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Article

Effect of Heat Treatment Temperature on the Microstructure and Mechanical Properties of Fe-18Mn-0.6C-xAl

by
Li Xiao
1,2,
Yuqi Zhang
1,2,
Huan Huang
1,2,
Bochao Zhang
1,2,
Ningning Ji
1,2,
Shuang Li
3 and
Jun Chen
3,*
1
CNOOC Key Laboratory of Liquefied Natural Gas and Low-Carbon Technology, Beijing 100028, China
2
CNOOC Gas & Power Group, Beijing 100028, China
3
State Key Laboratory of Digital Steel, Northeastern University, Shenyang 110819, China
*
Author to whom correspondence should be addressed.
Metals 2025, 15(8), 927; https://doi.org/10.3390/met15080927 (registering DOI)
Submission received: 21 May 2025 / Revised: 23 June 2025 / Accepted: 25 July 2025 / Published: 21 August 2025
Browse Figures
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Abstract

High-Mn steels are commonly fabricated by hot rolling and on-line cooling for cryogenic applications, because there exists an aging embrittlement zone in most high-Mn steels, and this shortcoming makes it difficult to optimize their mechanical properties by heat treatments. Hence, 0.6C-18Mn-0/3/5Al (in wt.%) steels were designed to investigate the effects of Al on their strength and toughness. The addition of 5 wt.% Al can increase yield strength from 357 to 461 MPa and the Charpy impact absorbed energy from 56 to 119 J. Although there is still a cryogenic aging embrittlement zone in each steel, we found that the addition of Al can narrow this brittle zone. Moreover, the absorbed energy is lowered by around 89%, 48%, and 40% for the 0Al, 3Al, and 5Al steels at −196 °C, respectively. Additionally, impact plastic deformation mechanisms were also revealed in the steels with a heat-treating temperature of 600 °C, revealing that the main deformation mechanism shifts from numerous partial dislocation slip to twinning plus strong planar slip as the addition of Al increases.

1. Introduction

High-Mn twinning-induced plasticity (TWIP) steels have been extensively investigated in light of their applications in the automotive industry, owing to their combination of strength and plasticity [1,2,3]. In fact, besides their strength and plasticity, their cryogenic impact toughness was also investigated back in 1985 [4] and was found to be better than that of 9Ni steels. Since 2012, high-manganese austenitic steel has become an attractive option again for cryogenic applications [5]. Numerous reports have focused on alloy design, microstructure, mechanical properties, strengthening and toughening mechanisms, and fabrication processes [5,6,7,8,9,10]. However, a few studies have reported the effect of heat treatment temperature on the evolution of the microstructure and mechanical properties. Moreover, heat treatment is not permitted under the ASTM standard [11].
One reason for this could be that high-Mn austenitic steels become brittle at certain temperatures, as supported by our previous studies [12], where numerous intergranular (Cr,Mn)23C6 carbides were observed in the steel after aging at 800 °C for 5 h, and Cr and C peaks were also detected at grain boundaries. Huang et al. [13] proposed that the susceptibility to fracture of Fe-Mn-C TWIP steel increases with Mn content. They attributed this phenomenon to the segregation of Mn at grain boundaries. Lee et al. [14,15] also indicated that the precipitation of M23C6 strongly deteriorates ductility and impact toughness. Moreover, this brittle temperature range is typically wide, which strongly suppresses the improvement of mechanical properties by heat treatment. Although heat treatment is usually applied to cold-rolled high-Mn steels to optimize grain size, particle size, and texture, the issue of heat-treatment-induced embrittlement has not attracted sufficient attention owing to its weak effect on tensile elongation at room temperature (RT). However, heat-treatment-induced embrittlement is very pronounced during cryogenic impact toughness tests. Hence, it is important to reduce or alter the brittle temperature range to optimize the microstructure by heat treatment.
In the present work, three Al-bearing Fe-18Mn-0.6C high-Mn steels were fabricated. The evolution of their microstructures with heat treatment temperature was analyzed in detail. We found that there is still a cryogenic impact embrittlement zone in each steel. In addition, the effect of heat treatment temperature on the degree of plastic deformation was also investigated.

