Effect of Thermomechanical Processing on the Microstructure and Mechanical Properties of a Medium Mn Steel Sheet with Nb Addition

Abstract

Medium Mn steels have gained significant attention for their potential to achieve exceptional combinations of strength and ductility through controlled microstructural engineering, and intercritical annealing is recognized as a key processing route to optimize austenite formation and mechanical properties in these alloys. However, the precise relationship between annealing parameters and resulting microstructural evolution in medium Mn steels remains inadequately understood, particularly regarding the optimal conditions for achieving superior mechanical performance. Therefore, this study systematically investigated the microstructural development of 7 wt.% Mn steel during controlled hot-rolling and intercritical annealing to establish processing-structure-property relationships. Hot-rolling produced predominantly lath martensite microstructure with minimal austenite presence. Intercritical annealing was conducted at 700°C and 780°C for 1, 2, and 5 h, resulting in ferrite/martensite and austenite microstructures with austenite content ranging from 10 to 25 wt.%. NbC precipitates were identified after electrochemical extraction. On the other hand, samples annealed at 700°C demonstrated optimal mechanical properties in terms of ultimate tensile strength and elongation, while both hot-rolled samples and those treated at 780°C exhibited brittle behavior with low elongation. These findings establish that intermediate annealing temperatures provide the most favorable microstructural balance for achieving superior mechanical performance in medium Mn steels.

Key words:
Medium Mn steel; Intercritical Annealing; Retained Austenite

1. Introduction

Medium Mn steels are categorized as third-generation advanced high-strength steels (AHSS)¹, characterized by an ultimate tensile strength (UTS) exceeding 1000 MPa and a ductility greater than 20%². These steels undergo initial processing through hot or cold rolling, followed by intercritical annealing (IA). This heat treatment produces a complex microstructure primarily composed of retained austenite (${\gamma }_R$), ferrite ($\alpha$), and martensite (${\alpha }^{\text{'}}$)³,⁴. The multiphase microstructure contributes to their exceptional mechanical properties, as it facilitates transformation-induced plasticity (TRIP) and/or twinning-induced plasticity (TWIP). Such effects emerge during plastic deformation, enabling excellent plasticity without compromising tensile strength and ductility⁵-⁷. Consequently, these attributes make medium Mn steels promising materials for the automotive industry, particularly for stamped structural parts⁸. The microstructure and desired mechanical properties are achieved through optimal selection and control of process parameters during hot-rolling and IA. Despite the balanced combination of strength and ductility, manufacturers continue to explore methods to further enhance the mechanical properties of medium Mn steels⁹-¹¹.

In the literature concerning other classes of steels, such as high-strength low-alloy steels (HSLA), the addition of microalloying elements like Nb or Ti has been shown to significantly refine grain structure. These elements raise the temperature at which austenite ($T_{NR}$) ceases to recrystallize. This increase benefits the final stages of hot-rolling by allowing deformation to accumulate in the austenite grains during the process. As a result, these elongated austenite grains, with their high dislocation density, enhance the nucleation rate of ferrite during IA. This process ultimately leads to considerable microstructural refinement¹². Recently, some studies in the literature have considered these findings in the context of medium Mn steel.

In a recent publication, Varanasi et al.¹³ observed that NbC decreased the mobility of the austenite-ferrite interface during IA, resulting in a lower austenite volume fraction in steel containing 0.06 wt% Nb and 0.06 wt% C. In another study, Mohrbacher et al.¹⁴ examined the precipitation of NbC at various processing stages in steel with 0.03 wt% Nb and 0.2 wt% C. They found that the volume fraction of NbC was lower during hot-rolling compared to IA. Additionally, they noted that NbC presence reduced ferrite recrystallization kinetics. Wu et al.¹⁵ reported that adding 0.04 wt% Nb and 0.1 wt% V decreased the recovery and recrystallization kinetics of martensite during IA. They proposed that dislocation hardening was the primary mechanism of hardening, while grain refinement played a secondary role.

