Recent research efforts to develop advanced–/ultrahigh–strength medium-Mn steels have led to the development of a variety of alloying concepts, thermo-mechanical processing routes, and microstructural variants for these steel grades. However, certain grades of advanced–/ultrahigh–strength steels (A/UHSS) are known to be highly susceptible to hydrogen embrittlement, due to their high strength levels. Hydrogen embrittlement characteristics of medium–Mn steels are less understood compared to other classes of A/UHSS, such as high Mn twinning–induced plasticity steel, because of the relatively short history of the development of this steel class and the complex nature of multiphase, fine-grained microstructures that are present in medium–Mn steels. The motivation of this paper is to review the current understanding of the hydrogen embrittlement characteristics of medium or intermediate Mn (4 to 15 wt pct) multiphase steels and to address various alloying and processing strategies that are available to enhance the hydrogen-resistance of these steel grades.
The production routes for medium-Mn steels and resulting microstructures are shown in the schematic of Figure 1. Note that many of the microstructural variants presented in Figure 1 are developed through processing of a “cold-rolled” product. It is anticipated, however, that some of the processing and microstructural engineering concepts developed for the cold-rolled product can be extended to hot-rolled or plate-rolled products.
A key attribute of both lamellarized and equiaxed microstructures is the presence of either UFG or nanolaminate
[ 13 ] retained austenite, respectively, which contributes to enhanced work hardening and high ductility associated with the TRIP and/or TWIP effect. For medium-Mn steels, submicron-sized retained austenite grains, typically smaller than 500 nm, are obtained by a relatively simple IA step or austenite reversion annealing. It is also important to note that IA is essential to obtain a substantial amount of retained austenite with adequate stability as carbon (C) and Mn, both strong austenite stabilizers, partition from ferrite or martensite to austenite during IA. The origin of the development of a high fraction of UFG austenite in the intercritically annealed medium-Mn steel is largely associated with the Mn austenite-stabilizing effect [ 2 , 14 , 15 ]. When the bulk Mn content is increased, the reverse transformation to austenite starts at a lower temperature and, thus, a relatively high fraction of austenite may be produced before recrystallization takes place. Suh and Kim [ 2 ] pointed out that grain growth in the ultra-fine regions is restricted even at a relatively high IA temperature (~800 °C), presumably due to the nature of the two-phase microstructure.
Some alloy and microstructure variants of typical UFG duplex medium-Mn steel are also shown in Figure 1. For example, many investigators considered aluminum (Al) and silicon (Si) as essential alloying elements for medium-Mn steels. These elements are useful in raising the intercritical temperature and, thus, increasing the solute partitioning kinetics. In addition, it has been recognized that a bimodal grain structure can be achieved in medium-Mn steels by the addition of Al and Si, when the amount of both alloying elements exceeds 3 wt pct
[ 16 , 17 ]. Fe–xC-6 wt pct Mn steels were employed in the initial studies of Al-added medium-Mn TRIP steels [ 16 , 18 ]. Figure 2 shows the influence of Al and Si additions on the pseudo-binary phase diagram for an Fe–6.0Mn–xC alloy system (all in wt pct), calculated by using Thermo-Calc® software. The Al and Si additions expand both the α-ferrite and δ–ferrite stability regions, resulting in the formation of layers of coarse δ–ferrite grains in the hot–/cold-rolled microstructures (third example in Figure 1). For specific alloy and heat-treatment conditions, a steel specimen with a UFG microstructure often experiences pronounced localized deformation and plastic instability, associated with dynamic strain aging [ 19 ]. The introduction of coarse δ–ferrite grains in the UFG structure, resulting in a bimodal grain size distribution, may eliminate the localized deformation and plastic instability and enhance ductility [ 20 ]. In another alloying and processing variation, “warm” rolling at an IA temperature has been applied to Al-free and Al-containing medium-Mn steels, leading to an even finer-grained, lamellarized microstructure, elongated along the rolling direction [ 21 , 22 , 23 ] (fourth example in Figure 1).
