C, 675 °C and 700 °C for 1 h and cooled to room temperature in air. These specimens are designated as IA625, IA650, IA675 and IA700. The samples for scanning electron microscope (SEM) observation were first mechanically polished and then electrochemically polished in a mixed solution of 90% ethanol and 10% perchloric acid (vol.%); subsequently they were etched by 4 vol.% Nital solution. The quantitative analysis of the grain sizes of the ferrite and RA was conducted using SEM micrographs and Image-Pro Plus 6.0 (produced by Media Cybernetics company, USA) image analysis software [5]. The volume fraction of the RA was calculated using the integrated intensities of the α-200, α-211, γ-200, γ-220 and γ-311 diffraction peaks [20] and the following equation:

VA=1.4Iγ/(Iα+1.4Iγ) (1)

where Iγ is the integrated intensity of austenite and Iα is the integrated intensity of ferrite.

2.2. In situ HE-XRD experiments

The in situ HE-XRD experiments were performed at the 1-ID-E beamline at the Advanced Photon Source (APS) in Argonne National Laboratory (ANL); the experimental configuration is schematically shown in Fig. 1. A monochromatic X-ray beam (0.2 mm × 0.2 mm) with an energy of 71.676 keV (wavelength 0.017297 nm) was used to study the microstructural evolution and lattice strain change in the specimens during tensile deformation. Dog-bone-shaped tensile specimens with a gauge section of 10 mm (length) × 3 mm (width) × 0.5 mm (thickness) were mounted with the rolling direction parallel to the loading direction (LD). Tensile tests were performed at room temperature with a nominal strain rate of 1 × 10-3 s-1 using an MTS load frame. Peaks were fitted to pseudo-Voigt functions using MATLAB computer programs developed by the 1-ID-E beamline group. Using the peak position information from the 10° azimuth range of the Debye rings near the LD and transverse direction (TD), lattice strains were obtained by calculating the lattice spacing change relative to the lattice spacing of the undeformed state.

Fig. 1.   Experimental configuration of the tensile test with in situ synchrotron-based HE-XRD experiments.

3. Results and discussion

3.1. Microstructures and mechanical properties

Fig. 2(a)-(d) shows SEM micrographs of the IA625, IA650, IA675 and IA700 samples. The microstructural constituents consisted of ferrite and austenite, labeled as α and γ, respectively. The grains of the ferrite and austenite after intercritical annealing were mostly equiaxed, and distinct from the lamellar morphology of the ferrite and austenite found in the hot-rolled and intercritically annealed medium-Mn TRIP steel [12]. The ferrite grain sizes of the IA625, IA650, IA675 and IA700 samples were 0.38 μm, 0.43 μm, 0.47 μm and 0.48 μm, respectively. The austenite grain sizes of the IA625, IA650, IA675 and IA700 samples were 0.58 μm, 0.71 μm, 0.78 μm and 0.98 μm, respectively. The engineering stress-strain curves of the experimental steel after intercritical annealing at different temperatures are presented in Fig. 3. As the intercritical temperature increased, the yield strength (YS) decreased and the UTS increased. The TE decreased with increasing intercritical temperature after reaching a maximum value at 650 °C. Moreover, the Lüders strain decreased with increasing intercritical temperature and completely disappeared for the IA700 sample. The evolution of the Lüders strain as a function of the intercritical temperature observed in this investigation is in agreement with previous studies [10,21]. The best combination of UTS and TE was achieved by the IA650 sample, followed by the IA675, IA625 and IA700 samples; the detailed mechanical properties are listed in Table 1.

Fig. 2.   SEM micrographs of Fe-0.12C-10.16Mn-1.87Al steel subjected to intercritical annealing at different temperatures: (a) 625 °C, (b) 650 °C, (c) 675 °C and (d) 700 °C.

Fig. 3.   Engineering stress-strain curves of Fe-0.12C-10.16Mn-1.87Al steel tested during in situ HE-XRD experiments at room temperature.

Table 1   Mechanical properties of Fe-0.12C-10.16Mn-1.87Al steel subjected to intercritical annealing at different temperatures.