In the references cited in the final paragraphs, cold rolling is most commonly the processing stage preceding IA¹³-¹⁵, with some instances of warm or hot-rolling¹⁰,¹¹. These processing conditions play a crucial role in determining the microstructural evolution of steel during IA. To the best of the authors’ knowledge, there is a paucity of research on the microstructural response of medium Mn steels microalloyed with niobium (Nb) during IA, particularly when hot-rolling precedes it. Therefore, this study aims to investigate the microstructural evolution and resultant mechanical properties of this steel when Nb is added.

2. Experimental

A medium Mn steel with a nominal composition of 0.15C-7Mn-1.5Al-2Si-0.1Nb-0.0035N-Fe (wt.%) was cast in a vacuum induction furnace. Sections measuring 70 × 40 mm with a thickness of 30 mm were cut from the cast ingots and heated to 1200°C for 120 min to homogenize the microstructure. Given that the ingots were solidified in permanent molds, it was anticipated that coarse Nb carbides would not precipitate during casting due to the high cooling rate. However, their formation is possible during the homogenization step and during hot-rolling (Figure 1). Subsequently, the sections were hot-rolled to a thickness of 3 mm, with a thickness reduction of 1 mm per rolling pass. A reheating step at 1000 °C was employed after every three rolling passes. The hot-rolling process was conducted using previously established parameters¹⁶. The thermomechanical process for the hot-rolling passes is illustrated in Figure 1.

Figure 1
Thermomechanical cycle used during the hot-rolling process.

The hot-rolling cycles were established to ensure that the steel temperature did not drop below 850 °C, preventing ferrite formation and excessive rolling load in the mill. The Continuous Cooling Transformation (CCT) diagram was calculated using JMatPro software Version 13.1, while the corresponding pseudo-binary phase diagram was generated with Thermo-Calc software and the TCFE9-Steel/Fe-Alloying v9.0 database (Figure 2). The CCT diagram helps explain phase formation during cooling from the final hot-rolling step. Meanwhile, the pseudo-binary phase diagram identifies the intercritical zone and aids in selecting annealing temperatures of 700 and 780 °C. The CCT diagram was constructed based on the overall steel chemistry and the amount of Nb in solid solution in austenite at 1200 °C, as indicated by thermodynamic simulations. Additionally, it considered an average grain size of 47 µm, which was measured in samples from the last rolling pass.

Figure 2
Calculated phase diagram by using Thermo-Calc software.

Three annealing holding times (1, 2, and 5 h) were used to observe changes in microstructure, with a particular focus on the austenite volume fraction. Thermo-CALC software was employed to determine the distribution of Nb among different phases in thermodynamic equilibrium at annealing temperatures. IA heat treatments were subsequently carried out using a Thermo-Scientific/FB1415M Muffle Furnace. Samples measuring 10 mm in width and 30 mm in length were cut from the as-rolled steel. These samples then underwent acid pickling to remove oxide scale before being intercritically annealed and finally quenched in brine.

Metallographic analyses were performed on both the as-rolled and annealed conditions using optical microscopy (OM) on an Olympus GX51 and scanning electron microscopy (SEM) on a JEOL JSM-6610LV operated at 30 kV. To reveal the steel microstructure, samples were color-etched in two steps: a 1–2 s pre-etch in Nital followed by a 20 s etch in sodium metabisulfite solution.

Phase identification for each condition employed a Bruker D8 Advance high-resolution X-ray diffractometer with Cu Kα radiation (λ = 0.15405 nm). Scans were collected from 40° to 110° 2θ with a 0.02° step size. Quantitative analysis was carried out with MAUD software (Rietveld refinement) using the PDF-2/Release 2010 database.

To verify the presence of Nb carbides, an electrochemical extraction was applied to a specimen annealed for 1 h at 700 °C—the condition that later proved to give the best mechanical properties. Powder for XRD and SEM was obtained by dissolving the matrix in an electrolyte of 10 vol% HCl and 90 vol% ethanol, following ASTM E963¹⁷. The resulting powders were filtered, dried, examined by SEM, and characterized by XRD under the same operating parameters, except for an extended scan range of 20–130° 2θ. NbC peaks were indexed using PDF card 00-038-1364 from the International Centre for Diffraction Data.