2. Experimental Procedure

2.1. Materials

Three high-Mn steels with different amounts of Al were prepared, and their specific chemical compositions are shown in Table 1. Moreover, the stacking fault energies (SFEs) of the three steels were also estimated according to a thermodynamic model [16]. Ingots were first prepared through vacuum induction melting and then reheated at 1200 °C for 3 h. Subsequently, the ingots were hot-rolled to a thickness of 12mm using 7 passes at a temperature range of 1140–950 °C and then water-quenched to RT. The steels were further isothermally treated at 150, 300, 500, 600, 700, 800, 900, and 1000 °C for 5 h and then water-quenched to RT.

2.2. Room Temperature Tensile and Cryogenic Impact Tests

Standard cylindrical tensile samples with a gauge length of 50 mm and a gauge diameter of 6 mm were machined along the rolling direction (RD). They were tested at RT at a constant crossbar speed of 5 mm/min using an AG-X Plus PC-controlled tensile machine (Shimadzu Co., Ltd., Kyoto, Japan). Standard Charpy V-notch impact samples with gauge dimensions of 10 × 10 × 55 mm3 were also machined along the RD. They were held in liquid nitrogen for 20 min first and then tested on a pendulum impact tester (MTS systems Co., Ltd., Eden Prairie, MN, USA) with an impact hammer energy of 450 J and an impact speed of 5.24 m/s.

2.3. Microstructure Characterizations

The metallographic specimens were prepared from the hot-rolled and heat-treated plates, mechanically polished, and then further electrolytically polished in a solution of 10% perchloric acid and 90% ethanol to remove the strained layer. The fractured impact specimens were split at the center of their thickness. The cross-sectional regions close to the fracture surface were mechanically polished first and then further polished by a three-ion-beam polishing instrument (Leica EM TIC 3X, Wetzlar, Germany) to remove the strained surface layer. They were examined using a Zeiss Ultra 55 scanning electron microscope (SEM) (Zeiss, Oberkochen, Germany) to obtain electron back-scattered diffraction (EBSD)–inverse pole figure (IPF) maps, and the data were post-processed using HKL CHANNEL 5 software. In addition, the failure surfaces were examined using an FEI QUANTA 600 tungsten filament SEM (Thermo-Fisher, Waltham, MA, USA) to determine fracture modes. In order to perform transmission electron microscopy (TEM) observations, some specimens were prepared from cross-sections close to the failure surface; these were mechanically thinned to around 50 μm and then further thinned using a twin-jet electrolytic polisher (Struers TenuPol-5, Ballerup, Denmark). They were examined in a field emission transmission electron microscope operated at 200 kV (FEI Tecnai G2 F20) (Thermo-Fisher, Waltham, MA, USA).

3. Results and Discussion

3.1. Tensile and Cryogenic Impact Properties

The tensile engineering stress–strain curves are shown in Figure 1, and mechanical properties including yield strength (YS), ultimate tensile strength (UTS), and total elongation (TEL) are also listed in the figure. All the steels exhibit continuous yielding behavior, and the tensile properties are virtually unaffected by heat treatment at a temperature range of 150–800 °C. However, there is a steep decrease in YS and UTS after heat treatment at 1000 °C due to the large grain size of the steels. In addition, pronounced serrations are found in the stress–strain curves of the 0Al steels but are absent from the curves of Al-bearing steels. In general, the appearance of serrations is interpreted as dynamic strain aging (DSA) produced by the interaction between the Mn-C pair and the twinned partial dislocations [17,18]; in addition, Al can inhibit the formation of twinning partial dislocations and reduce the activity and diffusion rate of C atoms [19]. Hwang [20] indicated that the serrated flow in the tensile curve disappears when the ratio of Al to C is beyond about 3.0. This ratio in the 3Al and 5Al steels is far higher than 3.0, and dynamic strain aging is thus restrained.
Figure 2 shows the energy (−196 °C) absorbed by the tested steels at different heat treatment temperatures. The three steels exhibit similar trends in their absorbed energy. There is nearly no change in the absorbed energy of the 0Al steel below 400 °C, whereas the absorbed energy begins to drop sharply when the heat treatment temperature is greater than 400 °C, down to a minimum value at 600 °C, and then gradually increases with further increasing temperature, thus forming a cryogenic impact brittle zone. Although this brittle zone is still present in the 3Al and 5Al steels, the onset of the brittle zone shifted to 500 °C, thus narrowing the brittle zone. In addition, the absorbed energy of the 0Al steel after heat treatment at 600 °C is only 6 J, which is 89% lower compared to that of the hot-rolled 0Al steel, whereas those of the 3Al and 5Al steels are 48% and 40% lower, respectively. In addition, this aging embrittlement was also observed in Fe-Mn-Cr-Cu high-Mn steels, in which the peak temperature of aging embrittlement is approximately 800 °C [13]. Hence, it can be deduced that the addition of Al can lower the peak temperature of aging embrittlement.