Similar to the Q&T/Q&P applications, the application of hot or “warm” stamping to medium-Mn steels produces microstructures with retained austenite in addition to martensite, which provides an ultrahigh strength (tensile strength up to 1.6 GPa) and a ductility (total elongation up to 44 pct) much higher than that of conventional 22MnB5 press-hardened steel (7–8 pct)
[ 32 , 33 , 34 ]. There are additional advantages of the application of medium-Mn steel to hot stamping over the conventional low-alloy boron-added steels; for example, Mn-associated enhanced hardenability reduces the cooling rate requirement and enables alternative cooling approaches [ 32 , 34 ]. In addition, compared to low-alloy steels, lower reheating (austenitizing or IA) temperatures can be employed for hot stamping of medium-Mn steel, which reduces both the heating time and potentially production costs, and also influences the behavior of metallic coatings that may be employed.
More specifically, the effects of constituent morphologies, austenite fraction and stability, alloying elements, e.g., Al, Si, and copper (Cu), and processing routes on the HE sensitivity are presented here.
A number of investigations employed SSRT to compare the HE characteristics of two different (lamellarized and equiaxed) morphologies of a medium-Mn steel microstructure, as summarized in Table 1. In SSRT, relative elongation loss due to H (HE index) is commonly used as a metric to measure a HE sensitivity of a material, and the HE index is given by the following equation:
On the other hand, Han et al.
[ 36 ] focused on the fact that the fracture and cracking mechanisms are different in the two conditions. For the lamellarized microstructure, H-induced fracture occurred preferentially along “prior” austenite grain boundaries, whereas fracture in the equiaxed microstructure occurred either across the ferrite grains or along the granular retained austenite grain boundaries. Greater degrees of crack deflection were present in the equiaxed condition, which was interpreted to improve H-resistance, as compared to the lamellarized condition. A more recent study [ 39 ] performed three-dimensional (3D) atom probe tomography (APT) of the lamellarized condition, used by Han et al. [ 36 ]. The 3D-APT analyses revealed noticeable segregation of C, Mn, and P at prior austenite grain boundaries, which may explain the lamellarized microstructure being more prone to prior austenite grain boundary cracking. The unfavorable influence of the combination of H and segregated elements (e.g., Mn, Si, P, and S) on the severity of HE has been demonstrated in other publications [ 40 , 41 ]. Therefore, control of impurity levels and segregation can also be a factor determining which microstructures better resist HE.
According to the results summarized in Table 1, the effect of microstructural morphology on H absorption is usually not significant for microstructures with a similar amount of austenite (e.g., the studies by Han et al.
[ 36 ] and Jeong et al. [ 10 ]). In contrast, only the study by Cameron et al. [ 35 ] was unique, given that the equiaxed morphology absorbed a higher H content, nearly an order of magnitude than that of the lamellarized microstructure, for the same charging condition. The larger H absorption may partly be explained by the higher fraction of austenite, where H is more soluble than in martensite (and ferrite), in the equiaxed microstructure.
Retained austenite is the key constituent controlling activation of the TRIP and/or TWIP effect, which have a pronounced influence on the strain hardening behavior, in medium-Mn steels. Some studies suggest that retained austenite is also beneficial for the resistance of medium-Mn steel to HE because of its high H-trapping capability or its positive contribution to overall ductility (toughness). As an example, Wang et al.