SpecimenYS (MPa)UTS (MPa)TE (%)UTS × TE (GPa%)
IA625
IA650
1058
964
1099
1163
26.6
44.8
29.2
52.1
IA675820133031.141.2
IA700394142018.225.8

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The work-hardening capability of TRIP steel is typically characterized by the incremental work-hardening exponent,

nincr=d ln σ/d ln ε (2)

where σ is true stress and ε is true strain. The true stress-strain curves of the studied steel after intercritical annealing at different temperatures are presented in Fig. 4(a). To investigate the work-hardening behavior of the studied steel, we plotted the incremental exponent nincr as a function of true strain in Fig. 4(b), in which the straight line corresponds to the instability criterion nincr=ε. In the case of the IA625 sample, the nincr oscillated around the nincr=ε-line, suggesting the formation of localized deformation throughout the entire strain region [21]. For intercritical annealing at 650 °C, nincr at the onset of straining was relatively low but steadily increased over a large strain range. Therefore, the highest TE was obtained in the IA650 sample. As the intercritical annealing temperature increased to 675 °C, nincr was higher at the onset of straining and increased faster relative to that of the IA650 sample. In the case of the IA700 sample, nincr first increased sharply and showed a relatively high maximum at low strains, followed by a sharp decrease during further deformation. Thus, the IA700 sample demonstrates the highest UTS and lowest TE among all samples. The overall trend of the nincr of the experimental steel was an increase with increasing intercritical temperature, and an increase (followed by a decrease) with increasing true strain.

Fig. 4.   (a) True stress-true strain curves of Fe-0.12C-10.16Mn-1.87Al steel; (b) Incremental work-hardening exponent nincr of Fe-0.12C-10.16Mn-1.87Al steel as a function of true strain.

3.2. Kinetics of the deformation-induced martensitic transformation

The deformation behavior of the samples after intercritical annealing at 625 °C, 650 °C, 675 °C and 700 °C for 1 h showed significant differences from sample to sample. This suggests that the austenite stability plays an important role in tailoring the mechanical properties. The diffraction patterns obtained from the GE3 detector of the undeformed and the fractured samples, and the calculated volume fraction of austenite during deformation, are shown in Fig. 5. The initial volume fraction of the RA increased with increasing intercritical annealing temperature, and the stability of the RA was strongly dependent on the intercritical annealing temperature. For the IA625 sample, the austenite was quite stable, and the volume fraction of RA decreased just slightly at an engineering strain of 0.05. As the engineering strain increased, another slight drop in volume fraction of RA was observed, as it transformed to martensite prior to sample failure. For the IA650 sample, the stability of metastable RA was lower in comparison with that of IA625, where the TRIP effect could be triggered by a proper external load. The volume fraction of RA decreased suddenly at an engineering strain of 0.05 after the sample yielded. With increasing plastic strain, the volume fraction of RA decreased step by step. According to previous research [13,22], after the Lüders band propagates through the entire gauge length of the sample, the stepwise RA transformation to martensite is caused by Portevin-Le Chatelier (PLC) band propagation within the studied steel. In the IA675 sample, the austenite stability was lower than that in the IA650 sample, and Lüders band propagation led to approximately 30% of the RA volume fraction to transform to martensite. Following Lüders band propagation, the RA transformed to martensite gradually until the sample failed. As the intercritical temperature increased further, IA700 sample, the stability of RA was even lower than that in the IA675 sample, and the deformation-induced martensitic transformation started at a relatively early stage of plastic deformation. Approximately 74% volume fraction of RA transformed to martensite prior to sample fracture at an engineering strain of 0.18.

Fig. 5.   (a) XRD patterns of the undeformed samples; (b) XRD patterns of the fractured samples; (c) evolution of the volume fraction of RA as a function of engineering strain for the Fe-0.12C-10.16Mn-1.87Al steel.