3.2. Microstructure Characteristics Before Deformation

Figure 3 shows the EBSD-IPF maps of the hot-rolled and heat-treated 0Al, 3Al, and 5Al steels. The grains in the hot-rolled steels are in a recrystallized state, indicating that sufficient recrystallization occurred during hot rolling. The grain boundaries are also relatively straight, showing relatively low energy levels. There are no obvious color contrast changes within the same grain, also indicating that the grains are mainly recrystallized ones. Moreover, the orientation of the grains is random. In addition, Figure 3a,g,m show that grain size can be refined by the addition of Al.
It can be clearly seen that there are no pronounced changes in grain morphology and size after heat treatment at 300–800 °C; this may also be a major reason that there are no obvious changes in the tensile properties at RT. However, all the steels exhibit a large grain size after heat treatment at 1000 °C. Interestingly, although the grain of the hot-rolled steels can be refined by the addition of Al, the grain refinement effect of Al seems to be lost after heat treatment at 1000 °C, showing that the grain size increases with increasing Al content.
In addition, in order to examine all the grains in the steels, the grain size was estimated according to an area method. When the heat treatment temperature is 800 °C or below, the average grain size ranges of the 0Al, 3Al, and 5Al steels are 8.0–9.3, 6.3–7.3, and 6.0–6.6 μm, respectively. However, their grain sizes sharply increase to 32.6, 38.6, and 43.1 μm after heat treatment at 1000 °C.

3.3. Microstructure Characteristics After Cryogenic Impact Deformation

3.3.1. Impact Fracture Morphology

Typical SEM fractographs of the fractured impact samples are shown in Figure 4. The hot-rolled 0Al, 3Al, and 5Al steels show a typical ductile failure mode. The fracture surfaces are mainly covered by dimples, but the features of these dimples are different. The 3Al and 5Al steels exhibit similar morphology, whereas some large and deep dimples as well as severe tore regions can be observed in the 3Al steel. The dimple size of the 0Al steel is clearly larger than that of the 3Al and 5Al steels, while some flat-bottomed areas can be observed, indicating small plastic deformation. Thus, the 0Al steel has a lower amount of absorbed energy.
Compared with the hot-rolled steels, there is a large change in failure morphology after heat treatment at 600 °C. The fracture surfaces of the 3Al and 5Al steels still show similar morphology and are mainly covered by very small and shallow dimples. Combined with a change in absorbed energy, it can be deduced that the formation of large dimples in high-Mn austenitic steels always implies better cryogenic impact toughness. Of course, this is also associated with dimple depth and bottom morphology. The 0Al steel exhibits a typical brittle failure mode; cleavage-like fractures—caused by crack propagation along {111} planes in austenite steel at low temperatures and high strain rates [21,22,23]—and numerous secondary cracks can be observed. When the heat treatment temperature is as high as 1000 °C, numerous large dimples can be observed, indicating good cryogenic impact toughness. Although the depths of the dimples are relatively small, their bottoms are rugged, indicating large plastic deformation. In addition, the 3Al and 5Al steels show tore morphology, indicating that they have a higher amount of absorbed energy.

