Optimization of Mechanical Properties of High-Manganese Steel for LNG Storage Tanks: A Comprehensive Review of Alloying Element Effects

Abstract

High-manganese austenitic steel represents an innovative variety of low-temperature steel used in the construction of liquefied natural gas (LNG) storage tanks. This steel boasts remarkable characteristics such as exceptional plasticity, superior toughness at cryogenic temperatures, and robust fatigue resistance, all while providing significant cost benefits. By utilizing high-manganese steel, the material manufacturing costs can be considerably lowered, simultaneously ensuring the long-term stability and safety of LNG storage tanks. The alloying design is pivotal in attaining superior performance in high-manganese steel. Choosing the right chemical components to control the stacked fault energy (SFE) of high-manganese steel and fine-tuning its structure can further improve the balance between strength and plasticity. Summarizing the advancements in alloying design for high-manganese steel is of great importance, as it offers a foundational dataset for correlating the chemical composition with the performance. Therefore, this paper outlines the deformation mechanisms and the principles of low-temperature brittleness in high-manganese austenitic steel, and from this foundation, it explicates the precise functions of alloying elements within it. This aims to provide a reference for future alloying designs and the industrial deployment of high-manganese steel in LNG storage tanks.

1. Introduction

Since 2000, the global natural gas sector has experienced a renaissance. Concurrently, the global trade volume of LNG has seen a swift ascent, soaring from 137 billion cubic meters in 2000 to an impressive 521 billion cubic meters in 2021, which now represents 39.2% of the total trade in natural gas [1,2,3]. As demand for LNG storage and transportation continues to climb, research into materials for LNG storage tanks is becoming increasingly vigorous. Given that LNG storage tanks must endure the severe low temperature of −162 °C during operation, the materials used for these tanks necessitate exceptional low-temperature performance, resistance to brittle cracking, and the capacity to arrest cracks [4,5].

Currently, invar alloys, aluminum alloys, and austenitic stainless steel are among the mainstream materials for LNG storage tanks, with the widely adopted and top-performing 9% nickel steel standing out as the preferred choice [6,7]. However, these materials come with drawbacks, such as lower design strength and poor weldability [8]. Of particular concern is the issue of pricing; with Ni prices consistently high and even rising annually, a 10% increase would add approximately $680,000 to the construction cost of a single 160,000 m³ storage tank project [9]. Consequently, there is a pressing need to develop cost-effective steel alternatives for LNG storage tanks without compromising on quality. Currently, two principal alternatives to 9% Ni steel materials have been identified: one option includes steels with 7% Ni and 5% Ni, utilizing Cr in place of Ni [10,11]; the other is an innovative nickel-free material—high-manganese austenitic steel specifically designed for low-temperature use.

High-manganese austenitic low-temperature steel, representing the second generation of LNG storage tanks, has attracted considerable interest [12,13,14]. Its cryogenic toughness, thermal expansion rate, and fatigue resistance are on par with those of 9% Ni steel [15,16]. Additionally, it boasts significantly enhanced plasticity, which is roughly threefold that of 9% Ni steel. Furthermore, the cost of metallic Mn is merely one-tenth that of Ni, and the unit price of high-manganese steel is just 46% that of 9% Ni steel. Consequently, this can lead to a reduction in the manufacturing costs of LNG storage tanks to a quarter of the expense associated with 9% Ni storage tanks [17]. For instance, manufacturing a LNG tanker with a total volume of 170,000 m³ can save approximately $14,600,000 in costs. As global LNG storage tanks trend towards larger scales, the thickness required for these tanks is also escalating, sometimes exceeding 100 mm. Embracing the strategy of “substituting Mn for Ni” holds considerable importance for enhancing safety and diminishing construction costs.

Thanks to its face-centered cubic (FCC) crystal structure, high-manganese steel maintains its ductility without undergoing a ductile-to-brittle transition across a broad temperature spectrum, from room temperature (RT) down to liquid nitrogen temperature (LNT, −196 °C) [18,19]. This inherent characteristic ensures the material’s cryogenic toughness to a considerable degree. However, its relatively low yield strength restricts its use in low-temperature environments [20,21,22]. An overemphasis on strengthening could compromise cryogenic toughness, increasing the risk of brittle fractures. Consequently, enhancing the strength without impairing the cryogenic toughness to achieve an optimal balance of strength and ductility has become a focal point in the research and development of high-manganese steel materials. Alloying is recognized as a vital approach to enhancing the mechanical properties of high-manganese steel [23,24,25]. Currently, prevalent industrial techniques for producing and managing alloy steel or micro-alloy steel encompass controlled rolling with heat treatment, regulated cooling, continuous casting at high temperatures with plastic deformation, direct rolling, and immediate quenching. These processes generally involve the material’s deformation, where the control of second-phase particles in steel is attained through mechanisms that induce deformation and phase transformation. Figure 1 exhibits the microstructure of common high-manganese austenitic steel (including second phase precipitates). In the production of medium and thick plates from high-manganese steel for storage tanks, a combination of hot rolling and subsequent heat treatment is commonly employed to control the resulting microstructure and properties (as shown in Figure 1). Upon undergoing hot processing, alloying elements display unique action mechanisms within high-manganese austenitic steels at low temperatures [26,27].

It is crucial to elucidate the laws governing the impact of alloying elements and processes on the strength and cryogenic toughness, which is essential for the alloy design of high-manganese steels intended for low-temperature applications. Consequently, this paper reviews the current knowledge and advances in research concerning the impacts of alloying elements on the strength and toughness of high-manganese steel. The goal was to offer insights that may guide future alloy design and the industrial use of high-manganese steel in LNG storage tanks.

2. Deformation Mechanism of High-Manganese Steel

The superior characteristics of high-manganese steel are largely due to the twinning-induced plasticity (TWIP) effect and the transformation-induced plasticity (TRIP) effect, as referenced in studies [29,30,31]. The TWIP effect primarily occurs through the induction of deformation at grain boundaries, where the twin formation impedes the dislocation movement. This process significantly increases the material’s ductility and plasticity, as detailed in study [32]. Such an enhancement is key to the material’s ability to endure extreme deformation without failure, contributing to the high-manganese steel’s exceptional properties. The TRIP effect encompasses a phase transformation from austenite to martensite, which generates additional dislocations, thereby elevating the dislocation density. This increase in dislocation density contributes to an enhanced strength of the material, as documented in source [33]. This phase transformation is instrumental in augmenting the strength of the steel without compromising its ductility.