Tensile tests were conducted on machined flat specimens with a 50 mm gauge length following ASTM E8¹⁸. At least two specimens per condition were tested on a Shimadzu AGI-100 universal testing machine at a crosshead speed of 1 mm min⁻¹.

3. Results and Discussion

3.1. Effect of hot-rolling in the microstructure

Figure 3 illustrates the CCT diagram, simulated using JMATPro software, which considers the composition of the studied steel. The M_s temperature is 116°C, and the M_f temperature is -29.3°C. The surface temperature of the sheet post-hot-rolling was measured via a thermocouple at various time intervals, with an average cooling rate determined to be approximately 350°C/min. Based on this cooling rate, it can be concluded that the ferrite transformation curve is not intersected during cooling to room temperature, resulting in the transformation of nearly all austenite to martensite. Consequently, the primary phases in the microstructure are expected to be martensite and retained austenite.

Figure 3
Calculated CCT diagram using JMatPro.

Figure 4 displays the XRD pattern of the steel. After indexing the peaks, they were identified as being associated with martensite, with some low-intensity peaks attributed to austenite. The XRD pattern was refined using MAUD software, revealing a volume fraction of 92% martensite and 8% retained austenite. Micrographs in Figure 5 demonstrate a complete lath-like morphology corresponding to the martensite phase. The austenite is not visible in these micrographs due to its minimal presence and fine scale. The predominance of martensite is reported to be typical in medium Mn steel after hot-rolling, attributed to high hardenability from the alloying elements in such steels.

Figure 4
X-Ray diffraction pattern of the steel sheet in the hot-rolled condition.

Figure 5
SEM micrograph showing the martensitic microstructure of hot-rolled steel sheet.

Hu et al.¹⁹ reported that a steel with the composition 0.25C–7.17Mn–2.6Al (wt.%) formed a nearly fully martensitic microstructure with some equiaxed austenite grains after four hot-rolling passes, finishing at a temperature near 800 °C. These findings align closely with our results, as their austenite volume fraction is also low after this thermomechanical processing. In another study, Kim et al.²⁰ observed a fully lath martensitic microstructure in a steel composed of Fe-0.3C-5.9Mn-1.5Si-2.0Al-0.1Nb (wt.%) following hot-rolling. In contrast, when this steel underwent hot torsion testing, the resultant microstructure included 20% equiaxed ferrite, 3% retained austenite, and the remainder lath martensite. They attributed the presence of ferrite to the dynamic transformation of austenite to ferrite, promoted by the higher strain values encountered during torsion compared to conventional hot-rolling.

3.2. Microstructure evolution in the steel during intercritical annealing

Figure 6 illustrates the X-ray diffractogram of intercritically annealed samples treated for varying durations. As observed, the most prominent peaks correspond to the ferrite phase, while the smaller diffraction signals belong to the austenite phase. Importantly, distinguishing the individual contributions of ferrite and martensite to the overall peak intensity was not feasible, as both phases can diffract at identical positions. Consequently, the peaks were indexed as ferrite/martensite.

Figure 6
X-Ray diffractogram of the studied medium Mn steel at different IA conditions.

Additionally, the intensity of the austenite peaks increased significantly after IA at 700°C but decreased at 780°C. This behavior is attributed to the diminished stability of austenite at higher annealing temperatures, due to an increase in austenite volume fraction during isothermal holding and the associated reduction in Mn content within this phase, as previously documented²¹.

Furthermore, Table 1 provides the quantification of retained austenite and ferrite/martensite phases, determined using MAUD software. As noted, the maximum amount of retained austenite, approximately 26 wt.%, was achieved at 700°C, whereas less than 15 wt.% was retained at 780°C. For reference, intercritically annealed samples at 700 °C with holding times of 1, 2, and 5 h are labeled IA700-1, IA700-2, and IA700-5, respectively. Similarly, those at 780°C with holding times of 1, 2, and 5 h are labeled IA780-1, IA780-2, and IA780-5.

Table 1
Phases amount (wt.%) after intercritical annealing for the studied steel, measured by MAUD software.