3.3.2. EBSD Observations of Cross-Sectional Regions Close to Fracture Surfaces

The deformation regions adjacent to main cracks were further analyzed through EBSD, as shown in Figure 5, whose insets are kernel average misorientation (KAM) maps. The high local misorientation indicates a high dislocation density, a wide distribution of misorientation, and a high accumulation of plastic deformation [24]. Figure 5a,d,g clearly exhibit that plastic deformation can occur far away from a main crack, and the KAM values are mainly over 2°. After heat treatment at 600 °C, there is a sharp decrease in plastic capacity, indicating that the plastic deformation regions became small. In particular, the KAM values for the 0Al steel in most of the observed regions are very close to 0°. Moreover, the main crack propagation path is relatively straight.
Interestingly, although Figure 2 shows that all the steels heat-treated at 1000 °C possess a relatively higher amount of absorbed energy compared to their hot-rolled counterparts, the severe plastic deformation regions are not significantly large, except for that for the 3Al steel. Figure 5c,f,i show that only a few grains, traversed by a main crack, can be observed. Hence, it is possible that the observed grains have unfavorable deformation orientations, resulting in large regions with small KAM values.

3.3.3. TEM Observations of Cross-Sectional Regions Close to Fracture Surfaces

In addition to the major deformation mechanism of dislocation slip, mechanical twinning and martensite transformation acting as secondary deformation mechanisms play a crucial role in the strain hardening rate, thereby leading to higher strength and ductility. Moreover, the secondary deformation mechanisms of high-manganese steels are predominantly governed by their SFE [25]. An SFE below 18 mJ/m2 results in ε/α′-martensite transformation during deformation. When the SFE ranges between 18 and 45 mJ/m2, deformation twinning occurs. An SFE above 45 mJ/m2 promotes dislocation slip. However, the deformation mechanisms can also be affected by deformation methods such as cryogenic impact deformation.
Figure 6 shows typical deformation microstructure characteristics near the fracture surfaces of the 0Al steel impact samples heat-treated at 600 °C, showing numerous stacking faults and a few dislocations. The stacking fault energy (SFE) of the 0Al steel is only 10.7 mJ/m2 at −196 °C. Although the SFE is relatively small, martensite is not observed. On the one hand, Figure 5b shows that the plastic deformation is very small. On the other hand, there is a certain distance between the TEM observation region and the main crack. An a/6<112> Shockley partial dislocation typically emits from grain boundaries and then glides into the grain interior [26], leaving behind a stacking fault, which exhibits fringe contrast when inclined to the sample surface. Dislocation contrast can also be observed in the fringe contrast regions, and these dislocations may be partial dislocations, as shown in Figure 6b. Figure 6a,c show that two directional stacking faults can be observed, indicating that the movement of partial dislocations on the conjugate {111} planes was activated. In addition to numerous stacking faults, weak dislocation tangles and isolated dislocations can also be observed, also indicating a small plastic deformation in the 0Al steel heat-treated at 600 °C.
Figure 7 shows typical deformation microstructure characteristics near the fracture surfaces of the 3Al steel impact samples heat-treated at 600 °C. The SFE of the 3Al steel increases to 20.5 mJ/m2 at −196 °C. Thus, numerous deformation twins can be observed, as shown in Figure 7a–c. The two twining systems were activated, and the two {111} planes are both in the edge-on position relative to the <011> viewing direction. The interior angle between the two variant twins is very close to 70.53°, as expected of any two neighboring {111} twin planes [1]. The twin thickness can be easily estimated as the two {111} planes are in the edge-on position, indicating that their thickness is essentially below 10 nm. Although some twin bundles are very thick, they are still composed of numerous fine twins. In addition, Figure 7b shows that steps appear at the intersection of two variant twins, indicating a strong twin–twin interaction.
The dislocation density in the 3Al steel is also much higher than that in the 0Al steel, which is consistent with the KAM maps. Numerous bands containing highly dense dislocation walls (HDDWs) along one variant {111} slip plane are visible, as shown in Figure 7d, indicating strong planar slip and dense dislocation sheets on the primary slip system [27]. The dislocation bands along the two slip planes can also be observed in Figure 7e. Therefore, planar slip is also a plastic deformation mechanism in the 3Al steel.
Figure 8 shows typical deformation microstructure characteristics near the fracture surfaces of the 5Al steel impact samples heat-treated at 600 °C. Figure 8a,b clearly show different dislocation configurations. Although dislocation tangles can be observed in Figure 8a, there are still dislocation traces along the {111} planes. The HDDWs along two variant {111} planes are clearly observed in Figure 8b, and numerous dislocations are also located between HDDWs. Besides these different dislocation configurations, numerous fine deformation twins can also be observed, as shown in Figure 8c. These results indicate that the 3Al and 5Al steels have the same deformation mechanisms.