SFE plays a pivotal role in determining the deformation mechanisms of high-manganese steel when subjected to deformation and thermal influences. Adjusting the SFE is a critical method for regulating the mechanical properties of this material, allowing for tailored performance under varying conditions. In the context of high-manganese steel, SFE functions as the minimal driving force necessary for twinning and acts as the critical stress threshold for the martensite phase transformation from austenite, as outlined in reference [34]. Research indicates that under conditions of external stress, different behaviors emerge based on the SFE values. Specifically, a martensitic phase transformation is triggered when the SFE is less than 18 mJ/m² [35]. For an SFE range between 12 mJ/m² and 35 mJ/m², twinning deformation predominates, leading to the extensive creation of deformation twins [36,37,38]. Furthermore, when the SFE exceeds 45 mJ/m², the dominant deformation mechanisms transition to full and partial dislocation slips [37]. This delineation of behaviors based on SFE values underscores its critical role in dictating the steel’s response to stress, influencing both its microstructural evolution and mechanical properties.

Factors influencing SFE encompass the chemical composition, the temperature, and the size of austenite grains, as indicated by references [39,40]. It is noteworthy that the impact of alloying elements on SFE varies significantly [29]. Figure 2a illustrates that Al and Cu markedly elevate the SFE, whereas Cr has a diminishing effect on it. The SFE experiences a gradual rise followed by a decrease with the incremental addition of Si. Moreover, C is highlighted as another significant element that conspicuously raises the SFE. The influence of temperature on the SFE is depicted in Figure 2b, providing a visual representation of this relationship. For a specific chemical composition, a rise in temperature results in an increase in the SFE. Allain et al. [41] utilized a thermodynamic model to compute the SFE of Fe-22Mn-0.6C steel at temperatures of −150 °C, 0 °C, and 420 °C, with corresponding values of 10 mJ/m², 19 mJ/m², and 80 mJ/m², respectively. Consequently, different deformation modes were triggered at these varying temperatures due to the changes in SFE levels. The relationship between grain size and SFE is articulated through a calculation model detailed in references [29,31].

$$\Gamma ={2\rho \mathsf{\Delta }G}^{\gamma \to \epsilon }+{2\rho \sigma }^{\frac{\gamma }{\epsilon }}+{2\rho \mathsf{\Delta }G}_{ex}$$

(1)

where $\rho =\frac{4}{\sqrt{3}}\frac{1}{a^{2}N}$ is the molar surface density along the (111) plane; $a$ is the lattice constant; $N$ is the Avogadro constant; ${\mathsf{\Delta }\mathrm{G}}^{\mathsf{\gamma }\to \mathsf{\epsilon }}$ is the Gibbs free energy of γ-phase transformation from austenite to ε-martensite; ${\mathsf{\sigma }}^{\frac{\mathsf{\gamma }}{\mathsf{\epsilon }}}$ is the interface energy between austenite and ε-martensite; and ${\mathrm{\Delta }G}_{ex}=170.06\mathit{exp}\left(-\frac{d}{18.55}\right)$ is the additional part of the influence of austenite grain size effect on the SFE. It is evident that diminishing the grain size of austenite can lead to an elevation in the SFE, with the effect being notably more substantial for smaller grains. Furthermore, research indicates that once the austenite grain size surpasses 30 μm, the influence of grain size on the SFE becomes negligible [42].

In high-manganese steel, the size of the austenite grain is a critical factor in the nucleation of deformation twins. The interplay between the SFE and the grain size significantly influences the critical stress required for twin formation, a relationship that can be quantified by a specific calculation model [43].

$${\tau }_T=\frac{{\gamma }_{isf}}{b_p}+\frac{G·b_p}{D}$$

(2)

where ${\tau }_T$ is the critical resolved shear twinning stress; ${\gamma }_{isf}$ is the intrinsic SFE; $b_p$ is the Burgers vector of the Shockley partial dislocation.; G is the shear modulus; and D is the grain size. Lee et al. [44] produced austenitic Fe-24Mn-4Cr-0.5C high-manganese steel with grain sizes ranging from 5 to 117 μm by fine-tuning the heat treatment process, and assessed the SFE across various grain sizes. They found that as the grain size decreased from 117 μm to 5 μm, the SFE rose from 24 mJ/m² to 32 mJ/m². Both the experimental findings and calculations revealed that the critical stress required for twin nucleation rises as the grain size decreases. This was attributed to the fact that in fine-grained samples, a high dislocation density, a densely packed dislocation configuration, and the interactions among dislocations hinder the activity of dislocations and their partial movement, which are essential for twin formation, thereby leading to an increase in the twin nucleation stress. Recent studies suggest that when the grain size is reduced to a certain threshold, the grain boundary’s impedance of dislocations diminishes, allowing some dislocations to traverse the grain boundary. In some instances, this can precipitate a shift from a dislocation movement mechanism to a grain boundary slip mechanism. Such a shift culminates in a decline in the material’s strength, a phenomenon termed the inverse Hall–Petch relationship [45].

3. Low Temperature Brittleness of High-Manganese Steel

High-manganese austenitic steel typically exhibits failure through brittle fracture at low temperatures. The principal contributors to intergranular brittleness in these steels at reduced temperatures are the embrittlement from strain-induced martensitic transformation [46,47] and the weakening effect of Mn and impurity elements segregating at the grain boundaries [48,49,50]. (A detailed discussion on element segregation is provided in the subsequent section of this paper.)

Chen et al. [51] determined a correlation between the SFE and the energy absorbed during impact at −196 °C for high-manganese steels with different compositions. They discovered that superior low-temperature impact toughness is typically achieved by maintaining the SFE within the range of 40–44 mJ/m². When the SFE of FCC crystals is low, full dislocations readily decompose into Shockley partial dislocations with a Burgers vector of b = 1/6<112>; consequently, a stacking fault forms between these partial dislocations. The stacking fault region may serve as a two-dimensional precursor for ε-martensite formation. The movement of the stacking fault facilitates the transformation of this two-dimensional precursor into a three-dimensional structure, eventually leading to the development of ε-martensite [52]. While ε-martensite is often viewed as an intermediate phase in the γ→α’ phase transformation, its presence is not mandatory for α’ martensite formation. In numerous instances, austenite will directly transition into α’ martensite without passing through the ε-martensite phase [53,54]. Given its typically low SFE, high-manganese steel is inclined to initially form ε-martensite featuring a discontinuous structure from the γ phase through the mechanism of stacking faults, before converting into α’ martensite.

Research [55,56] has demonstrated that ε-martensite plays a crucial role in determining the cryogenic toughness of high-manganese steel. Additionally, the interactions between ε-martensite, mechanical twins, and grain boundaries are frequently linked to brittle fractures when exposed to low temperatures. Figure 3 depicts an Electron Backscatter Diffraction (EBSD) analysis of Fe-17Mn-0.6C steel, which experienced a transgranular fracture as a result of ε-martensite transformation during tensile testing at −150 °C. Supporting this observation, Kim et al. [57] found that the Charpy V-notch impact energy is inversely related to the ε-martensite volume fraction in austenitic Fe—(0.4, 1.0) C-18Mn steels, as illustrated in Figure 4. Due to its inherent brittleness, the transformed ε-martensite tends to facilitate crack propagation, leading to the brittle fracture characteristics [58].