	IA700-1	IA700-2	IA700-5	IA780-1	IA780-2	IA780-5
Austenite	23	26	26	10	14	12
Ferrite/Martensite	77	74	74	90	86	88
Sigma	1.6	1.4	1.7	1.1	1.3	1.5
Rwp	6.3	5.5	6.7	4.4	5.1	5.8

Correspondingly, Zou et al.²² reported approximately 22% retained austenite in medium Mn steel with a nominal composition of 0.005C-5.0Mn-0.5Si-1.4Ni-0.12V (wt.%), processed by hot-rolling in the non-recrystallization region, followed by annealing at 630°C for 5 h and subsequent air cooling to room temperature. In another study, Liu et al.²³ investigated medium Mn steel with a nominal composition of 0.12C-5.0Mn-1Al-0.2Mo-0.05Nb (wt.%), undergoing hot-rolling, IA, and tempering processes. The hot-rolled sheets were initially annealed at 600, 625, 650, and 690 °C for 30 min, then quenched and tempered at 200°C for 15 min, and finally cooled to room temperature. They achieved at least 19 wt.% of retained austenite at 600°C, with a peak of 29 wt.% at 650°C.

To clarify the impact of IA on the microstructure of steel, Figure 7 presents OM and SEM images of samples annealed at 700 °C. Figures 6a-f illustrate the microstructures of IA700-1, IA700-2, and IA700-5 samples, respectively. The microstructure comprises ferrite and austenite, which display two distinct morphologies: plate-like and equiaxed. The former appears white, while the latter is light and dark brown²⁴. Figure 8 showcases the OM and SEM images of IA780-1 (parts a, b), IA780-2 (parts c, d), and IA780-5 (parts e, f) samples, which exhibit morphologies similar to those of the IA700 samples. In the equiaxed region, ferrite (white) and austenite (brown) are present, whereas the plate-like area features a mix of blue and white, indicating the presence of martensite and ferrite, respectively. The formation of martensite in this area is linked to the decreased stability of austenite at the higher IA temperature.

Figure 7
OM and SEM images showing the ferrite/austenite phases formed after intercritical annealing at 700 °C with annealing times of: a) and b) 1h; c) and d) 2h; e) and f) 5h.

Figure 8
OM and SEM images showing the ferrite/martensite + austenite phases in the microstructure formed after intercritical annealing at 780 °C with annealing times of: a) and b) 1h; c) and d) 2h; e) and f) 5h.

When correlating these findings with the phase quantification provided in Table 1, it is evident that at 700°C, martensite does not form. The micrographs show no evidence of light or dark blue coloring, confirming that the phase reported as martensite/ferrite in the table consists entirely of ferrite. In contrast, the samples annealed at 780 °C exhibit blue regions within the microstructure, indicating martensite formation. Thus, the reported martensite/ferrite content in Table 1 is accurate, as distinguishing the individual contributions of these phases was not feasible.

Li et al.²⁵ proposed that the zone of phase nucleation is responsible for the observed morphologies. Growth of austenite and ferrite along the interface of the initial microstructure’s martensite laths results in a lath-like morphology. Conversely, equiaxed morphology arises when ferrite and austenite nucleation occurs at the intersection of martensite laths. They argued that as IA time increases, reduced interfacial energy causes the coarsening of the equiaxed zone and transforms the lath-like zone into an equiaxed one. Although this effect was not evident in IA-700 samples, slight coarsening of the equiaxed zone was observed in IA-780 microstructures, with no detected morphological change in the lath-like area. Li et al.²⁵ asserted that the two morphologies differ in stability because of varying dislocation densities, which are higher in the equiaxed zone, leading to lower stability in that area.

In contrast, Figure 9 demonstrates that thermodynamic simulations using Thermo-Calc software predicted that most Nb would exist as NbC in the microstructure below 1000 °C. Given the high dissolution temperature of this carbide (approximately 1300 °C), precipitation likely commenced during hot-rolling and may have completed during the IA treatment.

Figure 9
Calculated amounts of Nb in each phase, in equilibrium conditions for the given steel composition.