4. Conclusions

The effects of heat treatment temperature on the microstructure and properties of Fe-18Mn-0.6C-0/3/5Al (in wt.%) steels were investigated, and all three steels showed a cryogenic impact brittle zone with the addition of different amounts of Al. The dependence of microstructure characteristics and deformation mechanisms on heat treatment temperature and Al content were analyzed and discussed. The main conclusions are as follows:
(1)
There is a cryogenic impact brittle zone (400–800 °C) in the three steels, and a very small amount of absorbed energy is usually observed at around 600 °C. The addition of Al can shift the onset of the brittle zone from 400 to 500 °C, thus narrowing the brittle zone. In addition, the absorbed energy can be lowered by around 89%, 48%, and 40% for the 0Al, 3Al, and 5Al steels after heat treatment at 600 °C compared to their hot-rolled counterparts, respectively. However, there is no tensile ductility brittle zone in the three steels, and the tensile properties are virtually unaffected by heat treatment at a temperature range of 150–800 °C.
(2)
It is found that grain morphology and size are not greatly affected by heat treatment at 300–800 °C, and the average grain size ranges of the 0Al, 3Al, and 5Al steels are 8.0–9.3, 6.3–7.3, and 6.0–6.6 μm, respectively. However, their grain sizes sharply increase to 32.6, 38.6, and 43.1 μm after heat treatment at 1000 °C.
(3)
The degree of plastic deformation is greatly lowered after heat treatment at 600 °C, exhibiting cleavage-like fracture in the 0Al steel as well as ductile fracture with small and shallow dimples in the 3Al and 5Al steels; this implies that the addition of Al can enhance the degree of plastic deformation. The main deformation mechanism of the 0Al steel is partial dislocation slip, which forms numerous stacking faults, while those of the 3Al and 5Al steels are twining and strong planar slip.

Author Contributions

Conceptualization, J.C.; methodology, L.X., Y.Z. and S.L.; writing—original draft preparation, L.X., Y.Z. and S.L.; writing—review and editing, H.H., B.Z. and N.J.; funding acquisition, J.C. All authors have read and agreed to the published version of the manuscript.

Funding

Open Fund of CNOOC Key Laboratory of Liquefied Natural Gas and Low-Carbon Technology (KJQZ-2024-1105).

Data Availability Statement

The original contributions presented in this study are included in the article. Further inquiries can be directed to the corresponding author.

Conflicts of Interest

Authors Xiao Li; Yuqi Zhang; Huan Huang; Bochao Zhang; Ningning Ji; Shuang Li were employed by the companies CNOOC Key Laboratory of Liquefied Natural Gas and Low-Carbon Technology and CNOOC Gas & Power Group, The remaining authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

