Carbon content-tuned martensite transformation in low-alloy TRIP steels

Ultrahigh strength and good ductility are obtained for two low-alloy transformation-induced-plasticity steels fabricated by the quenching and partitioning (Q&P) processing, respectively. Compared to 0.19 wt.% C steel in which γ → α′-martensite transformation is the dominant mechanism under deformation, the relatively high C content of austenite in 0.47 wt.% C steel is responsible for the transformation from γ to ε-martensite, suggesting that the transformation is not solely determined by the stacking fault energy. The study shows that during the Q&P process, strong and ductile steels with specific transformation procedures can be obtained by adjusting volume fraction and carbon content of the retained austenite.

Mechanical testing and microstructural characterization. Uniaxial tension along the rolling direction (RD) of the as-fabricated steels was performed at room temperature. The strain rate was 5 × 10 −3 s −1 . During loading strains were measured with a contactless laser extensometer. Samples used for the X-ray diffraction (XRD) analysis were etched with a solution of nitric acid (30 vol.%) and hydrochloric acid (70 vol.%). The XRD measurements were performed on an X'Pert diffractometer equipped with the Co K α1 anode. The carbon concentration of RA was determined with an assumption that the lattice parameters are affected by the chemical compositions 19 , using the method described in supplemental materials. Electron backscattering diffraction (EBSD) measurements were made for the rolled planes of the steels, using a SU-70 Hitachi field-emission scanning electron microscope (SEM) at an operating voltage of 20 kV. EBSD specimens were electro-polished with an electrolyte consisting of perchloric acid (10 vol.%), ethanol (80 vol.%) and deionized water (10 vol.%) for 20 s at 20 V. The scanning step was 0.1 µm and the EBSD data were analyzed using the OIM software. Deformation microstructures were further analyzed with a field-emission-gun Tecnai transmission electron microscope (TEM) operated at 200 kV. TEM specimens cut from the homogeneously deformed regions were prepared by mechanically polishing down to 50 µm thickness and punching into 3 mm diameter discs. Subsequently, twin-jet electropolishing at −10 °C was conducted using an electrolyte with perchloric acid (10 vol.%) and ethanol (90 vol.%) at 32 V.

Results and Discussion
Grain morphologies and phase distributions of the as-fabricated steels were characterized by EBSD (Fig. 1). For medium-carbon steel the average grain size is ~3 μm, which is slightly smaller than that in 0.19 wt.% C steel (~5 μm). The RA volume fraction is as high as 32 ± 2% in medium-carbon steel, which is significantly higher than that in low-carbon steel (~21 ± 2%). The XRD analysis show that the carbon concentration of RA in 0.47 wt.% C steel is 1.63 wt.% whilst that in 0.19 wt.% C steel is 1.17 wt.%. The intercritical temperature governs the Figure 1. Inverse-pole-figure (IPF) maps illustrating the morphologies on planes composed of normal direction (ND) and rolling direction (RD) for the as-fabricated 0.47 wt.% C steel (a) and the 0.19 wt.% C steel (b), respectively. The IPF maps show the grain orientations with respect to RD. The phase distribution maps corresponding to the IPF maps (a,b) are shown in (c,d), respectively, with blue for FCC phase (retained austenite), red for BCC phase (ferrite and α′-martensite) and yellow for HCP phase (ɛ-martensite).
concentration of alloying elements. Carbon element is the stabilizer of austenite and C atoms will diffuse from ambient ferrite to austenite during annealing. This leads to the high C concentration in austenite. Nevertheless, the initial carbon contents were different between the two steels. The larger C content facilitated the diffusion from ferrite to austenite, leading to the higher concentration of carbon in medium-carbon steel. On the other hand, medium-carbon steel undergoes the two-step Q&P process including the extra quenching at 250 °C for 8 s. Thus, the reverse phase transformation may be another responsible mechanism. Consistent observation has been reported in a 1.04 Cu-bearing steel undergone the three-step heat treatment, in which copper was enriched inside the austenite and a large copper content (15~19 wt.%) was found as the reverse transformation occurred 20 .
True stress-strain curves of the steels are presented in Fig. 2, in which a combination of high ultimate tensile strength and superior elongation can be seen for the both steels. Medium-carbon steel shows that the yield strength (σ y ) is 630 MPa, the ultimate tensile strength (σ UTS ) is 1680 MPa, and the elongation to failure is 0.38, respectively. Low-carbon steel shows the decreased σ y (520 MPa) and σ UTS (1100 MPa), but the increased maximum elongation of 0.50. The work-hardening curves in Fig. 2 show that work-hardening ability of the medium-carbon steel is significantly higher than that of the low-carbon steel.
