Microstructural and Mechanical Comparison Between the Plastic Deformation and Annealing Conditions of Low-C high-Mn (Fe-27Mn-1Si-0.027C) Steel for Cryogenic Applications

Abstract

The high-Mn steel with the composition of Fe-27Mn-1Si-0.027C (wt.%) was selected. Two samples with different grain sizes were prepared by hot rolling followed by annealing: one was annealed at 1100°C for 90 minutes after 60% hot rolling, and the other was annealed at 1100°C for 60 minutes after 80% hot rolling. The study focused on the relationship between the microstructural evolution of the low-C high-Mn steel and its tensile and impact properties influenced by the stacking fault energy (SFE) at room and cryogenic temperatures. Through SEM and SEM-EBSD experimental techniques, ε-martensite phase was observed, and its mechanical response was compared under room and cryogenic temperatures. In consequence, the low-C high-Mn steel did not present the TWIP effect, but the TRIP effect (i.e., dislocation slip deformation) was detected due to the presence of ε-martensite formed by mechanical straining, increasing its volumetric fraction as the testing temperature decreased. Even with the negative presence of the ε-manartensite phase, the mechanical properties of the low-C high-Mn steel (i.e., strength, ductility and toughness) at the cryogenic temperature testing were similar to those of commercial cryogenic alloys currently used for low temperature combustible storage applications.

Keywords:
High-Mn content; stacking fault energy; ε-martensite; cryogenic mechanical testing

1. Introduction

The behavior of metals at working temperatures lower than the ambient is largely determined by their crystalline structure. For cryogenic applications, near liquid nitrogen temperatures, metals with a face-centered cubic (FCC) atomic structure are widely used because the tensile strength is more temperature-dependent than the yield strength and frequently become more ductile as working temperatures drop¹. On the other hand, metals with a body-centered cubic (BCC) atomic structure have low resistance to ductile-brittle transition behavior (i.e., loss of ductility in a limited temperature range)¹.

In this context, FeMnC alloys, with high-Mn content (i.e., 20-30 wt.%), represent a recent development in the field of austenitic steels and they signify interesting candidate’ materials for extreme low temperature working conditions. These high-Mn steels are being studied by several research groups worldwide²-⁴. Thus, TWIP and TRIP effects, well-differentiated deformation mechanisms that can occur in these high-Mn steels, are of great scientific and economic importance.

Their occurrence strongly depends on the chemical composition, defining the stacking fault energy (SFE) value. Thus, the SFE governs the way in which the alloy behaves structurally. On the one hand, at higher-Mn contents, the TWIP effect is well activated in the range of 18-60 mJ/m² of SFE, thus causing austenitic mechanical twinning due to the accumulation of cold work. On the other hand, the TRIP effect is manifested at SFE values below 18 mJ/m², in which the alloy experiences the austenite-martensite transformation caused by higher localized strains⁵,⁶. Hence, due to the localized plastic strain, elastic stress accumulation and the earlier transformations, two types of martensite may form in high-Mn austenitic steels that manifest the TRIP effect. The first one is ε-martensite with a hexagonal compact (HC) crystal structure, and the second one is α'-martensite presenting a body-centered cubic (BCC) or a body-centered tetragonal (CCC/CBT) crystal structure⁷. The kinetic transformation sequence is stated as ã_FCC (austenite) → å_HCP (HCP-martensite) → á^'_BCC (BCC-martensite).