The prevalence of brittle martensitic phases can be regulated through the modification of grain size and chemical composition, as documented in studies [37,55,59,60,61,62]. Refining the grains augments the aggregate grain boundary area, which in turn impedes the expansion of newly formed martensite within austenitic grains. Consequently, this containment by the grain boundaries results in a reduction of the martensite start temperature (Ms) [63].

As depicted in Figure 5 [64], the formation of ε martensite and its interaction with other ε-martensite and grain boundaries contributes to the reduced impact toughness of the coarse grain Fe-30Mn-0.11C austenitic steel at LNT. Conversely, in fine-grained samples, neither martensite was detected at RT nor LNT, and the fracture surface mainly consisted of dimples of sub-micrometer or greater depth, demonstrating superior impact toughness. Koyama et al. [55] focused on the effect of varying grain sizes on the mechanical properties of Fe-17Mn-0.6C TWIP steels at low temperature. They observed that at −150 °C, brittle fractures manifested in steels with grain sizes ranging from 10 to 37 μm. However, reducing the grain size to 3.5 μm significantly mitigated the occurrence of brittle fractures. In their research on the Fe-15Mn alloy, Takaki et al. [65] discovered that a reduction in grain size nearly halted the phase transformation to α’ martensite and precipitated a steep decline in the content of ε-martensite, particularly within a grain size range below 30 μm. This finding is illustrated in Figure 6.

Takaki et al. [66] put forward the theory of the “Spatial Confinement Effect”, predicated on the strain energy involved in martensitic phase transformation. This theory posits that grain refinement inhibits the transformation to martensite as depicted in Figure 7. As martensite laths form, they impose anisotropic strains on the surrounding austenite that has yet to transform. To counteract these strains, various martensite variants typically crystallize at the same time within a single austenite grain, serving to minimize the cumulative strain energy. Nevertheless, as the austenite grains become more refined, there is a progressive shift in the martensite morphology from a stratified structure to a simplified one. When the austenite grain size reaches a dimension comparable to that of the martensitic laths, the spatial limitations hinder the simultaneous formation of multiple variants. This implies that in ultrafine austenite grains, the minimization of strain energy resulting from the martensitic transformation is no longer achievable. As the transformation shifts from a multi-variant to a single-variant mode, the austenitic phase can be obviously stabilized. In scenarios where multivariate transformation is challenging, the martensitic transformation in metastable austenite will be suppressed.

4. Effect of Alloying Elements on High-Manganese Steel

Since the 21st century began, intensive research and development have been conducted globally on high-manganese austenitic low-temperature steel. Various standards have been established by organizations such as the China Special Steel Enterprise Association, China Classification Society, ISO International Standards Organization, and the American Society for Testing and Materials, each imposing slightly divergent prerequisites for its composition, as illustrated in

Table 1. Chemical composition requirements of high-manganese austenitic cryogenic steel (wt.%). Reprinted from Ref. [67]. Data from [68,69,70,71].

Standard No.	Standards Organizations	C	Si	Mn	P	S	Cr	Cu	B	N
T/SSEA 0060―2020	Special Steel Enterprises Association of China	0.35–0.55	0.10–0.50	22.50–25.50	≤0.012	≤0.005	3.00–4.00	0.30–0.70	≤0.005	≤0.050
GD28―2021	China Classification Society	0.35–0.55	0.10–0.50	22.50–25.50	≤0.030	≤0.010	3.00–4.00	0.30–0.70	≤0.005	≤0.050
ISO 21635:2018	ISO International Standards Organization	0.35–0.55	0.10–0.50	22.50–25.50	≤0.030	≤0.010	3.00–4.00	0.30–0.70	≤0.005	≤0.050
A1106M―17	American Society for Testing and Materials	0.35–0.55	0.10–0.50	22.50–25.50	≤0.030	≤0.010	3.00–4.00	0.30–0.70	≤0.005	≤0.050

[67]. Notably, the Chinese Special Steel Enterprise Association has the strictest requirements for the P and S content in high-manganese austenitic low-temperature steel, while the other elements’ content aligns across the four standards. These standards’ establishment and implementation have provided essential reference points for high-manganese austenitic low-temperature steel utilization, facilitating advancements and progress in this field.

The four standards exhibit relatively broad requirements for the element content range in high-manganese austenitic low-temperature steel. According to existing literature, the mechanical properties distribution of high-manganese austenitic low-temperature steel can be summarized from the perspectives of RT yield strength, RT tensile strength, and LNT Charpy impact energy, as delineated in Figure 8 [67], where the red dashed line represents the benchmark performance criteria (yield strength ≥ 400 MPa; Tensile strength: 800–970; LNT Charpy impact energy ≥ 80 J). Considering the fulfillment of the four standard composition and performance requirements, as well as the cost-effectiveness of alloying, the Fe-24Mn-0.45C-3Cr high-manganese austenitic low-temperature steel appears to be the optimal case material and can serve as a crucial reference for the composition optimization design of high-manganese austenitic low-temperature steel for LNG storage tanks.

Beyond the realm of chemical composition, the prowess of high-manganese austenitic steel is significantly swayed by the nuances of thermal processing methodologies. The four prevailing standards require high-manganese austenitic low-temperature steel plates to be delivered using the hot-rolling followed by the controlled-cooling process. Moreover, a number of researchers have delved into the effects of the various heat treatments [28,72,73,74,75] and welding procedures [76,77,78] on the intrinsic characteristics of this steel variety. While various processing techniques have been employed by different scholars in their experiments, this treatise shall principally concentrate on unraveling the intricate interplay between the alloy composition and the performance attributes of high-manganese austenitic steel tailored for low-temperature environments.

Figure 8. Properties distribution of high-manganese steel austenitic low-temperature steel (a) RT yield strength, (b) RT tensile strength, and (c) LNT Charpy impact energy. Reprinted from Ref. [67]. Data from [68,69,70,71,79].

Figure 8. Properties distribution of high-manganese steel austenitic low-temperature steel (a) RT yield strength, (b) RT tensile strength, and (c) LNT Charpy impact energy. Reprinted from Ref. [67]. Data from [68,69,70,71,79].

4.1. Mn

Mn, the key component in high-manganese steel, is instrumental in broadening the austenite phase region, lowering the martensite start (Ms) temperature, and enhancing the stability of the austenite phase. In the study by Bhattacharya et al. [80], the findings indicated that the primary factor influencing the impact toughness is the stable volume fraction of residual austenite. However, the presence of a higher Mn content in steel, coupled with its slower diffusion rate, leads to the material’s increased fraction of residual austenite.