Microstructures of the steels after tensile tests characterized by EBSD are shown in Fig. 3. No obvious variation in grain size appears in either steel compared to that prior to loading. However, significant differences can be identified between the phase distribution maps. For medium-carbon steel the BCC phases composed of ferrite and α′-martensite are dominant. Meanwhile, the HCP phase tangled with a few FCC phase is uniformly distributed around the BCC phases. The ɛ-martensite and the RA show the volume fractions of 30% and 3%, respectively ( Fig. 3c and Table 1). During tensile deformation, the transformed volume fraction of RA to ɛ-martensite is as high as 27% whilst that transformed to α′-martensite is only 2%. In contrast, in the deformed low-carbon steel the total volume fraction of BCC phases is as high as ~98%, without the occurrence of intermediate ε-martensite ( Fig. 3d and Table 1). This suggests that the majority of RA has been transformed to α′-martensite during tensile loading.
TEM observations show the ɛ-martensite dominated microstructures in medium-carbon steel after deformation (Fig. 4). The high density of dislocations pile up between the intervals of ε-plates ( Fig. 4a,b), suggesting that ε-martensite act as obstacles to slip. This is interesting because previous investigation on austenitic steels showed that during the transformation from austenite to α′-martensite, ε-martensite merely exists as an intermediate phase 21 . Here, streaks in the selected-area-electron-diffraction (SAED) patterns are related to the formation of HCP phase (Fig. 4c). The ε-martensite and ferrite coexist following a relationship of , namely the Kurdjumov-Sachs (K-S) misorientation relationship is fulfilled 22 . When the slip of stacking faults (SFs) is superposed on every second {111}γ, HCP structure can be formed. Therefore, it is conceivable that the irregular overlapping SFs gradually evolve into a regular sequence. High-resolution TEM (HRTEM) image for the region in Fig. 4a shows that the deformed microstructure consists of a bundle of thin ε-martensite plates. In contrast, low-carbon steel exhibits the different deformation behavior. Comparing the microstructures before and after tensile testing (Fig. 5a,b), one can see that RA (γ) has been transformed to α′-martensite under deformation. High density of dislocations appears in both ferrite and α′-martensite, contributing to the high ultimate tensile strength and good ductility of the bulk material. The SAED patterns of the γ phase (Fig. 5a) suggest an FCC lattice. Figure 5c,d reveal the SAED patterns of the α′-martensite and ferrite, taken along a zonal axis of [001] α′ and [111]α, respectively.
The stress-strain curves indicate that the σ UTS increases with increasing austenite from the steel with 0.19 wt.% C to that with 0.47 wt.% C. However, the total elongation of the steels goes along an inverse way. According to ref. 23 , martensitic transformation may lead to an increased work-hardening rate, delaying necking and thus effectively improves the ductility. As was demonstrated by our previous work 11 , the RA with the larger carbon concentration showed the elevated stability, enabling the slower but more gradual martensitic transformation during straining. Here the RA in medium-carbon steel is too stable to transform to α′-martensite because of its high carbon content. Both of the situations lead to the small increase in elongation 8 . Here, it is reasonable www.nature.com/scientificreports www.nature.com/scientificreports/ that as the carbon content in RA is moderate, low-carbon steel shows the larger elongation. This is related to the gradual phase transformation during the whole straining up to a true strain of 0.5, leading to the continuous strain-hardening behavior. It has been well established that uniform plastic deformation represents the resistance to necking for a material, i.e., a measure of the strain-hardening capability 24 . SEM images and elemental distribution maps detected by electron probe microanalysis (EPMA) of the steels can be found in the supplemental materials (Fig. S1). In the both steels RA is enriched in C, Al and Mn, illustrating that those elements are the stabilizers of austenite. However, the level of carbon content in medium-carbon steel is higher than that in low-carbon steel, despite of the identical levels of Al and Mn concentrations. Here we prove that RA is with the identical contents of Al and Mn but the significantly different carbon concentration between the steels. Therefore, the different phase transformation procedures should be mainly determined by the carbon concentration of RA.