Under a certain amount of deformation straining, the volumetric fraction of ε-martensite (also called mechanical ɛ-martensite) increases with deformation until a peak strain is reached. But at larger deformations, ɛ-martensite decreases, while the amount of α'-martensite gradually increases⁸. Also, decreasing the temperature for achieving a cryogenic temperature criterion, the SFE drops in about 20-60%, thus inducing the martensitic transformation. Mechanical ɛ-martensite transformation in the simplest Fe-Mn-C system may also be influenced by the alloying content adjustment as a result of the SFE alteration⁸-¹⁰. It has been found that Cr additions can reduce deformation twinning and directly reduce the amount of formed ɛ-martensite, resulting in an absorbed energy enhancement¹¹. To further prevent ɛ-martensite transformation, Al and N (e.g., Cr+N) are added as well¹²-¹⁴. Likewise, fine austenitic grain size control represents an alternative practice to prevent ɛ-martensite transformation¹⁵. Curtze and Kuokkala¹⁶ have shown that by conserving the SFE for the deformation twinning regime at cryogenic temperatures, it is also possible to obtain ɛ-martensite by the martensitic transformation due to the existence of low-solute segregated band structures.

Nonetheless, martensite transformation during low-temperature deformation, grain boundary impurity segregation, dislocation density “slipping-off” fracture, and even hydrogen embrittlement are the principal causes for cryogenic mechanical properties degradation¹⁷,¹⁸. Alongside, the presence of ɛ-martensite necessitates several analyses to understand microstructural evolution, including precipitation and atom segregation to grain boundaries. Additionally, prior heat treatment conditions and associated deformed microstructures are considered to explain the mechanical response of high-Mn steels under cryogenic temperatures¹⁹.

Thus, thermal stability of the austenite at lower temperature can be achieved by controlling the decrease of the SFE values, which directly limits ε-martensite formation²⁰. An additional material hardening can be obtained due to both mechanical twinning and ε-martensite interactions. However, ɛ-martensite and its interaction with other microstructural features mostly promotes brittle behavior²¹. As well, the banded morphology of ɛ-martensite deteriorates cryogenic toughness²². Conversely, heat-treated microstructures help to increase the absorbed energy at cryogenic temperatures because the homogeneous grained condition limits the free path for crack propagation²³. In this manner, the control of grain size can also influence the SFE values at cryogenic temperatures and restrict the martensitic transformation²⁴,²⁵. Thus, mechanical performance can be improved with a slight ductility deterioration²⁶-²⁸.

In the specific case of the low-C high-Mn steels, they can represent a good way for impact behavior improvement at cryogenic temperatures by avoiding the ductile-brittle transition. As it can be inferred, the presence and effects of ɛ-martensite in low-C high-Mn steels still need a systematic exploration focused on tailored cryogenic applications. Hence, the aim of the present study was to investigate the relationship between the microstructure evolution and mechanical properties at room and cryogenic temperatures of a low-C high-Mn steel.

2. Experimental Procedures

2.1. Fabrication of the low-C high-Mn steel

The studied low-C high-Mn steel (termed as Fe-27Mn-1Si-0.027C) was manufactured by the centrifugal casting process with dimensions of 525x50x150 mm. The chemical composition of the steel is Fe-27Mn-0.027C-1.0Si-0.41Cr-0.19Ni (in wt.%) and was determined by Optical Emission Spectroscopy (OES) performed on a Shimadzu Spectrometer, model PDA 7000. The hot rolling of the steel, employing the as-cast samples with an initial thickness of 9.0 mm, was conducted in a laboratory rolling mill (Coelho Machines Manufacturer/ Model LE-200B) operated at a rotational speed of 8 m/min. Hot-rolled steel plates were obtained with a total reduction in the cross-section of 60% and 80%, respectively. After the last pass, the hot-rolled steel plates were fast-cooled in water. Then, quite different annealing heat treatments, in terms of the holding time, were applied to the hot-rolled plate samples with the intention of obtaining a final free-segregated microstructure, conditions that are comparable between the employed samples. In a first scenario, the 60% deformed sample was annealed at 1100 °C with a total holding time of 90 minutes. In a second scenario, the 80% deformed sample was annealed at the same temperature of 1100 °C, but a total holding time of 60 minutes was used. After that, both samples were fast cooled in water. The reader should note that, after hot rolling, the samples were not cold rolled. So, at this experimental stage, the two heat-treated steel conditions were subject of characterization and thus the samples were identified as (60%-annealed-by-90-min) and (80%-annealed-by-60-min), respectively.