Zhu et al. [81,82,83] introduced the Fe-Mn-C atom cluster-strengthening mechanism, underpinned by the theory of short-range ordered segregation, through both theoretical calculations and electronic probe detection. They observed a heterogeneous distribution of C and Mn atoms within the high-manganese steel matrix, leading to the formation of C-Mn crystalline cells. The Mn and Fe atoms imposed a significant constraint on the C atoms, effectively restricting their movement along the dislocations and thereby creating a substantial drag effect on these dislocations as they moved [84]. When the dislocations crossed paths with the areas of Fe-Mn-C atomic cluster segregation, many C-Mn bonds were severed. This event caused atoms adjacent to the slip plane to reorganize, leading to a decrease in short-range order. The aforementioned structural modifications led to a localized reinforcement within the material, which elevated its capacity for work hardening throughout the deformation process. Consequently, it can be inferred that fine-tuning the levels of C and Mn elements in the alloy can effectively augment the strength of high-manganese steel. Zhao et al. [85] explored the tensile deformation characteristics of high-carbon, high-manganese steels with Mn concentrations varying between 7% and 19%. Their findings indicated that an increase in the Mn content not only simultaneously improved the strength and ductility, but also the work-hardening capacity of high-manganese steel. This enhancement was attributed to the synergistic impact of Mn on promoting deformation twinning and facilitating dynamic strain aging (DSA) [86,87].

The influence of Mn on SFE is also noteworthy. Prior research has shown [30] that SFE decreases with an increase in the Mn content, within a specific range, as depicted in Figure 9. Nevertheless, within the typical content range for high-manganese steel, SFE generally rises with greater Mn content, mitigating the low-temperature brittleness associated with martensitic phase transformation [88]. Lee et al. [16] conducted research on how varying amounts of Mn affect the mechanical properties of Fe-(17–22) Mn-0.45C-2Al steels at low temperatures. Their study revealed that 22Mn steel, where no traces of martensite could be found, exhibited stronger low-temperature tensile and impact properties compared to the presence of α’-martensite 17Mn and 19Mn. Kim et al. [89] performed dynamic compression deformation tests on austenitic Fe-(22–26) Mn-0.4C steel with 22%, 24%, and 26% Mn content. They observed that at low temperatures, the SFE of the 22Mn steel exhibited a 30% reduction from its RT value, which induced martensitic phase transformation. As the Mn content increased, a twinning-induced mechanism became apparent, which enhanced the low-temperature toughness of the 24Mn and 26Mn steels in comparison to the 22Mn steel. Sohn et al. [90] investigated the mechanical properties of four high-manganese steels, Fe-(19, 22) Mn and Fe-(19, 22) Mn-2Al, at both room and low temperatures. Their findings indicated that with a decrease in the testing temperature, the impact absorption energy of the 19Mn and 22Mn steels diminished significantly. Under the same conditions, the 19Mn2Al and 22Mn2Al steels showed only a slight reduction. Figure 10 presents the EBSD phase maps of the cross-sectional area at the fracture initiation region from low-temperature Charpy impact test specimens. It illustrates that the deficiency in cryogenic toughness can be ascribed to the abundant presence of ε or α’-martensite within the fracture initiation zone. These martensites act as initiation points for cracks and facilitate the rapid propagation of cracks, culminating in a quasi-cleavage fracture mode on the surface of Charpy impact cracks. Conversely, an increase in Mn and Al content raises the SFE, which promotes the formation of numerous deformation twins in the 19Mn2Al steel. Once these deformation twins have fully developed, the α’-martensite transformation occurs at a reduced rate. Owing to the scarcity of martensite as potential crack initiation sites, this results in a relatively high impact toughness at low temperatures.

Additionally, there are concerns regarding the negative impact that a high concentration of Mn may have on the properties of high-manganese steel. When the manganese content goes beyond 30 wt.%, the brittle β-Mn phase may be formed, resulting in a significant reduction in ductility while increasing the likelihood of brittle fractures in high-manganese steels [91,92]. The traditional viewpoint on grain boundary weakening posits that the inclusion of Mn augments the SFE, consequently elevating the energy barrier for dislocation movement. This increased energy requirement hampers the dislocation mobility within the crystal lattice, leading to heightened stress concentration at dislocation sites. Such persistent stress concentrations within the grains can set the stage for the onset of a brittle fracture [93].

At present, the most debated theory concerns the brittleness at low temperatures that results from the segregation of Mn and impurity elements at grain boundaries. Strum et al. [94] mentioned that the addition of Mn diminishes the cohesion at grain boundaries, thereby facilitating brittle fractures. Nonetheless, this perspective is not strongly supported by experimental evidence. Some researchers maintain that Mn intensifies the adverse impact of impurities on the material toughness [95], while others assert that the segregation of Mn at grain boundaries is responsible for increased brittleness [96]. The concept of Mn segregation at grain boundaries was first introduced by Morris in 1986 [97]. Subsequently, in 1988, Xue [48] and Chai [50] employed Auger spectroscopy to uncover that Fe-Mn binary alloys containing more than 30 wt.% Mn were prone to brittle fractures at low temperatures, a phenomenon attributed to disproportionate Mn segregation at the grain boundaries. In 1988, Strum and Morris [94] challenged the previously established idea of Mn segregation at grain boundaries. They noted that in the case of mixed fractures in 31Mn-5Cr-0.16N steel, the Mn content was comparable across surfaces of both intergranular and transgranular fracture zones. They proposed that what had been perceived as Mn segregation could actually be an artifact resulting from selective sputtering during the Ar+ sputtering process. Xue and colleagues [98] later cast doubt on this rebuttal, proposing that the suggested phenomenon might not occur in Fe-Mn alloys at all. Following this, Herbig et al. [49] employed atom probe tomography (APT) to examine the local chemical composition at grain boundaries within Fe-28Mn-0.3C series TWIP steel. Their findings indicated an absence of Mn enrichment or depletion at the interfaces. Thus, it would be erroneous to exclusively ascribe the intergranular fracture of high-manganese austenitic steel at low temperatures to Mn segregation at the grain boundary. The influence of Mn or impurity segregation on the cryogenic toughness of high-manganese steel is neither unequivocal nor singular. There may be other contributing factors at work, and occasionally it could be a combination of multiple mechanisms interacting. Consequently, more comprehensive research is essential to thoroughly unravel this complex issue.