The kinetics of austenite to martensite transformation plays an important role in governing the mechanical properties of TRIP steel. The model proposed by Olson and Cohen (OC) is usually used to characterize the transformation kinetics of RA with respect to the true strain [23]:

f α'=1-exp[-β(1-exp(-αε))m] (3)

where fα' is the fraction of martensite transformed from RA, α and β are parameters used to characterize the stability of the RA during plastic deformation, where larger values imply lower stability [24], and m is a constant with a value of 2 for TRIP-assisted multiphase steel [25]. Fitting with Eq. (3) was conducted for the experimental data, and the results are shown in Fig. 6. The OC model may not be particularly appropriate to describe the burst transformation of austenite to martensite during Lüders band and PLC band propagation. However, the overall trend of the kinetics of austenite-to-martensite transformation in all specimens still follows the sigmoidal shape predicted by the OC model. The exact α and β parameters for the investigated steel are listed in Table 2. Although α for the IA625 sample is higher than those of the IA650 and IA675 samples, the value of β for the IA625 is the lowest among all samples. As the intercritical annealing temperature increases, the stability of RA decreases, as expected. In addition, the martensite formation rate, r, is dependent on the true strain and can be expressed as r(ε)=dfα'/dε [26]. The fitting data for fα' in the OC model was used to calculate the martensite formation rate, and the martensite transformation rate as a function of true strain for the studied steel is presented in Fig. 7. On the one hand, the martensite transformation rate increased as the intercritical annealing temperature increased. On the other hand, the martensite transformation rate increased and then decreased with increasing true strain. The martensite transformation rate with respect to true strain follows an inverted U-shaped correlation.

Fig. 6.   Normalized austenite transformed fraction of Fe-0.12C-10.16Mn-1.87Al steel during tensile deformation: (a) IA625 and IA650, (b) IA675 and IA700.

Table 2   Parameters α and β for Fe-0.12C-10.16Mn-1.87Al steel.

Sampleαβm
IA6258.851170.134252
IA6504.843231.181252
IA6757.591992.039292
IA70029.338651.753752

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Fig. 7.   Martensite transformation rate for Fe-0.12C-10.16Mn-1.87Al steel as a function of true strain.

3.3. Lattice strain of the constituent phases

The micromechanical behavior of the experimental steel can be understood from the lattice strains of the constituent phases obtained from the in situ HE-XRD experiments. The α-211 and γ-311 diffraction peaks were used to calculate the body-centered cubic (bcc) and face-centered cubic (fcc) lattice strains, respectively, as these measurements are insensitive to intergranular strain [27,28]. Synchrotron-based HE-XRD with high spatial resolution makes it possible to separate the overlapping ferrite and martensite peaks despite their similar crystal structures [29,30]. We fit the asymmetric α-211 diffraction peaks along the LD with two Gaussian functions characterized as the ferrite and martensite as shown in Fig. 8. For austenite and ferrite, the strain-free d-spacing could be obtained from the initial undeformed state. The d0 of the martensite is identified as the reference state by plotting the d-spacing vs. sin2Ψ, where Ψ is the azimuth angle relative to the LD and represented by cos Ψ=cos η cos θ [30,31].

Fig. 8.   The {211} diffraction peaks of IA675 sample (a) diffraction patterns of α-211 and α'-211 deformed at an engineering strain of 0.15; (b) diffraction patterns of α-211 and α'-211 deformed at an engineering strain of 0.24.