2.2. Structural characterization

After the heat treatment, the X-ray diffraction (XRD) technique was employed in order to identify the phases present in both annealed samples. Thus, XRD analysis was performed using a Panalytical X’Pert Pro diffractometer equipped with Co (Kα₁=0,178897 nm) radiation source with a scanning step of 0.013° and a counting time of 13.7 s from 45 to 115° in 2θ. As the 2θ position of each diffraction peak can be affected by some heterogeneous local strains, a mathematical adjustment procedure was carried out by applying the Sherrer formula²⁹. In this regard, a computer code based on references⁹,¹⁰,¹⁶ was developed and adapted to the current chemical composition. The volume fraction of the austenite (γ) and martensite (ε) phases was obtained by the direct comparison method, as described by Cullity³⁰ and taking into account the different values in the scattering factors of these two phases for high-manganese steel. The performed calculation in this work is based on the scattering factor (f) values of Fe, Mn, Si, Cr, and Ni³¹ for a certain angle of incidence and a given wavelength. In consequence, the f was obtained considering equal chemical composition in both austenite and ɛ-martensite phases. Also, the quantification of the stacking fault energy (SFE) was calculated considering the thermodynamic model of Allain et al.⁹ for the general Fe-Mn-C system. The values ​​of the free phase transformation energy, chemical composition, temperature and the ASTM grain size were also taken into account for the calculations.

2.3 Microstructural characterization

Afterwards, both annealed samples were metallographically prepared by using the standardized procedure, including mechanical grinding, general polishing with diamond paste and a final surface polish using colloidal silica suspension. A chemical etching consisting of 40% HNO₃ + 60% H₂O was employed for revealing the microstructure by immersion under 3 seconds. The grain size measurement was conducted in accordance with ASTM E112³² procedure by quantifying the number of grains per unit of area based on metallographic images captured in an optical microscope at a specified magnification. Details of the microstructure were obtained using a field emission gun - scanning electron microscope (FEG-SEM) under the secondary electron mode. In addition, electron backscatter diffraction (EBSD) analysis was carried out in a SEM FEI Quanta FEG 450, operated at 20 kV, using a working distance of 13 mm and with a tilted angle of 70° and a scan step size of 0.3 µm. Data post-processing was rendered using HKL Channel 5 software. Thereby, grains presenting a misorientation above 15° with respect to their adjacent grain neighbors were identified as high-angle boundaries, meanwhile, the misorientation between 2 and 15° was identified as low-angle boundaries. Texture images were generated using the ATEX data processing software package. The orientation distribution functions (ODFs) were plotted at φ₂=0° and at the 45° section of the Euler space.

2.4. Mechanical characterization

Vickers hardness tests were performed on both annealed samples in accordance with the ASTM E92³³ procedure employing a Vickers Hardness Tester System A - Series 810 - 810-440A with a total load of 1 kg_f for 13 seconds as load time for each indentation, considering 10 measurements and an indentation separation of about 500 µm. So, the final hardness was reported as the average value. Also, tensile tests were performed at room temperature and at -196 °C as the cryogenic testing temperature on both annealed samples. Liquid nitrogen was used as the cryogenic medium. Tension samples were obtained with a gauge length of 16 mm and overall length of 55 mm and tested using MTS 370 universal testing machine under a strain rate of 2×10⁻² s⁻¹ as stated by the ASTM E8/E8M³⁴. After tensile testing, fractured surfaces were subjected to fractographic analysis using a FEG-SEM. Charpy impact tests were also performed on the same room and cryogenic sample conditions following the ASTM E23³⁵ procedure in which samples were obtained in the rolling direction employing a hammer of 300 J. The fractured surfaces of the Charpy tested samples were subjected to fractographic analysis using a FEG-SEM.