[ 42 ] investigated the influence of austenite fraction and film thickness, associated with austenite grain size, on H absorption and HE characteristics of a medium-Mn steel (Fe–0.01C–9Mn–3Ni–1.4Al, in wt pct), as shown in Figure 5. Figure 5a shows an electron backscatter diffraction (EBSD) phase map of a fully lath martensitic microstructure in the as-quenched condition. Wang et al. [ 42 ] produced two additional microstructures by applying an austenite reversion treatment to the as-quenched martensite. When the steel specimen was austenite reversion-treated at 600 °C with a hold time of 1 h, the microstructure contained austenite with a volume fraction of 10 pct and an average thickness of 200 nm, as shown in the EBSD phase map in Figure 5b. As the hold time at 600 °C increased to 8 h, the austenite fraction increased to 35 vol pct and the average thickness of the austenite films increased to 500 nm (Figure 5c). The H-TDA results by Wang et al. [ 42 ] (Figure 5d) indicate that the specimen containing a greater amount of retained austenite absorbs more H for a given H-charging condition because H solubility in austenite is much higher than in ferrite or martensite. The engineering stress–strain curves obtained by SSRT (Figure 5e) show that the H-induced ductility loss became less significant with increasing austenite fraction. Similar to the study by Wang et al. [ 42 ], Du et al. [ 43 ] produced three different microstructures from a medium-Mn steel (Fe–0.065C–5.45Mn–0.2Si, in wt pct): an as-quenched martensitic microstructure and two intercritically annealed microstructures containing different amounts of retained austenite within the martensitic matrix. Du et al. [ 43 ] also presented a beneficial effect of retained austenite with respect to the H-resistance in the medium-Mn steel. Figure 6 shows the SSRT results obtained by Du et al. [ 43 ], indicating that with increasing fraction of retained austenite, the H content of the specimen significantly increased, and the H-induced elongation loss became less significant.
A detrimental effect of mechanically-induced martensite is also highlighted in a study by Wang et al.
[ 46 ], who investigated the influence of pre-strain on HE of a medium-Mn steel (Fe–0.25C–8.7Mn–2.69Al–0.5Si, in wt pct). Wang et al. [ 46 ] reported that pre-straining resulted in an increase in dislocation density and partial transformation of austenite to martensite, leading to a substantial increase in the HE susceptibility. Liu et al. [ 47 ] also explored the pre-strain effect on H-induced cracking of medium-Mn steel, emphasizing the importance of austenite mechanical stability to suppress H-induced damage, similar to the study by Wang et al. [ 46 ].
It should be noted that, when the annealing conditions are varied, the steel strength, which is also a factor influencing the H-sensitivity, changes, in addition to the microstructure. For example, the studies by Wang et al.
[ 43 ] and Du et al. [ 43 ] clearly showed that the HE susceptibility decreased with increasing retained austenite fraction (Figure 5 and Figure 6). However, in these studies, the specimen with a higher fraction of retained austenite exhibited a reduced yield strength, which may have contributed to the increased H-resistance. In contrast, the study by Shao et al. [ 44 ] clearly showed that, as IA time increased, HE became increasingly severe (Figure 7), despite the decreases in yield and tensile strength. Therefore, it was clear, in this study, that the H-resistance of medium-Mn steels was influenced by the mechanical stability of retained austenite.
A few investigations evaluated the HE characteristics of substantially Al- and Si-alloyed medium-Mn steels containing coarse δ-ferrite grains (Table 2). Figure 8a,b shows the microstructures of medium–Mn steels containing lower (1.1 wt pct) and higher Al contents (3.1 wt pct), respectively
[ 48 ]. The microstructure of the low Al medium-Mn steel exhibited a relatively uniform distribution of UFG austenite and ferrite (Figure 8a), while the high-Al/medium-Mn steel contained coarse δ-ferrite layers in addition to the UFG austenite–ferrite regions (Figure 8b). Figure 8c,d shows the tensile properties of the low-Al and high-Al steels in H-free (air) and H-charged conditions, evaluated by means of SSRT. In the H-free condition, the low Al steel exhibited a high ultimate tensile strength and high work-hardening rate, associated with a pronounced TRIP effect (Figure 8c). Comparatively, the high-Al steel had an increased ductility and a low work-hardening capacity (Figure 8d). Ryu et al. [ 48 ] measured austenite fraction of each specimen as a function of the tensile strain and confirmed that the mechanical stability of austenite in the high-Al steel was higher compared to the low-Al steel. The SSRT results of H-charged specimens (Figure 8c, d) indicate that the H-induced loss in ductility was more pronounced in the low–Al/medium-Mn steel containing less stable austenite.