The elastic modulus and Poisson's ratio for the individual planes of the experimental steel were obtained from the elastic stage of the curve and are summarized in Table 3. The elastic moduli of α-211 and γ-311 are close to the values reported in previous studies [26,32]. As the intercritical annealing temperature increases, the elastic modulus for the individual planes decreases, while the Poisson's ratio increases. The lattice strains εhkl along the LD and TD for different {hkl} planes of the constituent phases of the studied steel are plotted as a function of true strain in Fig. 9. During the elastic stage, both ferrite and austenite are loaded elastically in all samples. The lattice strains of α-211 and γ-311 along the LD increase, while the lattice strains of α-211 and γ-311 along the TD decrease, consistent with the Poisson effect. As the true strain increased after the samples yielded, the difference between the lattice strains of the austenite and ferrite indicated that load partitioning is taking place. For the IA625 sample, once the Lüders band propagates at some point in the gauge length of the specimen, the lattice strain of α-211 dropped by approximately 500 με (where 1 με represents a strain of 10-6). Following Lüders straining, the lattice strains of α-211, γ-311 and α'-211 increased slightly until the sample failed. In the IA650 sample, Lüders band propagation led to a large drop of lattice strain of γ-311. The lattice strain of α-211 remained almost unchanged during the plastic stage, while the lattice strains of γ-311 and α'-211 increased gradually until the sample failed. For the IA675 sample, Lüders band propagation led to a larger decrease of the lattice strain of γ-311 due to a greater volume fraction of RA transformation to martensite, as shown in Fig. 5. Following Lüders straining, the lattice strains of γ-311 and α'-211 increased continuously until the specimen failed. As the intercritical annealing temperature increased to 700 °C, the lattice strain of γ-311 decreased slightly after sample yield and then increased continuously, while the lattice strains of α-211 and α'-211 increased remarkably until sample necking. For all samples, the strain hardening of ferrite, austenite and martensite after Lüders straining could be roughly estimated through ∂εLD/∂εtrue [33]. The deformation-induced martensite demonstrated the highest straining hardening capability, followed by the austenite and ferrite. With increasing intercritical annealing temperature, the strain hardening of the constituent phases increased. The significant differences in the evolution of lattice strains of α-211, γ-311 and α'-211 as a function of true strain is attributed to the different austenite stability in the studied steel.

Table 3   Elastic modulus (GPa) and Poisson's ratio for individual (hkl) planes.

SampleEα-211ν α-211Eγ-311ν γ-311
IA6251920.221970.25
IA6501860.241960.28
IA675
IA700
185
157
0.24
0.25
190
162
0.28
0.31

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Fig. 9.   Lattice strains for reflections from constituent phases along the LD and TD as a function of true strain: (a) IA625, (b) IA650, (c) IA675 and (d) IA700.

3.4. Load partitioning contributed by martensitic transformation

Load partitioning in multiphase steel has been previously investigated through the use of in situ HE-XRD and in situ neutron diffraction techniques [[31], [32], [33], [34], [35]]. The load partitioning can be quantified by calculating the von Mises stress, σeff, of the ferrite, austenite and deformation-induced martensite [31,36]:

σ11=$\frac{E}{1+v}ε_{11}+\frac{vE}{(1+v)(1-2v)}(ε_{11}+ε_{22}+ε_{33})$ (5)

σ22=$\frac{E}{1+v}ε_{22}+\frac{vE}{(1+v)(1+2v)}(ε_{11}+ε_{22}+ε_{33})$ (6)

where σ11 is the axial and σ22 (=σ33) is the transverse stress, ε11 and ε22 (ε33) are the axial and transverse strains, respectively, and E and ν are Young's modulus and Poisson's ratio, respectively. Since the martensite is transformed from austenite during tensile deformation, E and ν of α' martensite are not obtained in this study. As the experimental steel consisted of three constituent phases during deformation, the phase stress of the martensite can be calculated using the rule of mixtures, considering the volume fraction of the constituent phases [26,37]:

σ=fασα+fγσγ+fMσM (7)

where fα, fγ and fM are the volume fractions of the respective phases and σα, σγ and σM are stresses in the ferrite, austenite and martensite, respectively. The phase stress of the constituent phases for the experimental steel is plotted as a function of true strain in Fig. 10. In the elastic range, the phase stress of the ferrite and austenite increased linearly with increasing true strain. The load of each phase during the elastic deformation process was almost equal, indicating no substantial load partitioning between the different phases as a result of their similar elastic moduli [26]. As the true strain increased, yielding of the constituent phases occurred along with significant redistribution of the external load among constituent phases due to their different plastic properties. For the IA625 sample, during Lüders band propagation, there was a sudden increase in the phase stress of austenite and a decrease in the phase stress of ferrite. Due to the very small amount of austenite transformed to martensite, the phase stress of martensite was very low in comparison to that of the ferrite and austenite. For the IA650 sample, the evolution of phase stress of the ferrite and austenite was similar to that of the IA625 sample; a high-volume fraction of martensite was transformed from the austenite during plastic deformation, and the phase stress of the martensite increased from 161 MPa to 928 MPa. For the IA675 sample, a higher volume fraction of martensite was formed in comparison with that in the IA650 sample, and the phase stress of the martensite increased from 651 MPa to approximately 1387 MPa-even higher than the phase stress of the ferrite and austenite at a true strain beyond 0.18. For the IA700 sample, the phase stress of the ferrite, austenite and martensite increased rapidly during plastic deformation. Fig. 10 shows that the phase stress of the martensite increased as the intercritical annealing temperature increased. Moreover, as the intercritical annealing temperature increased, the volume fraction of deformation-induced martensite increased as shown in Fig. 5. Thus, the load undertaken by the martensite increased with increasing intercritical annealing temperature.