3. Results and Discussion

3.1. Structural characterization

Figure 1 (a-b) shows the XRD patterns for both annealed samples. As it can be observed, it is confirmed that any other phases like α'-martensite are not present in both steel conditions, but on the contrary, austenite and ε-martensite are present. The high Mn content in the alloy and also the low carbon content favored the TRIP effect (i.e., disagreement slip deformation). There have been identified the characteristic diffraction peaks corresponding to the planes (111) and (220) for a material with FCC structure (γ-austenite) and various planes such as (100), (002), (101) and (103) that match with the HCP structure (ε-martensite). By observing the differences in intensity between the diffraction peaks of both annealed conditions, it is possible to anticipate differences in the percentage of γ-austenite and ε-martensite phases. In consequence, figure 1c presents the quantified volumetric fraction of the γ-austenite and ε-martensitic phases obtained by the direct comparison method. It is highlighted that the amount of ε-martensite has the highest value in the 80% deformation condition in comparison with the 60% one. A similar behavior was found by Li et al.³⁶ observing that the Mn content enhanced the stability of austenite and ε-martensite and also that the austenite grain refinement had a larger effect on the γ → ε transformation than in the γ → α′ transformation. It is worth noting that in deformed metallic materials, compression fields can cause an increase in the 2θ angle and tensile fields cause the contrary effect, it decreases. Therefore, the peak, which should be well defined, becomes less intense and wider²⁹. A crystal that has not undergone deformation presents more intense diffraction peaks. Peak broadening is caused by dislocations density (one-dimensional defects), whereas displacement of the diffraction peak position (i.e., the diffraction angle) is caused by crystal twinning. Coincidence site lattice (CSL) boundaries, including twining formation, is also a consequence of lattice distortion and diffraction peak asymmetry. It can be suggested that these effects are correlated with the increase of the ε-martensite and, additionally, the lattice distortion induced by the plastic strain on the steel³⁷-⁴⁰. In addition, the calculated SFE value for Fe-27Mn-1-Si-C steel at room temperature was ≅10.0 mJ/m² and at the cryogenic temperature (T = -196 °C) it was ≅5.0 mJ/m².

Figure 1
Low-C high-Mn steel annealed samples: (a-b) XRD patterns and (c) Phase content calculation by the direct comparison method.

3.2. Microstructure characterization

For the cross-section, Figure 2 shows the scanning electron microscopy micrographs of the Fe-27Mn-1Si-C steel for both the 60%-annealed-by-90-min sample and the 80%-annealed-by-60-min sample, respectively. In Figure 2a, it can be stated that the austenitic grains are deformed, but slightly larger compared to the sample with greater applied deformation, a difference confirmed in Table 1, which contains the average values of the calculated austenitic grain size for both annealed conditions. It is possible to observe a small difference in grain size of about 10 µm due to the total deformation applied to the material (i.e., 60% and 80%, respectively). The classical model proposed by G.B. Olson and M. Cohen⁴¹ indicates a dependence of SFE on grain size, where larger grain sizes tend to reduce the SFE. In this context, the grain size also contributes to the formation of ɛ-martensite, which is associated with the stacking fault energy (SFE) of the alloy. From Figure 2b, it is possible to observe the microstructure of the sample with greater deformation, presenting an increase in the amount of ɛ-martensite phase.

Figure 2
FEG-SEM secondary electron micrographs of the low-C high-Mn steel: (a) Sample: 60%-annealed-by-90-min and (b) Sample: 80%-annealed-by-60-min.

Table 1
Grain size of the low-C high-Mn steel in both annealed conditions.