4.2. Al

Al possesses an FCC crystal structure, akin to that of austenite. Introducing Al can serve as an exogenous nucleation agent, which results in the refinement of grains and an improvement in the strength of the material [99]. Jung et al. [100] conducted a study on the tensile mechanical properties of Fe–18Mn–0.6C–(0–2.5)Al TWIP steels with varying levels of Al content. Their findings indicated that adding Al to austenitic steel enhances its yield strength and elongation, albeit at the cost of its work-hardening capability. The improvement in yield strength originates from the distortion of the crystal lattice attributed to the dissolution of the Al element in austenite. Conversely, the diminished work-hardening ability stems from an increase in SFE, elevating the critical stress required for twin nucleation and thereby favoring dislocation slip as the primary mode of deformation. Moreover, Al plays a substantial role in markedly reducing the DSA effect. This is primarily because Al raises the diffusion activation energy of C, thereby shortening the interaction duration between stacking faults and the C-Mn atoms.

As highlighted earlier, incorporating Al into high-manganese steel can elevate the SFE [101], suppress the formation of ε-martensite, and promote the stability of single-phase austenite. Sohn et al. [84] revealed that incorporating 2.0 wt.% Al into 19Mn steel resulted in an approximate 14 J increase in the absorbed energy measured by the Charpy V-notch test at −196 °C. In a similar vein, the addition of Al to 22Mn steel yielded an improvement of around 11 J in the same impact toughness test. However, Chen et al. [102] observed a notable reduction in the impact energy absorbed at low temperatures when the Al content surpassed 5.0 wt.%. As depicted in Figure 11, the introduction of Al allowed 3Al and 5Al steels to achieve fully austenitic microstructures and develop high local orientation gradients in the vicinity of the principal crack. This suggests that during crack propagation, there was ample plastic deformation. In contrast, the crack propagation path in 8Al steel was relatively linear, which implies less effective crack inhibition. An overly high concentration of Al has been shown to diminish the austenite phase region while subsequently increasing the ferrite phase region [103]. δ ferrite, possessing a body-centered cubic (BCC) structure, demonstrates a transition from ductility to brittleness when subjected to low temperatures [104]. Figure 11e,f showcase a proclivity for microcracks to initiate at the γ/δ interfaces, spreading either along these interfaces or throughout the δ ferrite colonies. This behavior is largely due to the elevated stress concentrations present in these areas.

Moreover, an excess of Al has the potential to combine with nitrogen to form aluminum nitride (AlN), which is soluble in austenite at elevated temperatures. As the temperature declines, AlN precipitates and predominantly accumulates along grain boundaries. Such deposition is inclined to cause grain boundary embrittlement, which in turn can lead to a decrease in material toughness [105]. The process of pouring can readily result in the formation of oxide residues, which have a tendency to clog the sprue gate, posing challenges to the industrial production of high-manganese steel [106]. In conclusion, it is advisable to maintain the Al content in high-manganese steel at a relatively low level to mitigate these issues.

4.3. Cr

Incorporating Cr into the austenitic structure leads to lattice distortion, which contributes to solid solution strengthening and consequently enhances the yield strength of the material [23]. Similar to the interaction between Mn atoms, Cr can pair with C to form C-Cr atomic pairs, and the repulsion between Cr and manganese atoms is notably weak, which permits the formation of Cr-C-Mn atomic clusters [21]. Additionally, Cr exhibits a strong affinity for nitrogen atoms. Chen et al. [107] applied Cr and nitrogen alloying to traditional Fe-12Mn-1.1C high-carbon high-manganese steel and conducted a systematic set of comparative experiments. The results from tensile tests revealed a marked enhancement in yield strength, along with a modest improvement in elongation following the Cr + N alloying treatment. Under the same strain, Fe-12Mn-1.1C-2.14Cr-0.05N high-carbon high-manganese steel exhibited a higher incidence of twinning. Moreover, in low-cycle fatigue tests, this Cr + N alloyed steel developed twins at high strain amplitudes (0.8%), which substantially extended its fatigue life. Additionally, the incorporation of Cr elements has been shown to improve hardenability and corrosion resistance [108]. Yuan et al. [109] found that in Fe-26Mn-0.33C-2.8Al-(1.13–3.95)Cr-0.01N TWIP steels, the proportion of low-angle grain boundaries (2°–10°) rose with an increase in Cr content. By enhancing the stability of low-angle grain boundaries, which occur before the corrosion of high-angle grain boundaries with higher mismatch energy, the overall stability of the grain boundaries is improved, and this stabilization contributes to an augmented corrosion resistance overall.

Yuan et al. [109] observed that as the Cr content rises, the grain size of Fe-26Mn-0.33C-2.8Al-(1.13–3.95)Cr-0.01N TWIP steels post-annealing diminishes. This effect is attributed to the formation of Cr₂₃C₆ carbide during the annealing process. This carbide, characterized by its high melting point, displays minimal solubility in the annealing phase, which serves to obstruct grain boundary movement and consequently restricts the enlargement of grains [110]. At the same time, carbides can impede the slip of dislocations during plastic deformation, thereby contributing to an enhancement in the yield strength. Nonetheless, the presence of M₂₃C₆ or M₇C₃ carbides may adversely affect the fracture toughness, with the extent of this impact contingent upon their size, content, and distribution within the material [23,111].

In general, grain boundary segregation tends to become more pronounced at relatively low temperatures. However, when Kwon et al. [112] studied the fracture behavior of 8Mn steel under cold conditions, they found that intergranular fractures are predominantly a consequence of the matrix being stronger than the grain boundaries, a finding that deviates from the traditional view that attributes such brittleness to segregation. Chen et al. [113] reported that substantial grain boundary segregation along with a high concentration of (Cr, Mn)₂₃C₆ carbides leads to weakened atomic bonds at these boundaries, thus increasing their susceptibility to crack formation. Moreover, the heightened critical stress necessary for secondary twinning reduces the number of twins that form in the secondary twinning system, translating to a diminished ability of the material to deform under conditions of low temperature and dynamic stress. As mentioned earlier, the movement of Shockley partial dislocations leads to the creation of several isolated intrinsic stacking faults, which are also the basis foundation for the nucleation of twins [114]. Several studies suggest that when subjected to shear stress, a perfect dislocation may split into leading and trailing partial dislocations, with the emergence of stacking faults occurring between them. The presence of stacking faults triggers the movement of the partial dislocations. In particular, the stacking fault pulls the trailing partial dislocation forward and pushes the leading partial dislocation backward, and the separation distance between the partial dislocations is closely linked to the SFE [115]. Clearly, the nucleation of twins is closely connected to dislocation decomposition and stacking faults at grain boundaries, and equilibrium segregation at these grain boundaries significantly influences the toughness by affecting twin nucleation. Lee et al. [23] found that adding Cr to high-manganese steel impedes deformation twinning, which reduces the material’s toughness at 25 °C. Conversely, this addition actually enhances the steel’s low-temperature impact toughness at −196 °C due to its inhibitory effect on ε-martensite formation. Such results open up encouraging prospects for the utilization of high-manganese steel at low temperatures.