References

  1. De Cooman, B.C.; Estrin, Y.; Kim, S.K. Twinning-induced plasticity (TWIP) steels. Acta Mater. 2018, 142, 283–362. [Google Scholar] [CrossRef]
  2. Bouaziz, O.; Allain, S.; Scott, C.P.; Cugy, P.; Barbier, D. High manganese austenitic twinning induced plasticity steels: A review of the microstructure properties relationships. Curr. Opin. Solid State Mater. Sci. 2011, 15, 141–168. [Google Scholar] [CrossRef]
  3. De Cooman, B.C.; Kwon, O.; Chin, K.-G. State-of-the-knowledge on TWIP steel. Mater. Sci. Technol. 2012, 28, 513–527. [Google Scholar] [CrossRef]
  4. Kim, Y.G.; Park, Y.S.; Han, J.K. Low temperature mechanical behavior of microalloyed and controlled-rolled Fe-Mn-Al-C-X alloys. Metall. Trans. A 1985, 16, 1689. [Google Scholar] [CrossRef]
  5. Choi, J.K.; Lee, S.G.; Park, Y.H.; Han, I.W.; Morris, J.W., Jr. High Manganese Austenitic Steel For Cryogenic Applications Available to Purchase. In Proceedings of the Twenty-Second International Offshore and Polar Engineering Conference, Rhodes, Greece, 17–22 June 2012; pp. 29–35. [Google Scholar]
  6. Kim, H.; Ha, Y.; Kwon, K.H.; Kang, M.; Kim, N.J.; Lee, S. Interpretation of cryogenic-temperature Charpy impact toughness by microstructural evolution of dynamically compressed specimens in austenitic 0.4C–(22–26)Mn steels. Acta Mater. 2015, 87, 332–343. [Google Scholar] [CrossRef]
  7. Chen, J.; Liu, N.; Liu, Z.Y.; Wang, G.D. Alloying design as well as strength and toughness of high Mn steels for LNG tank building. China Metall. 2023, 33, 73–80. [Google Scholar] [CrossRef]
  8. Sohn, S.S.; Hong, S.; Lee, J.; Suh, B.-C.; Kim, S.-K.; Lee, B.-J.; Kim, N.J.; Lee, S. Effects of Mn and Al contents on cryogenic-temperature tensile and Charpy impact properties in four austenitic high-Mn steels. Acta Mater. 2015, 100, 39–52. [Google Scholar] [CrossRef]
  9. Wang, X.; Sun, X.; Song, C.; Chen, H.; Han, W.; Pan, F. Enhancement of yield strength by chromium/nitrogen alloying in high-manganese cryogenic steel. Mater. Sci. Eng. A 2017, 698, 110–116. [Google Scholar] [CrossRef]
  10. Chen, J.; Lu, S.; Hou, Z.-Y.; Song, W.-W.; Liu, Z.-Y.; Wang, G.-D.; Furuhara, T. Atomic-scale understanding of twin intersection rotation and ε-martensite transformation in a high Mn twinning-induced plasticity steel. Acta Mater. 2024, 271, 119832. [Google Scholar] [CrossRef]
  11. A1106/A1106M: 2017; Standard Specification for Pressure Vessel Plate, Alloy Steel, Austenitic High Manganese for Cryogenic Application. ASTM International: West Conshohocken, PA, USA, 2017. [CrossRef]
  12. Chen, J.; Ren, J.-K.; Liu, Z.-Y.; Wang, G.-D. Interpretation of significant decrease in cryogenic-temperature Charpy impact toughness in a high manganese steel. Mater. Sci. Eng. A 2018, 737, 158–165. [Google Scholar] [CrossRef]
  13. Huang, Y.; Zhao, A.-M.; Mi, Z.-L.; Jing, H.-T.; Li, W.-Y.; Hui, Y.-J. Formability of Fe-Mn-C Twinning Induced Plasticity Steel. J. Iron Steel Res. Int. 2013, 20, 111–117. [Google Scholar] [CrossRef]
  14. Lee, D.; Yoon, G.-Y.; Park, S.-J.; Lee, J.S.; Heo, Y.-U. M23C6-assisted formation mechanism of grain boundary cellular structure and its role on the brittle fracture in an austenitic Fe-Mn-Al-C lightweight alloy. J. Alloys Compd. 2024, 972, 172865. [Google Scholar] [CrossRef]
  15. Lee, D.; Trang, T.; Kim, D.-H.; Yoon, G.-Y.; Park, S.-J.; Lee, J.S.; Heo, Y.-U. Boron-induced unique precipitation behavior at the grain boundary and its effect on the brittle fracture in a Fe-Mn-Al-C-based austenitic lightweight alloy. Scr. Mater. 2025, 259, 116574. [Google Scholar] [CrossRef]
  16. Curtze, S.; Kuokkala, V.T.; Oikari, A.; Talonen, J.; Hänninen, H. Thermodynamic modeling of the stacking fault energy of austenitic steels. Acta Mater. 2011, 59, 1068–1076. [Google Scholar] [CrossRef]
  17. Jo, Y.H.; Jung, S.; Choi, W.M.; Sohn, S.S.; Kim, H.S.; Lee, B.J.; Kim, N.J.; Lee, S. Cryogenic strength improvement by utilizing room-temperature deformation twinning in a partially recrystallized VCrMnFeCoNi high-entropy alloy. Nat. Commun. 2017, 8, 15719. [Google Scholar] [CrossRef]