Steel Sample	Number of Intercepts	Grain Size (µm)
60%-annealed-for-90-min	157	63 ± 2
80%-annealed-for-60-min	183	51 ± 1

The EBSD maps for both annealed conditions are shown in Figure 3. The IQ band contrast maps are presented in Figure 3 (a,e) and Figure 3 (b,f) presents the IPF maps for the austenitic phase (i.e., γ-IPF). Also, colored it can be seen the grain orientation of the austenite phase and in white the ε-martensite phase as shown in Figure 3 (c,g). Similarly, in Figure 3 (d,h) it can be observed the identification and quantification of the phases, in which the measured values ​​correspond to 61 and 48% for the γ-austenite and 39 and 52% for the ε-martensite for both annealed conditions, respectively. This indicates that the ε-martensite increase in the 80%, annealed by 60 min sample, by reducing the austenitic phase. The crystallographic orientation in relation to the direction of loading is a very important aspect on the formation of ε-martensite. Not every austenitic grain has an equal tendency to form martensitic phases. Note that some austenite grains contain annealing twins with a different crystal orientation and ε-martensite in both annealed conditions. The formation of ε-martensite occurs by the slip of Shockley partial dislocations with each {111}_γ. The Shoji-Nishiyama (S-N) orientation relationship (OR) is established between γ-austenite and ε-martensite, i.e., ${\left(111\right)}_{\gamma }//{\left(0001\right)}_{\epsilon },{[110]}_{\gamma }//{\left[11\overline{2}0\right]}_{\epsilon }$ ⁴²,⁴³. However, the formation of ε-martensite is an intermediate phase transformation step to form α'-martensite.

Figure 3
EBSD maps for the annealed low-C high-Mn steel: (a,e) IQ band contrast images generated for the 60%-annealed-by-90 min and 80%-annealed-by-60-min samples, respectively, (b,f) γ-IPF maps generated for the 60%-annealed-by-90 min and 80%-annealed-by-60-min samples, respectively, (c,g) ε-IPF maps generated for the 60%-annealed-by-90 min and 80%-annealed-by-60-min samples, respectively, and (d,h) Austenite and ɛ-martensite phase maps for the 60%-annealed-by-90 min and 80%-annealed-by-60-min samples, respectively.

The Taylor factor analysis (shown in Figure 4) was performed to evaluate the mechanical resistance of the grains as the strain increased. The uniform distribution of the Taylor factor values throughout the bulk material (see Figure 4 (a, b, d, e)) suggests that the plastic strain induced by the martensitic transformation is absorbed by several austenitic grains surrounding the newly formed ε-phase nuclei. This result appears to support the hypothesis proposed by other authors⁴⁴ that a high Taylor factor indicates greater difficulty in deformation through dislocation slip, promoting the formation of ε-martensite with increasing the strain. Also, it has been showed that high-Mn steels display substantial residual stresses resulting from the non-equilibrium solidification phenomenon encountered during their plastic deformation⁴⁵. When subjected to strain, some of these residual stresses are relieved through deformation-induced martensitic transformation, resulting in residual plastic strain due to the volume expansion associated with the γ→ε phase transformation⁴⁶. The pronounced misorientation and its uneven distribution shown in Figure 4 (c,f), suggesting that the plastic strain from the martensitic transformation is accommodated by a smaller number of austenitic grains surrounding the newly nucleated ε-phase. In this scenario, a heterogeneous stress field develops near the grain boundaries due to the constraints imposed by the grain boundaries and the impingement of ε-plates by neighboring grains⁴⁷. The present high-Mn steel has a low value of SFE, value at the cryogenic testing temperature, making it prone to more stacking faults formation. When subjected to loading, dislocation glide is naturally inhibited due to the alloy's low SFE. Consequently, slip becomes highly localized on both pre-existing and strain-induced stacking faults with various crystallographic orientations. These localized regions give rise to different variants of deformation-induced ε-martensite, as depicted in Figure 4. Thus, the ε-martensite variants that form on one of the (111)_γ planes can be achieved by rotating the basal plane (0001)_ε by 70.5°. The direction $\left[11\overline{2}\right]$_γ in austenite aligned with the directions $\left[1\overline{2}10\right]$_ε in the ε-phase⁴⁸. As the load increases, the ε-martensite plates undergo plastic deformation, resulting in changes in the misorientation angles. From the determined Euler angles for the phases showed in Figure 4 (a,b,d,e), ε-martensite variants were calculated based on the Humbert model. The Euler angles found in the present work are in agreement with the values reported by Antunes et al.⁴⁵.