4.4. C

C is the primary element impeding a dislocation slip by acting as a pinning agent. Its atomic radius is smaller than that of Fe and Mn, which allows it to accumulate in the regions of compressive stress surrounding dislocations. This accumulation leads to the formation of a Cottrell atmosphere as a result of interactions with dislocations. Such interactions heighten the resistance to dislocation movement, thereby enhancing the material’s strength while simultaneously yielding a modest decrease in plasticity [116]. Choi et al. [117] demonstrated that in Fe-15Mn-2Al-1Si-xC TWIP steels, the solid solution strengthening effect yields an increase of 91.3 MPa in strength for each 1 wt.% increment in C content. Kusakin et al. [93] carried out analogous research and found that in Fe-18Mn-(0.3–0.6)C-1.5Al TWIP steels, the solid solution strengthening effect of C was quantified at approximately 250 MPa for each weight percent of C. They posited that this strengthening effect might be associated with the creation of octahedral clusters containing Mn around the C atoms.

A deficit of C hinders the formation of a singular austenite phase in high-manganese steel, whereas an excessive quantity results in the formation of a grain boundary network of cementite or the precipitation of large carbides, which predisposes the high-manganese steel to brittle fractures. Koyama et al. [118] investigated Fe-33Mn with varying C content ranging from 0 to 1.1% and observed that both the strength and elongation of the material improved as the C content increased, at a deformation rate of 10⁻² s⁻¹. However, when the tensile rate was lowered to 10⁻⁵ s⁻¹, a DSA effect was noted in materials with a higher C content, which led to the premature fracture of Fe-33Mn-1.1C. They also noted the potential for blue embrittlement in steels with an excessively high C content.

It is widely recognized that the precipitation of carbides adversely affects cryogenic toughness due to the resultant stress concentrations. This condition facilitates the easy nucleation of microcracks at the interface between the carbide and the matrix. Nevertheless, the impact of carbide precipitation on toughness is also contingent upon their size, content, and distribution. As illustrated in Figure 12, these vanadium carbides (VCs) are relatively small, approximately 20 nm in diameter, and their lattice parameters closely resemble those of the austenitic matrix, permitting a certain level of lattice coherence. This compatibility fosters a more even distribution of strain across both the carbide particles and the matrix. As a result, the elastic interaction between the dislocations involved in cutting and the carbides is diminished, which in turn lowers the strain accumulation and aids in mitigating crack initiation [119].

In Ren’s research [120], it was found that fine, nanoscale VCs could enhance the strength of Fe-23.7Mn-0.8 C high-manganese steel while preserving high toughness. The introduction of a high C content fosters recrystallization and recovery processes, resulting in fewer deformed grains in the hot-rolled microstructure. This consequently aids the glide of dislocations and encourages the creation of twins. Moreover, higher levels of C increase the SFE, favoring twinning and a dislocation slip as the main mechanisms of deformation, while simultaneously preventing the martensite phase transformation within this context. It is noteworthy that carbides typically hinder twinning growth. However, Yen et al. [121] suggested that significant interactions between twins and carbides were the main manifestation when the twin’s thickness is considerably less than that of the carbide within Fe-21.6Mn-0.63C-0.87V steel. Consequently, in the context of nanoscale VCs, there is no substantial impediment to the propagation of twinning. This allows for fully activated twinning to trigger the dynamic Hall–Petch effect during impact deformation, thereby enhancing the capacity for plastic deformation.

The combined alloying approach using C and N applied to high-manganese steel also demonstrates marked enhancements in strength. More precisely, this C + N composite treatment elevates the free electron density in the steel, aids in the short-range atomic ordering within the austenite structure, and encourages a homogeneous distribution of solute atoms like Cr and Mn throughout the austenite [122,123,124]. Kang et al. [125] engineered Fe-18Mn-(0.6–0.8) C-(0.2–0.3) N high-carbon high-manganese steels with a combined total of 0.9% C and N. This composition, following the C + N composite alloying treatment, displayed an abundance of ultrafine nano-twins during plastic deformation. The presence of these nano-twins contributed to enhanced strength and ductility and also resulted in a lowered ductile-to-brittle transition temperature.

5. Effect of Microalloying Elements on High-Manganese Steel

Currently, the microalloying elements primarily added to high-manganese steel include V, Nb, and Ti. These elements readily undergo chemical interactions with C or N to form carbides or nitrides. Given that these two compounds share the same crystal structure and have comparable lattice constants, they can dissolve in each other, typically forming either a solid solution or carbonitride precipitates. The finely dispersed precipitates act as formidable barriers to dislocation movement, thus providing a considerable precipitation strengthening effect. Furthermore, these precipitates can hinder the growth of austenite grains and inhibit austenite recrystallization. They also modify the microstructural parameters of the material following phase transformation, which in turn affects the material’s mechanical properties.

The solubility of most carbonitrides in ferrite or austenite can be captured by the solid solubility product equation. For austenite, the equations representing the solid solubility products for compounds such as VC, VN, NbC, NbN, TiC, and TiN are detailed in references [126,127]:

$$\mathrm{l}\mathrm{g}\left[\mathrm{V}\right]\left[\mathrm{C}\right]=6.72-9500/T$$

(3)

$$\mathrm{l}\mathrm{g}\left[\mathrm{V}\right]\left[\mathrm{N}\right]=3.40-8330/T$$

(4)

$$\mathrm{l}\mathrm{g}\left[\mathrm{N}\mathrm{b}\right]\left[\mathrm{C}\right]=2.96-7510/T$$

(5)

$$\mathrm{l}\mathrm{g}\left[\mathrm{N}\mathrm{b}\right]\left[\mathrm{N}\right]=2.89-8500/T$$

(6)

$$\mathrm{l}\mathrm{g}\left[\mathrm{T}\mathrm{i}\right]\left[\mathrm{C}\right]=2.75-7000/T$$

(7)

$$\mathrm{l}\mathrm{g}\left[\mathrm{T}\mathrm{i}\right]\left[\mathrm{N}\right]=0.32-8000/T$$

(8)

where [V], [Nb], [Ti] are the mass percentage of the part in solid solution, %.