The beneficial effect of coarse δ–ferrite grains or a bimodal grain size distribution on the mechanical behavior in the H–free condition is well known
[ 17 , 51 ]. That is, it has been suggested that there is enhanced strain partitioning in coarse, soft δ-ferrite grains, and thus, γ in the UFG region is stabilized, until larger strains are reached during deformation. This strain-partitioning behavior is believed to positively affect overall ductility and toughness. However, the direct role of δ–ferrite with respect to H-resistance is not clear in the studies summarized in Table 2. More often than not, the overall HE characteristics of a medium-Mn steel appear to be affected more significantly by the austenite fraction and mechanical stability rather than by the presence of δ–ferrite. However, there is some evidence that δ–ferrite layers might play a role in cracking or influence the fracture path in H testing. Wang et al. [ 49 ] reported that H-induced cracks are often arrested by the δ–ferrite layer. Warm rolling at an IA temperature is reported to provide a refined microstructure and significantly improved the H–resistance of a medium–Mn steel containing δ–ferrite. Zhang et al. [ 50 ] considered the refined δ–ferrite layers in a warm-rolled specimen to be beneficial, hypothesizing that the soft δ–ferrite exhibits low HE susceptibility, resulting in an effective crack termination mechanism. The contribution of the refined δ–ferrite layers in a warm-rolled specimen to deflection of the H–induced cracks is highlighted in Figure 9 [ 50 ]. Scanning electron microscopy (SEM) fractographs compare the path and morphology of the H-induced cracks in medium-Mn steels after conventional IA and after intercritical warm rolling. The crack in the intercritically annealed steel was slightly deflected at the δ-ferrite regions (Figure 9a). In the warm-rolled microstructure, the fine lamellar structure with δ–ferrite layers resulted in delamination-type fracture; the crack that encountered the δ–ferrite layer was significantly deflected and propagated in the rolling direction, rather than in the transverse direction (Figure 9b).
In addition to its role in δ–ferrite stabilization, Al additions significantly increase SFE, which is associated with mechanical stability of austenite. This dual role of Al may be important in the future design of a H-resistant ferrite–austenite microstructure with increased mechanical stability of austenite. There has been considerable discussion of the interaction between Al and H in “high” Mn TWIP steel
[ 52 , 53 , 54 ]. In summary, Al has been found to slow the diffusion of H in a high-Mn TWIP steel, which leads to the suppression of H uptake and embrittlement [ 53 ]. It was also suggested that an Al2O3 oxide layer that forms at the surface of TWIP steel interferes with H absorption and, thus, prevents the H uptake [ 54 ].
Carbide and nitride precipitates containing molybdenum (Mo), vanadium (V), niobium (Nb), and titanium (Ti) have long been known to be effective in mitigating H-induced fracture of steels. However, the influence of micro alloy precipitates has not been broadly explored in medium-Mn steels. Park et al.
[ 55 ] evaluated the HE resistance of medium-Mn steels micro alloyed with Nb, Ti, and V compared to a reference steel (Fe–6Mn–0.08C, in wt pct). They reported that the effects of the microalloying additions on the H-resistance were insignificant for the investigated test conditions. They explained that the insignificant effects of the precipitates were due to the two-phase nature of the microstructure, i.e., micro alloy carbides formed primarily in the ferritic regions, but the H atoms are expected to be trapped preferentially in the austenite, where H is more soluble than in ferrite. It is an important observation that, in a duplex microstructure, carbides and nitrides are mostly present in the ferrite, rather than the austenite, and may not actively interact with H atoms in the microstructure.
Cu is not a carbide/nitride former but can also precipitate in steel, contributing to precipitation strengthening. Li et al.