The results reported by Pan et al.²⁶ reinforce this argument. They found that the volume fraction of Nb precipitates increased during IA of a previously warm-rolled steel with a lower Nb content (0.05 wt.%) compared to our study. Additionally, the IA temperature they used, 650 °C, was similar to the temperature applied in our research.

Wu et al.¹⁵ observed a similar pattern of increasing precipitate volume fraction as the IA time extended from 3 to 30 min in a medium manganese steel composed of 5.2 wt.% Mn, 0.04 wt.% Nb, and 0.12 wt.% V. The IA temperature used was 660°C, closely comparable to the 700°C applied in our study. In a recent publication, Varanasi et al.¹³ utilized atom probe tomography to determine that the distribution of Nb within the microstructure of cold-rolled medium Mn steel was homogeneous before IA. They also noted that the precipitation of NbC was preceded by the formation of the carbide Fe_100–2xNb_xC_x during the early stages of IA.

Cuboidal particles with high Nb and C content were present at the prior austenite grain boundaries (Figure 10a and b). To verify whether these particles were NbC, the matrix of a selected sample (IA700-1) was dissolved using electrochemical procedures, and the particles were extracted. The XRD pattern of the extracted particles revealed peaks corresponding to NbC, along with some peaks associated with ferrite phase debris (Figure 10c). SEM micrographs of the extracted powder (Figure 9d) also showed NbC particles, most of which were smaller than 120 nm; however, a few isolated coarser particles measuring up to 500 nm were observed. The presence of NbC particles, particularly the finer particles, is expected to enhance the tensile properties of the steel sheet. Speer et al.²⁷ noted that the precise role of NbC in the microstructure of third-generation automotive steels remains as an active area of research.

Figure 10
a) SEM micrograph showing particles precipitated in the previous austenite grain boundaries in a IA700-1, b) EDS pattern showing the high content of Nb and C in the precipitate, c) X-Ray diffraction of the extracted particles showing the presence of NbC carbides mainly, d) SEM micrograph showing the extracted particles identified as NbC.

To monitor changes in the volume fraction of austenite during IA for all samples, we measured the average microhardness, as shown in Figure 11. The graph indicates that the highest microhardness, approximately 50 HRc, was observed in the HR sample, where a small amount of austenite was present, and martensite predominated. The IA-780 samples had an average microhardness of 46 HRc with an austenite content below 15%. In contrast, the IA-700 samples exhibited a microhardness of 32 HRc, with an austenite content close to 26%.

Figure 11
Average microhardness values for both hot-rolled and intercritical annealed samples.

3.3. Tensile mechanical properties

The tensile mechanical properties of hot-rolled and IA steel are depicted in Figure 12a. The 700°C treatment demonstrated superior mechanical properties compared to the 780°C treatment, particularly with respect to total elongation values. Specifically, the total elongation for the 700°C treatment was nearly 35%, while it was only 8% for the 780°C treatment. Within the 700°C IA steel samples, the optimal condition was achieved at a treatment duration of 1 h. At this duration, the yield strength approached 900 MPa, whereas it decreased to 800 MPa and 785 MPa after 2 and 5 h of IA, respectively. In this sample group, total elongation was consistent across all tests, ranging from 32% to 35%. For the UTS, values for 1 and 2 h were similar at approximately 1075 MPa. However, the IA700-5 sample showed a slightly higher UTS, close to 1100 MPa. This increase is attributed to a decrease in austenite content as the IA time exceeded a certain critical value, leading to increased ferrite/martensite formation and, consequently, a higher UTS, as previously reported²⁸.

Figure 12
Tensile test results from hot-rolled, a) IA-700 and b) IA-780 samples.

In the IA780 trials, as depicted in Figure 12b, all sheets demonstrated brittle performance. These samples closely resembled the hot-rolled sheet in terms of their low elongation (less than 8%) and high UTS, with values approaching 1500 MPa across all conditions. This UTS was 90 MPa lower than that of the hot-rolled condition. The decrease was attributed to the higher amount of martensite present in the hot-rolled condition.