M (C_xN_1−x) = x [M][C] + (1 − x) [M][N]

(9)

Considering the existence of thermodynamic dynamic equilibrium reaction Equation (9), where M is the alloying element and x is the mole fraction of MC in the carbonitrides, combined with the solid solubility product equation it can be concluded that the solid solubility product of vanadium carbonitride is the largest, being two to three orders of magnitude greater than that of the other compounds [127]. Nb and Ti require extremely high austenitizing temperatures to achieve complete solid solubility; however, excessively high temperatures can lead to severe grain coarsening. Conversely, the precipitation of V carbonitride is much more manageable during thermo-mechanical controlled processing (TMCP), which is why it is the most commonly utilized microalloying element in microalloyed steels. Scott et al. [128] investigated the impact of various microalloying elements on the mechanical properties of TWIP steels. Their comparative analysis, illustrated in Figure 13, revealed that Ti provided the most significant strengthening effect when the concentration of the alloying element was below 0.1 wt.%, as detailed in

Table 2. The strengthening coefficients of microalloying additions of Ti, V, and Nb below 0.1 wt.%. Reproduced with permission from ref. [128]. Copyright 2011 De Gruyter.

Microalloying Element	Strengthening Coefficient for Microalloying Additions <0.1 wt.%
Nb	187 MPa/wt.%
V	>530 MPa/wt.%
Ti	1380 MPa/wt.%

. The strengthening effect reaches a plateau when the addition of the alloying element surpasses 0.1 wt.%; with Nb microalloying, this leads to the formation of larger NbC particles, which in turn yields a diminished enhancement in strength and a reduction in plasticity. In contrast, V imparts a moderate strengthening effect without exhibiting the saturation plateau characteristic of Ti. The yield strength enhancement in two TWIP steels with distinct matrices, after the introduction of V as a microalloying element, was represented by fitting curves that were strikingly similar.

5.1. V

V exhibits an exceptionally high solubility in austenite, achieving nearly complete dissolution. Upon cooling, VN precipitates first, eventually transitioning into V(C,N). The presence of fine grains at the nano/micron-scale not only enhances the material’s strength but also improves its resistance to delayed fracture and stress corrosion cracking performance [129,130]. V carbonitrides are smaller in size compared to microalloyed elements such as Nb and Ti. Gladman’s theory posits that the precipitation strengthening effect is related to the size and volume fraction of the precipitated particles, which can be expressed by the subsequent formula [131]:

$$∆{\mathrm{S}}_{\mathrm{p}}=0.538Gb\left(\frac{f^{-\frac{1}{2}}}{d}\right)\mathrm{I}\mathrm{n}\left(\frac{d}{2b}\right)$$

(10)

where ∆Sp is the effect of precipitation relative strength; G is the shear modulus, MPa; b is the Burgers vector, nm; f is the volume fraction of precipitated particles; and d is the average diameter of precipitated particles, nm. It is evident that the precipitation strengthening effect is influenced by both the size of particles and their volume fraction; smaller particles and larger volume fractions lead to more significant strengthening effects. Furthermore, the size of the precipitated particles plays a crucial role in the formation of twins in high-manganese steel, exerting a substantial impact on its development. While Dumay et al. [29] observed no interaction between twin boundaries and V(C,N) precipitates in Fe-22Mn-0.6C-0.2V steel, even with 10% tensile deformation, it should be noted that this absence of a significant interaction only pertains to the behavior under low-strain conditions.

Scott et al. [128] demonstrated that at low strains (ɛ < 0.25), precipitated particles from microalloying elements did not impact the work hardening of high-manganese steel. Conversely, at higher strains (ɛ > 0.3), they observed a decrease in the strain-hardening rate, particularly in alloys with a higher volume of precipitates. This phenomenon was attributed to the dispersed particles inhibiting the growth of twins. Zhang et al. [129] substantiated these observations by comparison of Fe-22Mn-0.6-(0–0.19)V steels. They encountered two different scenarios when deformation twins came into contact with precipitated VC particles. In the first case, the deformation twins were able to traverse smaller VC particles, a process illustrated in Figure 14a. However, when the size of the VC particles exceeded 50 nm in diameter, a notable impediment to the growth of the twins was evident, as depicted in Figure 14b. This is consistent with the theory proposed by Yen et al. [121], which suggests that the interaction between deformation twins and VC carbides is influenced by the size of the particles. Specifically, they demonstrated that the interaction intensifies when the twin size is smaller than that of the carbides, and conversely, it diminishes when the twin size is larger. Figure 14 [120] also provides evidence that nanoscale carbides do not obstruct the growth of twins, yet they enhance the yield strength without detracting from cryogenic toughness. This observation serves as a pivotal point and guides future research endeavors aimed at optimizing microalloyed high-manganese steels for use in low-temperature conditions.

5.2. Nb

Besides hindering recrystallization through the pinning of grain boundaries and dislocations with carbonitrides, Nb can also induce a drag effect as solid solution atoms within the high-temperature austenite region. This drag effect contributes to the delayed onset of recrystallization, a phenomenon that is directly linked to the difference in radius between the solid-solution atoms and Fe atoms. The larger this difference, the more pronounced the inhibitory effect on recrystallization. Therefore, Nb effectively prevents austenite recovery, showcasing the strongest recrystallization inhibition capability. In contrast, V and Ti, having smaller radius differences with Fe atoms, introduce lesser lattice distortions and lower grain boundary energies, resulting in relatively weaker inhibitory effects on recrystallization.

Yu et al. [132] summarized the inhibitory effects of various alloying elements on austenite recrystallization by examining the relationship between the solute atom concentration and the onset time of recrystallization. Their findings indicated that Nb’s inhibitory effect on austenite recrystallization is significantly greater than that of V and Ti. When the solid solution content of Nb is below 0.020 wt.%, even trace amounts of Nb notably raise the recrystallization cessation temperature of austenite, which can reach up to 950 °C. As the Nb content increases from 0.020 to 0.050 wt.%, the effect on retarding recrystallization becomes more gradual, yet it remains considerably more substantial than the effects induced by V and Ti.

The precise mechanism by which Nb enhances the cryogenic toughness of high-manganese steels has not been fully elucidated. Huang et al. [133] proposed that the addition of Nb might elevate the SFE, thereby suppressing the martensitic phase transformation. Conversely, Kwon et al. [134] observed that Nb addition led to a reduction in twin dynamics within Fe-17Mn-0.6C-0.05NbTWIP steel, a phenomenon they attributed to the refinement of grains. Cao et al. [135] offered an alternative perspective. Through a comparative analysis of the microstructures of Fe-15Mn-0.9C and Fe-15Mn-0.9C-0.012Nb TWIP steel after pre-deformation (e.g., Figure 15), they noted that the Nb-enriched TWIP steel displayed a higher stacking fault density and finer and more evenly distributed spacing, along with a clear presence of high-density dislocations interspersed among the stacking faults. Figure 15c showcases an intricate composite structure comprising twins, dislocations, and lamellae, where T1 and T2 represent two distinct sets of twins, each varying in orientation and spacing. This suggests that as the deformation process progresses, the extent of deformation prompts the emergence of new twins with differing orientations amidst the pre-existing ones. The stress concentrations originating from the initial twins serve as a catalyst, providing the necessary driving force for the genesis of secondary twins. The synergistic activity of these twin systems significantly amplifies the TWIP effect, consequently augmenting the material’s plasticity. Variations in the impact on twinning behavior are likely attributable to the differing levels of Nb added to the alloy.