  18. Yang, H.-S.; Jang, J.; Bhadeshia, H.; Suh, D. Critical assessment: Martensite-start temperature for the γ→ε transformation. Calphad 2012, 36, 16–22. [Google Scholar] [CrossRef]
  19. Dastur, Y.N.; Leslie, W.C. Mechanism of work hardening in Hadfield manganese steel. Metall. Trans. A 1981, 12, 749–759. [Google Scholar] [CrossRef]
  20. Hwang, J.-K. Deformation behaviors of various Fe–Mn–C twinning-induced plasticity steels: Effect of stacking fault energy and chemical composition. J. Mater. Sci. 2020, 55, 1779–1795. [Google Scholar] [CrossRef]
  21. Tobler, R.L.; Meyn, D. Cleavage-like fracture along slip planes in Fe-18Cr-3Ni-13Mn-0.37N austenitic stainless steel at liquid helium temperature. Metall. Trans. A 1988, 19, 1626–1631. [Google Scholar] [CrossRef]
  22. Tomota, Y.; Xia, Y.; Inoue, K. Mechanism of low temperature brittle fracture in high nitrogen bearing austenitic steels. Acta Mater. 1998, 46, 1577–1587. [Google Scholar] [CrossRef]
  23. Liu, S.; Liu, D.; Liu, S. Transgranular fracture in low temperature brittle fracture of high nitrogen austenitic steel. J. Mater. Sci. 2007, 42, 7514–7519. [Google Scholar] [CrossRef]
  24. Lee, J.; Sohn, S.S.; Hong, S.; Suh, B.-C.; Kim, S.-K.; Lee, B.-J.; Kim, N.J.; Lee, S. Effects of Mn Addition on Tensile and Charpy Impact Properties in Austenitic Fe-Mn-C-Al-Based Steels for Cryogenic Applications. Met. Mater. Trans. A 2014, 45, 5419–5430. [Google Scholar] [CrossRef]
  25. Curtze, S.; Kuokkala, V.-T. Dependence of tensile deformation behavior of TWIP steels on stacking fault energy, temperature and strain rate. Acta Mater. 2010, 58, 5129–5141. [Google Scholar] [CrossRef]
  26. Wang, L.; Guan, P.; Teng, J.; Liu, P.; Chen, D.; Xie, W.; Kong, D.; Zhang, S.; Zhu, T.; Zhang, Z.; et al. New twinning route in face-centered cubic nanocrystalline metals. Nat. Commun. 2017, 8, 2142. [Google Scholar] [CrossRef]
  27. Gutierrez-Urrutia, I.; Raabe, D. Multistage strain hardening through dislocation substructure and twinning in a high strength and ductile weight-reduced Fe–Mn–Al–C steel. Acta Mater. 2012, 60, 5791–5802. [Google Scholar] [CrossRef]
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Figure 1. Engineering stress–strain curves of (a) 0Al, (b) 3Al, and (c) 5Al steels.
Figure 1. Engineering stress–strain curves of (a) 0Al, (b) 3Al, and (c) 5Al steels.
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Figure 2. Charpy impact absorbed energy of 18Mn-xAl steels with different heat treatment temperatures.
Figure 2. Charpy impact absorbed energy of 18Mn-xAl steels with different heat treatment temperatures.
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Figure 3. EBSD-IPF maps of (a–f) 0Al, (g–l) 3Al, and (m–r) 5Al steels.
Figure 3. EBSD-IPF maps of (a–f) 0Al, (g–l) 3Al, and (m–r) 5Al steels.
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Figure 4. SEM fractographs of the Charpy V-notch impact samples of the (a–c) 0Al, (d–f) 3Al, and (g–i) 5Al steels.
Figure 4. SEM fractographs of the Charpy V-notch impact samples of the (a–c) 0Al, (d–f) 3Al, and (g–i) 5Al steels.
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Figure 5. EBSD-IPF maps of cross-sectional regions close to the fracture surfaces of the (a–c) 0Al, (d–f) 3Al, and (g–i) 5Al steels. (a,d,g) Hot-rolled steels, (b,e,h) heat-treated steels with a temperature of 600 °C, and (e,f,i) heat-treated steels with a temperature of 1000 °C. The inset at the bottom left of each IPF map is a KAM map.
Figure 5. EBSD-IPF maps of cross-sectional regions close to the fracture surfaces of the (a–c) 0Al, (d–f) 3Al, and (g–i) 5Al steels. (a,d,g) Hot-rolled steels, (b,e,h) heat-treated steels with a temperature of 600 °C, and (e,f,i) heat-treated steels with a temperature of 1000 °C. The inset at the bottom left of each IPF map is a KAM map.