Based on the aforementioned results, IA700 samples qualify as third-generation automotive steels, as evidenced by their product of UTS and total elongation (TE) (PSE) reaching 37 GPa%. This value exceeds the minimum requirement of 20 GPa% for this classification. In contrast, neither the IA-780°C group samples nor those in the hot-rolled condition meet the third-generation criteria, as their PSE values are approximately 12 GPa%, failing to achieve the necessary threshold.

4. Conclusions

In the present study, we investigated the experimental characterization of the microstructure of the 0.15C-7Mn-1.5Al-2Si-0.1Nb-Fe alloy following hot-rolling and IA. The main conclusions are as follows:

The microstructure of the steel post-hot-rolling consisted of lath martensite and a minor amount of retained austenite. The presence of these phases aligns well with calculations made using the determined CCT diagram. Additionally, the formation of the ferrite phase was ruled out since the experimentally measured cooling rate after hot-rolling did not intersect the ferrite formation curve.

The application of IA treatment at 700°C led to a beneficial microstructural change, as evidenced by the documentation of a high amount of austenite (26 wt.%). In contrast, treatment at 780°C resulted in the formation of a substantially smaller quantity of austenite (less than 15%). Both treatments yielded a microstructure characterized by a mix of lath and equiaxial morphologies.

Thermodynamic simulations calculate the precipitation of Nb carbides, and these determinations were confirmed through both SEM and XRD following the steel’s electrochemical extraction procedure. The carbides primarily resided at prior austenite grain boundaries, with an average size smaller than 120 nm, although some isolated coarser carbides were also observed.

The IA700 tests yielded the optimal combination of tensile mechanical properties, whereas the IA780 samples demonstrated properties closely resembling those in the hot-rolled condition, indicating brittle behavior. The superior performance of the IA700 group was attributed to the higher quantity of austenite and its increased stability acquired during IA. Additionally, the presence of Nb carbides likely facilitated precipitation strengthening and microstructural refinement, resulting in a yield strength exceeding 750 MPa across all IA700 samples.

5. Acknowledgments

The authors gratefully acknowledge the technical support from Alfredo Ruiz and Rosalina Tovar. Also, access to the research facilities from the Universidad Michoacana de San Nicolás de Hidalgo is gratefully appreciated.

• Data Availability
  Full dataset supporting the findings of this study is available upon request to the corresponding author.