Chen et al. [136] examined the influence of varying amounts of Nb on the structural properties of Fe-24Mn-0.4C-(0–0.29)Nb high-manganese austenitic steels. They discovered that the introduction of trace amounts of Nb to the steel resulted in a negligible enhancement of the yield strength. The lack of effectiveness in precipitation hardening was linked to the substantial size of the NbC carbides, which did not provide the anticipated hardening effect. Moreover, the contribution of solid solution strengthening was limited, resulting from the relatively low quantities of Nb added. Nevertheless, a notable improvement was observed once the Nb addition was raised to 0.29 wt.%. This increase led to a significant rise in the number of grains exhibiting high local orientation gradients within the grain. Additionally, the presence of intense dislocation tangles and a lath recovery microstructure were noted, thereby substantially enhancing the yield strength. Simultaneously, the inhibition of recrystallization and the presence of relatively large grains favored the formation of twins, which positively influenced the cryogenic toughness. However, this benefit was countered by the larger radius and accelerated growth rate of NbC carbide. The advantageous impacts of twinning were overshadowed by the detrimental effects of the large-sized carbides, leading to a consistent decline in the material’s low-temperature impact energy.

5.3. Ti

Ti has the capacity to form stable nitrides or carbonitrides at elevated temperatures. Owing to its lower solid solubility, titanium nitride (TiN) is nearly insoluble in austenite. This characteristic enables TiN to impede the growth of austenite grains during the reheating phase prior to hot working. At lower temperatures, the strain-induced precipitation of finely dispersed TiC particles acts as a barrier to grain growth, thereby enhancing the strength and ductility of high-manganese steel. As the temperature rises, these TiC particles dissolve, leading to an increase in grain size. However, upon reaching 1250 °C, the precipitated TiN particles become effective at pinning the grain boundaries, thereby obstructing the growth of austenite grains [137].

Han et al. [138] conducted a study to assess the impact of Ti on the thermoplasticity and hot-cracking susceptibility of Fe-18Mn-0.65C-(0.01–1.4)Ti high-manganese steels, along with their performance at low temperatures. The research indicated that incorporating Ti enhanced the steel’s thermoplasticity, but diminished its susceptibility to hot cracking and severely degraded its low-temperature impact toughness. This decrease was attributed to the uneven distribution of TiC and TiN precipitates throughout the matrix. Wu et al. [139] delved into the causes behind the diminished cryogenic toughness of 0.059 wt.% Ti microalloyed steels. They discovered that the introduction of Ti shifted the fracture mechanism of the material to cleavage fracture. This transition was confirmed to be due to the presence of large-sized composite inclusions and a plentiful concentration of micron-level TiN.

Liu et al. [140] performed SEM observations on specimens exhibiting reduced toughness following 0.054 wt.% Ti microalloying, as depicted in Figure 16a,b. These images revealed larger-sized TiN particles situated along the crack paths on the smooth cleavage planes. Moreover, a number of cleavage cracks were observed initiating at the edges of these TiN particles, as detailed in Figure 16c,d. The fracture process in the impact specimen unfolds through the following sequence: Initial nucleation of sharp microcracks occurs at the large-size TiN particle tips. These microcracks then propagate from the interface between the inclusions and the matrix, progressively advancing into the matrix interior. The final stage involves further expansion, which proceeds either through or along the matrix boundaries. It has been calculated that the critical size for microcrack propagation control is 4.93 μm. Beyond this threshold, the TiN particles interrupt the metallic continuity and induce a stress concentration, which in turn generates microcracks in their vicinity. These microcracks then merge, grow, and extend, culminating in the material’s fracture and the consequent degradation of its toughness. To prevent the growth of TiN particles in steel, which can adversely affect the material’s properties, it is crucial to meticulously regulate the ratio of Ti to N. This ensures a minimal coarsening rate of TiN inclusions during the homogenization process prior to rolling [141,142]. Additionally, it is important that the product of the total Ti and N content remains below the solubility product of TiN at the solidus line temperature. This precaution is necessary to inhibit the precipitation of TiN particles prior to solidification [143].

6. Conclusions

High-manganese austenitic steel, as a new material for LNG storage tanks, has garnered significant interest for its outstanding performance and cost efficiency. However, its application is constrained by its low yield strength, and attempts to strengthen it often result in reduced low-temperature impact toughness. This paper outlined the deformation mechanism and the principles behind the low-temperature brittleness of high-manganese steel. It further explored how alloying elements impact high-manganese steel, with the objective of providing a theoretical foundation for the design of alloyed high-manganese steel for LNG tanks. The goal was to ensure future designs achieve both exceptional strength and toughness.

Based on the existing standards for high-manganese austenitic low temperature steels in terms of composition and property requirements, and considering the cost effectiveness of alloying, the Fe-24Mn-0.45C-3Cr system of steels appears to be the optimal case material and can be a crucial reference for the composition optimization design of high-manganese austenitic low temperature steels.

Solid solution elements like Al, Cr, and C, along with microalloying elements, play a pivotal role in solid solution strengthening and grain refinement. Mn and Cr atoms have a propensity to form atomic pairs with C atoms, exerting a significant drag effect on dislocation movement, thus enhancing the steel’s strength. Moreover, microalloying elements, such as V, Nb, and Ti, upon interacting with C and N within the steel, form carbides and nitrides that significantly improve the yield strength. When added in standard amounts (less than 0.3 weight%), these elements strengthen the steel in the order of Ti > V > Nb. However, the impact of Ti levels off after reaching a concentration of 0.1 weight percent.

The low-temperature performance of high-manganese austenitic steel is intimately linked to the presence of martensite laths and deformation twins within its structure, factors largely governed by the steel’s SFE. To achieve optimal low-temperature impact toughness, the SFE at RT should range between 40 to 44 mJ/m². The incorporation of Mn and Al can effectively suppress the transformation into the martensitic phase and thus enhance the steel’s low-temperature toughness. Nonetheless, an excessive amount of Al can hinder the formation of deformation twins and expand the ferrite phase region, which could ultimately lead to a reduction in low-temperature toughness. Nb, Ti, and Cr share the characteristic of forming large-sized carbides, which act as nucleation sites for cracks and inhibit the formation of deformation twins, negatively affecting the material’s low-temperature toughness.

Furthermore, alloying strategies involving Cr-N and C-N have been observed to increase the number of ultrafine twins, bolstering the material’s strength while maintaining toughness. In the same vein, high-manganese steels subject to V-microalloying and a subsequent suitable heat treatment can also result in the formation of nanoscale carbides, achieving this superior outcome.