[ 56 ] investigated precipitation behavior and HE in Cu–free (Fe–7Mn–2.5Ni–1.5Al–0.01C, in wt pct) and Cu-added medium-Mn steels (Fe–0.01C–7Mn-–2.5Ni–1.5Al–1.5Cu, in wt pct). In processing these two steels, Li et al. [ 56 ] used two-stage annealing; the first IA was performed at 630 °C for 1 h and the second step involved “tempering” at 500 °C. The hold times in the “tempering” step used for the Cu–free and Cu-added steels were 5 and 2 h, respectively. Figure 10a,b shows the 3D–APT results for the ferritic and austenitic regions in the Cu-added medium-Mn steel, indicating that Cu, as opposed to the microalloying example above [ 55 ], precipitated primarily in the austenite, rather than in the ferrite, despite the increased solubility of Cu in austenite compared to ferrite. In contrast, the APT map of the ferritic region (Figure 10a) shows non-uniformity in the Ni signal; TEM selected area diffraction analysis by Li et al. [ 56 ] revealed that the Ni-enriched regions in the ferrite are associated with B2 NiAl precipitates. Figure 10c shows the influence of Cu on the HE characteristics, evaluated by SSRT. The Cu-added steel exhibited a smaller elongation loss due to H, indicating a higher resistance to HE compared to the Cu-free steel. It should be noted that the interpretation of the primary role of the Cu precipitates with respect to the material’s H-resistance is different from the “conventional” mechanism from a viewpoint of precipitates being strong H-traps. The authors explained that the improved H–resistance in the Cu-added steel was due to increased strain compatibility between austenite and ferrite, resulting from the Cu precipitation. Li et al. [ 56 ] employed nano-indentation to support this view. Figure 10d shows the average indentation hardness of the ferrite and austenite phases in the Cu–free and Cu–added steels. In both the Cu-free and Cu-added steels, the ferrite was relatively hard (average nano-hardness > 8 GPa), possibly due to the NiAl associated precipitation strengthening of ferrite. In the Cu-free steel, the average hardness of ferrite (~10 GPa) is much harder than that of austenite (~5 GPa). In the Cu–added steel, however, Cu precipitation strengthening led to a significant increase in the average hardness of the austenite, which reduces the difference in hardness between the two constituent phases and, thus, improves the strain compatibility. The possible detrimental effect of strain incompatibility between neighboring constituent phases on the H–induced cracking of a medium–Mn steel was also discussed in a recent publication by Sun et al. [ 57 ]. It should be noted that, in the study by Li et al. [ 56 ], the improved H–resistance was noted despite the fact that the mechanical stability of the austenite was lower in the Cu-added steel. This study is one of only a few examples indicating that HE was less significant in the presence of a more pronounced TRIP effect.
Table 3 summarizes several studies related to processing- and microstructure-dependent fracture surface appearance and crack initiation sites of H–embrittled, medium–Mn steels. In general, different phase boundaries and grain boundaries in medium-Mn steel microstructures may serve as crack initiation sites; microstructural morphologies and constituents strongly influenced the H–induced cracking and fracture of medium–Mn steels. A few studies
[ 36 , 44 ] observed H–induced cracking at/near mechanically–induced martensite. The studies by Han et al. [ 36 ], Wang et al. [ 49 ], and Zhang et al. [ 50 ] indicated that the lamellarized austenite–martensite microstructure is often subject to H–induced cracking along the prior austenite grain boundaries. However, this type of fracture is not observed in an equiaxed ferrite–austenite microstructure which does not preserve its prior austenite grain structure. Instead, according to studies by Han et al. [ 36 ] and Shao et al. [ 44 ], the fracture surfaces of the equiaxed microstructure often display dimples filled with fine granular features. Examples of these features are indicated by yellow arrows in Figure 11a, the appearance of which is different from “typical” ductile micro voids indicated by white arrows in Figure 11a. A detailed image of a fine granular feature is shown in Figure 11b. Han et al. [ 36 ] and Shao et al. [ 44 ] interpreted the fine granular features to be associated with a region of either equiaxed austenite or mechanically–induced martensite. In particular, Han et al. [ 36 ] attributed the granular features to intergranular cracking along equiaxed austenite grain boundaries. Shao et al. [ 44 ] showed that the Mn content of the granular features was higher than that of surrounding microstructure implying they are associated with austenite (or mechanically–induced martensite) since austenite tends to be more enriched in Mn than ferrite or martensite in these steels.