6. References

• ¹ Kuziak R, Kawalla R, Waengler S. Advanced high-strength steels for automotive industry. Arch Civ Mech Eng. 2008;40(2):103-17. http://doi.org/10.1016/S1644-9665(12)60197-6
  » http://doi.org/10.1016/S1644-9665(12)60197-6
• ² Shi J, Sun X, Wang M, Hui W, Dong H, Cao W. Enhanced work-hardening behavior and mechanical properties in ultrafine-grained steels with large-fractioned metastable austenite. Scr Mater. 2010;63(8):815-8. http://doi.org/10.1016/j.scriptamat.2010.06.023
  » http://doi.org/10.1016/j.scriptamat.2010.06.023
• ³ Lee H, Jo MC, Sohn SS, Zargaran A, Ryu JH, Kim NJ, et al. Novel medium-Mn (austenite + martensite) duplex hot-rolled steel achieving 1.6 GPa strength with 20% ductility by Mn-segregation-induced TRIP mechanism. Acta Mater. 2018;147:247-60. http://doi.org/10.1016/j.actamat.2018.01.033
  » http://doi.org/10.1016/j.actamat.2018.01.033
• ⁴ Lee D, Kim JK, Lee S, Lee K, De Cooman BC. Microstructures and mechanical properties of Ti and Mo micro-alloyed medium Mn steel. Mater Sci Eng A. 2017;706:1-14. http://doi.org/10.1016/j.msea.2017.08.110
  » http://doi.org/10.1016/j.msea.2017.08.110
• ⁵ Bleck W, Guo X, Ma Y. The TRIP effect and its application in cold formable sheet steels. Steel Res Int. 2017;10(10):1700218. http://doi.org/10.1002/srin.201700218
  » http://doi.org/10.1002/srin.201700218
• ⁶ Field DM, Qing J, Van Aken DC. Chemistry and properties of medium-Mn two-stage TRIP steels. Metall Mater Trans, A Phys Metall Mater Sci. 2018;49(10):4615-32. http://doi.org/10.1007/s11661-018-4798-6
  » http://doi.org/10.1007/s11661-018-4798-6
• ⁷ Aydin H, Essadiqi E, Jung IH, Yue S. Development of 3rd generation AHSS with medium Mn content alloying compositions. Mater Sci Eng A. 2013;564:501-8. http://doi.org/10.1016/j.msea.2012.11.113
  » http://doi.org/10.1016/j.msea.2012.11.113
• ⁸ Li X, Chang Y, Wang C, Hu P, Dong H. Comparison of the hot-stamped boron-alloyed steel and the warm-stamped medium-Mn steel on microstructure and mechanical properties. Mater Sci Eng A. 2017;679:240-8. http://doi.org/10.1016/j.msea.2016.10.045
  » http://doi.org/10.1016/j.msea.2016.10.045
• ⁹ Hao L, Ji X, Zhang G, Zhao W, Sun M, Peng Y. Carbide precipitation behavior and mechanical properties of micro-alloyed medium Mn steel. J Mater Sci Technol. 2020;47:122-30. http://doi.org/10.1016/j.jmst.2019.12.024
  » http://doi.org/10.1016/j.jmst.2019.12.024
• ¹⁰ Pan H, Ding H, Cai M, Kibaroglu D, Ma Y, Song W. Precipitation behavior and austenite stability of Nb or Nb–Mo micro-alloyed warm-rolled medium-Mn steels. Mater Sci Eng A. 2019;766:138371. http://doi.org/10.1016/j.msea.2019.138371
  » http://doi.org/10.1016/j.msea.2019.138371
• ¹¹ Zhou N, Song R, Song R, Li X, Li J. Influence of Nb addition on microstructure and mechanical properties of medium-Mn low-density steels. Steel Res Int. 2018;89(8):1700552. http://doi.org/10.1002/srin.201700552
  » http://doi.org/10.1002/srin.201700552
• ¹² Hu B, Luo H, Yang F, Dong H. Recent progress in medium-Mn steels made with new designing strategies, a review. J Mater Sci Technol. 2017;33(12):1457-64. http://doi.org/10.1016/j.jmst.2017.06.017
  » http://doi.org/10.1016/j.jmst.2017.06.017
• ¹³ Varanasi RS, Gault B, Ponge D. Effect of Nb micro-alloying on austenite nucleation and growth in a medium manganese steel during intercritical annealing. Acta Mater. 2022;229:117786. http://doi.org/10.1016/j.actamat.2022.117786
  » http://doi.org/10.1016/j.actamat.2022.117786
• ¹⁴ Mohrbacher H, Yang JR, Chen YW, Rehrl J, Hebesberger T. Metallurgical effects of niobium in dual phase steel. Metals. 2020;10(4):504. http://doi.org/10.3390/met10040504
  » http://doi.org/10.3390/met10040504
• ¹⁵ Wu H, Guo H, Hu B, Enomoto M, Shang C. Ultra-high strength attributed to retardation of recrystallization during intercritical annealing in cold-rolled (V,Nb) microalloyed 5Mn steel. Mater Sci Eng A. 2020;794:139789. http://doi.org/10.1016/j.msea.2020.139789