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Figure 6. TEM images of typical deformation microstructures near the fracture surface of the 0Al steel with a heat-treating temperature of 600 °C. (a–d) Bright-field (BF) TEM images showing numerous stacking faults. (e) A BF TEM image showing weak dislocation tangles. (f) A BF TEM image showing isolated dislocations.
Figure 6. TEM images of typical deformation microstructures near the fracture surface of the 0Al steel with a heat-treating temperature of 600 °C. (a–d) Bright-field (BF) TEM images showing numerous stacking faults. (e) A BF TEM image showing weak dislocation tangles. (f) A BF TEM image showing isolated dislocations.
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Figure 7. TEM images of typical deformation microstructures near the fracture surfaces of the 3Al steel with a heat-treating temperature of 600 °C. (a) BF and (b,c) corresponding dark-field (DF) TEM images, exhibiting at least two activated twinning systems. (d,e) BF TEM images showing numerous dislocation bands and dislocation interactions. The inset at the bottom right in (a) is the selected area electron diffraction pattern (SADP) obtained from the circled region.
Figure 7. TEM images of typical deformation microstructures near the fracture surfaces of the 3Al steel with a heat-treating temperature of 600 °C. (a) BF and (b,c) corresponding dark-field (DF) TEM images, exhibiting at least two activated twinning systems. (d,e) BF TEM images showing numerous dislocation bands and dislocation interactions. The inset at the bottom right in (a) is the selected area electron diffraction pattern (SADP) obtained from the circled region.
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Figure 8. TEM images of typical deformation microstructures near the fracture surfaces of the 5Al steel with a heat-treating temperature of 600 °C. (a,b) BF TEM images exhibiting dislocation configurations. (c) A BF TEM image exhibiting fine deformation twins; the inset at the bottom right is the corresponding DF image.
Figure 8. TEM images of typical deformation microstructures near the fracture surfaces of the 5Al steel with a heat-treating temperature of 600 °C. (a,b) BF TEM images exhibiting dislocation configurations. (c) A BF TEM image exhibiting fine deformation twins; the inset at the bottom right is the corresponding DF image.
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Table 1. Chemical compositions and SFEs of tested steels (wt.%).
Table 1. Chemical compositions and SFEs of tested steels (wt.%).
SteelCSiMnAlPSNOFeSFE (25 °C), mJ/m2SFE (−196 °C), mJ/m2
0Al0.600.6118.040.040.00210.01100.0300.0006Bal.24.4010.74
3Al0.580.4717.873.010.00630.00560.0310.0006Bal.44.6620.54
5Al0.620.4918.375.160.00240.00710.0290.0007Bal.61.6229.71
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Xiao, L.; Zhang, Y.; Huang, H.; Zhang, B.; Ji, N.; Li, S.; Chen, J. Effect of Heat Treatment Temperature on the Microstructure and Mechanical Properties of Fe-18Mn-0.6C-xAl. Metals 2025, 15, 927. https://doi.org/10.3390/met15080927

AMA Style

Xiao L, Zhang Y, Huang H, Zhang B, Ji N, Li S, Chen J. Effect of Heat Treatment Temperature on the Microstructure and Mechanical Properties of Fe-18Mn-0.6C-xAl. Metals. 2025; 15(8):927. https://doi.org/10.3390/met15080927

Chicago/Turabian Style

Xiao, Li, Yuqi Zhang, Huan Huang, Bochao Zhang, Ningning Ji, Shuang Li, and Jun Chen. 2025. "Effect of Heat Treatment Temperature on the Microstructure and Mechanical Properties of Fe-18Mn-0.6C-xAl" Metals 15, no. 8: 927. https://doi.org/10.3390/met15080927

APA Style

Xiao, L., Zhang, Y., Huang, H., Zhang, B., Ji, N., Li, S., & Chen, J. (2025). Effect of Heat Treatment Temperature on the Microstructure and Mechanical Properties of Fe-18Mn-0.6C-xAl. Metals, 15(8), 927. https://doi.org/10.3390/met15080927

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