Figure 4
(a,d) Taylor factor maps generated for the 60%-annealed-by-90 min and 80%-annealed-by-60-min samples, respectively, (b) Euler orientation maps for the 60%-annealed-by-90 min and 80%-annealed-by-60-min samples, respectively, (c) Misorientation degree chart generated for the 60%-annealed-by-90-min sample and (f) Misorientation degree chart generated for the 80%-annealed-by-60-min sample.

3.4. Mechanical characterization

The true stress - true strain curves of the studied steel conditions can be seen in Figure 5. The cryogenic temperature tests for both annealed conditions present values ​​of ultimate tensile strength greater than 1200 MPa, a yield strength limit of about 400 MPa, and total elongation of nearly 30%. In contrast, tensile tests carried out at room temperature present better elongation close to 40%, but less mechanical resistance than the cryogenic ones. Then, for both annealing treatments it can be stated that under the current processing conditions for both testing temperatures, the ductility can be maintained and at the same time also a high mechanical resistance can be obtained. Table 2 resumes the measured mechanical properties for the high-Mn steel under both annealed conditions and room and cryogenic tensile testing conditions. So, it can be verified that the values ​​of yield stress, ultimate tensile strength (e.g., greater than 800 MPa), and total elongation are adequate ​​for the construction of pressure vessels for the storage of liquefied natural gas⁴⁹. In the same way, a gain in hardness is obtained in the 80%-annealed-by-60-min sample.

Figure 5
True stress - true strain curves for the annealed high-Mn steel samples under room and cryogenic temperature conditions.

Table 2
Low-C high-Mn steel: mechanical properties at room and cryogenic temperature for the both deformation and annealed conditions (values defined from the true stress-true strain curves).

Steel Sample	Test Temperature (°C)	Yield Strength (MPa)	Ultimate Tensile Strength (MPa)	Total Elongation (%)	Vickers Hardness (HV/25 °C)
60%-annealed -for-90-min	25	209	610	33.46	301 ± 5
	-196	429	936	23.25
80%-annealed -for-60-min	25	333	722	34.26	326 ± 3
	-196	446	972	27.98

The fractography images of the tensile fractured surfaces are shown in Figure 6. A mixed surface can be observed in the samples tested at room temperature presenting a ductile-dimpled fracture with microvoids (see Figure 6 (a,c)), which are features of ductile fracture and other quasi-cleavage regions (near cleavage) that characterize fragile fracture. Additionally, Figure 6 (b,d) shows the fractographies for the 80% deformed condition that showed similar fracture characteristics to the previous condition. Surfaces of the samples tested at cryogenic temperature showed lower ductility compared to those tested at room temperature. The ductility of the highly deformed 80% steels is lower, almost equaling another condition of less deformation due to the existence of about 10% ε-martensite more than the less deformed. In consequence, the variation in the ε-martensite fraction greatly influences tensile ductility. At both testing temperatures, the specimens showed the same type of mixed fracture as it has been found in similar alloys⁵⁰. In general, the fracture surface subject to tension of the high-Mn steel is complex with different mechanisms that seem to operate simultaneously⁵¹. The material presents of good toughness and ductility even at ultra-low temperatures, holding up very well the ductile-fragile transition. Thus, it can be said that the characteristics of ductile fracture, such as the presence of microvoids, are present in both deformations and tested temperature conditions.

Figure 6
SEM fracture surfaces of tensile tested high-Mn steel: (a) Sample: 60%-annealed-90-min and tested at 25 °C, (b) Sample: 60%-annealed-90-min and tested at -196 °C, (c) Sample: 80%-annealed-6-min and tested at 25 °C and (d) Sample: 80%-annealed-60-min and tested at -196 °C.