To reveal the partitioning of alloying elements under intercritical annealing and bainitic holding, DICTRA simulations invoking thermodynamic database TCFE6 and mobility database MOBFE1 25,26 were performed. Details of the model establishment are given in the supplemental materials (Fig. S2). The objective of partitioning during Q&P is to obtain RA enriched with carbon. Austenite nucleates at an interface of martensite during the intercritical annealing, hence the alloy compositions of RA and martensite are the same as that at the initial stage of annealing. After the intercritical annealing of medium-carbon steel, in austenite C and Mn are enriched  www.nature.com/scientificreports www.nature.com/scientificreports/  www.nature.com/scientificreports www.nature.com/scientificreports/ whilst in ferrite Al and Si exhibit the higher concentrations, as shown in Fig. 6. For low-carbon steel annealed at the identical condition, the similar trends appear for C, Mn and Al, whereas the distribution of Si varies between the steels because of its different content in the steels. The objective of bainitic holding is also to obtain the carbon-enriched RA. The calculated carbon distribution during bainitic holding is presented in Fig. 7. The calculations show that the homogenization of carbon in RA is completed at around 100 s and 130 s in medium-carbon and low-carbon steels, leading to the different carbon contents, i.e., 0.68 wt.% and 0.58 wt.% in RA, respectively. Therefore, partitioning at 420 °C for 180 s is sufficient for the enrichment of RA with carbon in the studied steels.
Lü et al. 27 confirmed that no α′-martensite formed in a Fe-Mn-C alloy during cold rolling, even though the area fraction of ε-martensite increased from 2.7% to 33.1% as the strain increased from 22% to 51%. Generally, ε-martensite or twinning acts as an intermediate phase during the transformation from γ-austenite to α′-martensite 19,28,29 . The γ → ɛ → α′-martensite transformation occurs through a two-stage process as in advance of the final transformation to α′-martensite austenite transforms to ε-martensite first 12,[30][31][32] . Here, our observations for medium-carbon steel support the result in ref. 27 . Namely, the formation of ε-plates is dominant during the whole period of phase transformation. To examine the role of SFE, we calculate it following the procedure provided in supplemental materials. The SFEs in the both steels increase with the increasing temperatures as well as the distances from the center to the edge of austenite (see supplemental materials, Fig. S3). For medium-carbon steel at room temperature, the SFE increase from 21 mJ/m 2 to 32 mJ/m 2 with the increasing distances from the center of austenite. For low-carbon steel, the value increases from 14 mJ/m 2 to 23 mJ/m 2 with the increasing distances. It is noteworthy that the deformation-induced transformation in medium-carbon steel shows the high room-temperature SFE exceeding the generally reported limit for high-Mn steels (i.e., ~20 mJ/m 2 ), below which the formation of ε-martensite occurs 8,33,34 . According to Sato et al. 8 , in alloys with the SFE larger than 20 mJ/m 2 deformation twins, instead of ε-martensite, are produced under deformation. However, the current results show that in 0.47 wt.% C steel the dominant deformation mechanism is the transformation from γ to ε-martensite. Therefore, the ε-martensite formation results from the high stability of RA with the large C content. Several studies on TWIP steels also reported inconsistent relationships between SFE range and the corresponding deformation mechanisms 35,36 .
To this stage, the larger σ UTS of medium-carbon steel compared to low-carbon steel is attributed to the different deformation mechanisms of the steels. For medium-carbon steel, the γ → ε-martensite transformation significantly enhances work hardening at the expense of the maximum elongation. For low-carbon steel, however, www.nature.com/scientificreports www.nature.com/scientificreports/ the γ → α′-martensite transformation does not lead to an increased work hardening but contributes to a larger elongation. Phase transformation is determined by the stability of austenite. As previously reported that, the increasing austenite stability promotes the gradual transformation with straining and thus contributes to the high ductility 37,38 . Here, the ε-martensite acting as the final phase in the deformed medium-carbon steel is associated with the high C content and the high stability of austenite.

Conclusions
In this work, two strong and ductile low-alloy TRIP steels with different carbon contents are fabricated by the Q&P processing, respectively. Under plastic deformation the retained austenite in 0.47 wt.% C steel transforms into ε-martensite, while in 0.19 wt.% C steel without the formation of ε-martensite as an intermediate phase, the majority of austenite transforms into α′-martensite directly. Compared to the low-carbon steel, γ → ε-martensite transformatioin in the medium-carbon steel results from high stability of the retained austenite with larger C content. Therefore, through the different phase transformation procedures that are dependent on volume fraction and carbon content of austenite the optimal combination of strength and ductility can be obtained.