  » http://doi.org/10.1016/j.msea.2020.139789
• ¹⁶ González-Castillo AC, Cruz-Rivera JJ, Ramos-Azpeitia MO, Garnica-González P, Garay-Reyes CG, Pacheco-Cedeño JS, et al. 3D-FEM simulation of hot rolling process and characterization of the resultant microstructure of a light-weight Mn steel. Crystals. 2021;11(5):569. http://doi.org/10.3390/cryst11050569
  » http://doi.org/10.3390/cryst11050569
• ¹⁷ ASTM: American Society for Testing and Materials. ASTM E963: Standard Practice for Electrolytic Extraction of Phases from Ni and Ni-Fe Base Superalloys Using a Hydrochloric-Methanol Electrolyte. West Conshohocken: ASTM; 2017.
• ¹⁸ ASTM: American Society for Testing and Materials. ASTM E8: Standard Test Methods for Tension Testing of Metallic Materials. West Conshohocken: ASTM; 2022.
• ¹⁹ Hu B, Luo H. Microstructures and mechanical properties of 7Mn steel manufactured by different rolling processes. Metals. 2017;7(11):464. http://doi.org/10.3390/met7110464
  » http://doi.org/10.3390/met7110464
• ²⁰ Kim W, Kang S, Jian B, Lee SW, De Cooman BC. High temperature deformation-induced transformation in Nb-bearing medium Mn steel. Metall Mater Trans, A Phys Metall Mater Sci. 2018;49(10):4389-93. http://doi.org/10.1007/s11661-018-4610-7
  » http://doi.org/10.1007/s11661-018-4610-7
• ²¹ Lee YK, Han J. Current opinion in medium manganese steel. Mater Sci Technol. 2015;31(7):843-56. http://doi.org/10.1179/1743284714Y.0000000722
  » http://doi.org/10.1179/1743284714Y.0000000722
• ²² Zou Y, Xu Y, Hu Z, Chen SQ, Han DT, Misra RDK, et al. High strength-toughness combination of a low-carbon medium-manganese steel plate with laminated microstructure and retained austenite. Mater Sci Eng A. 2017;707:270-9. http://doi.org/10.1016/j.msea.2017.09.059
  » http://doi.org/10.1016/j.msea.2017.09.059
• ²³ Liu C, Peng Q, Xue Z, Wang S, Yang C. Microstructure and mechanical properties of hot-rolled and cold-rolled medium-Mn TRIP steels. Materials. 2018;11(11):2242. http://doi.org/10.3390/ma11112242
  » http://doi.org/10.3390/ma11112242
• ²⁴ Zakerinia H, Kermanpur A, Najafizadeh A. Color metallography; a suitable method for characterization of martensite and bainite in multiphase steels. ISIJ Int. 2009;6:14-8.
• ²⁵ Li J, Song R, Li X, Zhou N, Song R. Microstructural evolution and tensile properties of 70 GPa·% grade strong and ductile hot-rolled 6Mn steel treated by intercritical annealing. Mater Sci Eng A. 2019;745:212-20. http://doi.org/10.1016/j.msea.2018.12.110
  » http://doi.org/10.1016/j.msea.2018.12.110
• ²⁶ Pan H, Ding H, Cai M. Microstructural evolution and precipitation behavior of the warm-rolled medium Mn steels containing Nb or Nb-Mo during intercritical annealing. Mater Sci Eng A. 2018;736:375-82. http://doi.org/10.1016/j.msea.2018.08.055
  » http://doi.org/10.1016/j.msea.2018.08.055
• ²⁷ Speer JG, Araujo AL, Matlock DK, de Moor E. Nb-microalloying in next-generation flat-rolled steels: an overview. Mater Sci Forum. 2017;879:1834-40. http://doi.org/10.4028/www.scientific.net/MSF.879.1834
  » http://doi.org/10.4028/www.scientific.net/MSF.879.1834
• ²⁸ Kaar S, Steineder K, Schneider R, Krizan D, Sommitsch C. New Ms-formula for exact microstructural prediction of modern 3rd generation AHSS chemistries. Scr Mater. 2021;200:113923. http://doi.org/10.1016/j.scriptamat.2021.113923
  » http://doi.org/10.1016/j.scriptamat.2021.113923

Edited by

• Associate Editor:
  Hugo Sandim.

• Editor-in-Chief:
  Luiz Antonio Pessan.

Data availability

Full dataset supporting the findings of this study is available upon request to the corresponding author.

Publication Dates

• Publication in this collection
  03 Oct 2025
• Date of issue
  2025

History

• Received
  27 Feb 2025
• Reviewed
  05 Aug 2025
• Accepted
  14 Aug 2025