Table 3 presents the obtained absorbed energy values of the low-C high-Mn steel. It can be seen that sample with the 80% of deformation has the best values of absorbed energy in both room and cryogenic conditions when compared with the 60% of deformation steel conditions. These values ​​correspond with the effect of the predominantly austenitic microstructure with the ɛ-martensitic phase fraction that turned out to be beneficial for the material toughness. As mentioned above, the stability of FCC materials prevents the fragile-ductile transition occurring in the material at low temperatures due to the high-Mn content. Although, greater fracture ductility at room temperature is also confirmed by the high absorbed energy values. It can also be stated by analyzing Figure 2 that the higher percentage of ε-martensite in almost 9% of the 80% deformed steel condition in relation with the 60% deformed steel condition has not affected the ductility behavior negatively.

Table 3
Absorbed energy values for the low-C high-Mn steel tested under room and cryogenic testing temperature conditions.

Steel samples	Absorbed energy at room temperature, 25 °C (J)	Absorbed energy at cryogenic temperature, -196 °C (J)
60%-annealed -for-90-min	96	46
80%-annealed -for-60-min	110	50

Fractured surface fractography representing the Charpy impact specimens tested at room and cryogenic temperature conditions are shown in Figure 7. A typical behavior of ductile-brittle mixed fracture can be observed with few dimples and near cleavage regions. A small difference in the ductility of the material across the fracture surface of the specimens tested at room and cryogenic temperature is observed, a fact confirmed by the lowest absorbed energy value presented in Table 3, for the 60% deformed steel. This means that the energy absorption values were reduced, but both deformation conditions (i.e., 60 and 80%, respectively) ended up having similar absorption values ​​at both room and cryogenic testing temperatures. It must be attributed to the fact that ε-martensite is less detrimental to the ductility behavior of the high-Mn steel than α'-martensite. The reason for this behavior is the fact that steel has high levels of Mn, stabilizing the austenitic field, thus avoiding the ductile-brittle transition that occurs in metallic materials with the FCC structure at very low temperatures⁵².

Figure 7
SEM fracture surfaces of the fractured Charpy specimens: (a) Sample: 60%-annealed-90-min and tested at 25 °C, (b) Sample: 60%-annealed-90-min and tested at -196 °C, (c) Sample: 80%-annealed-6-min and tested at 25 °C and (d) Sample: 80%-annealed-60-min and tested at -196 °C.

5. Conclusions

Based on the experimental results that were obtained and analyzed through the present study, it can be concluded that the low SFE (10 mJ/m² at room temperature and 5 mJ/m² at the cryogenic temperature) in the low-C high-Mn steel promotes strain-induced ε-martensite formation, with phase fractions of 52 and 39%, respectively, for both 60%-annealed-90-min and 80%-annealed-60-min steel samples. This ε-martensite phase fraction significantly influences tensile resistance and ductility, achieving an ultimate tensile strength exceeding 800 MPa and at the same time enhancing the toughness behavior at the cryogenic testing temperature. The 80% deformed steel demonstrated superior absorbed energy at both testing temperature conditions, mainly attributed to a brittle-to-ductile transition. These findings suggest the steel's suitability for cryogenic gas storage applications, where high toughness and fracture resistance are critical.

6. Acknowledgments

The authors acknowledge the support of the Research and Support Foundation of Maranhão (FAPEMA), the Brazilian National Council for Scientific and Technological Development (CNPq), the Coordenação de Aperfeiçoamento de Pessoal de Nível Superior-Brasil (CAPES), and the PRPGI at IFMA.

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

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Edited by

• Associate Editor:
  Hugo Sandim.

• Editor-in-Chief:
  Luiz Antonio Pessan.

Data availability

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

Publication Dates

• Publication in this collection
  01 Dec 2025
• Date of issue
  2025

History

• Received
  27 May 2025
• Reviewed
  03 Sept 2025
• Accepted
  26 Oct 2025