Fig. 10.   Phase stress of constituent phases as a function of true strain during deformation: (a) IA625, (b) IA650, (c) IA675 and (d) IA700.

The mechanical properties of TRIP-assisted multiphase steels, with deformation-induced martensitic transformation during plastic deformation, are significantly influenced by the strain/stress partitioning among constituent phases [[37], [38], [39]]. In the present study, the work-hardening capability of IA625 sample is very low because of the very low volume fraction of austenite transformed to martensite during deformation. The load is mainly undertaken by austenite, and the phase stress of the martensite is relatively low. In contrast, for the IA700 sample in our experiments, an approximate 74% volume fraction of RA transformed to martensite at an engineering strain of 0.18, as shown in Fig. 5, and the phase stress of the ferrite, austenite and martensite increased rapidly, leading to the highest UTS and lowest TE among all samples. A favorable combination of UTS and TE was achieved in the IA650 and IA675 samples. Two factors contributed to the desirable mechanical properties of these samples. On the one hand, a suitable stability of the RA was obtained in both samples, where the RA transformed to martensite moderately during deformation. On the other hand, effective load partitioning occurred among the ferrite, austenite and martensite. The martensite in our samples is capable of plastic deformation similar to that of high-dislocation-density martensite [40], as shown in Fig. 9. The strain-hardening capability of the martensite allows it to bear higher stress in comparison with the ferrite and austenite to accommodate the external load, facilitating effective load partitioning among constituent phases.

4. Conclusions

In this study, the micromechanical behavior of Fe-0.12C-10.16Mn-1.87Al medium-Mn TRIP steel fabricated by intercritical annealing at 625 °C, 650 °C, 675 °C and 700 °C for 1 h was systematically investigated via in situ synchrotron-based HE-XRD technique. The following conclusions can be drawn:

(1) As the intercritical annealing temperature increased, the volume fraction of austenite and UTS increased, while Lüders straining and YS decreased. The sample intercritically annealed at 650 °C demonstrated the largest TE and the optimal combination of UTS and TE. The incremental work-hardening exponent of the experimental steel increased with increasing intercritical temperature.

(2) The deformation-induced transformation kinetics from RA to martensite with respect to true strain followed the OC model. The martensite transformation rate of studied steel increased with increasing intercritical temperature. An inverted U-shaped relationship between the martensite transformation rate and true strain was observed.

(3) Load partitioning took place among the ferrite, austenite and martensite immediately after the studied steel yielded. Because the stability of the RA decreased as the intercritical annealing temperature increased, the load undertaken by the martensite increased.

Acknowledgments

This work was supported by the National Key Research and Development Program of China (No. 2017YFA0403804), the National Natural Science Foundation of China (NSFC) (Nos. 51471032 and 51527801), the Fundamental Research Funds for the Central Universities (Nos. 06111020 and 06111040) and the State Key Laboratory for Advanced Metals and Materials (Nos. 2016Z-01, 2016Z-12, and 2016Z-19). Minghe Zhang acknowledges financial support from the Chinese Scholarship Council (CSC). The authors thank Jun-Sang Park and Jonathan Almer, beamline scientists at APS, for their help during all stages of the experiment. The use of the Advanced Photon Source was supported by the US Department of Energy, Office of Science, Office of Basic Energy Sciences (No. DE-AC02-06CH11357).

The authors have declared that no competing interests exist.


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