Role of Mo/V carbides in hydrogen embrittlement of tempered martensitic steel

1 Introduction

High fuel efficiency is receiving massive attention in the automotive industry due to the issues of energy conservation and environmental protection. Engineers have endeavored to reduce vehicle weight using high-strength steels to improve fuel efficiency. Tempered martensitic steel is one of various types of high-strength steels. This steel has been favored in the industry because it is possible to easily obtain high strength by a simple heat treatment, whereas other steels (e.g. pearlitic steel) require a significant amount of plastic deformation to attain such a high strength (Chun, Park, & Lee, 2012c; Kim, Lee, Lee, Park, & Lee, 2009; Krauss, 1999).

Hydrogen embrittlement (HE) refers to a phenomenon where a material experiences brittle fracture by internal hydrogen elements. High-strength steels exhibit poor HE resistance because their high defect density promotes hydrogen-assisted crack growth (Hirth, 1980). Namimura (2002) reported that HE susceptibility dramatically increased when the ultimate tensile stress (UTS) of a steel exceeded 1.2 GPa. In addition, Chun, Lee, Bae, Park, and Lee (2012b) suggested that tempered martensitic steel is particularly vulnerable to the HE phenomenon due to its bcc nature.

A number of studies have been carried out to increase the HE resistance in tempered rmartensitic steels. Nie et al. (2012) suggested three approaches to attain this objective: improvement of intrinsic fracture resistance, reduction of hydrogen elements at crack tips, and suppression of hydrogen intrusion. Some researchers have modified the grain structure of materials, including grain size and shape, to decrease the local hydrogen concentration (Banerji, McMahon, & Feng, 1978; Padmanabhan & Wood, 1983). However, with these methods, it is difficult to increase the HE resistance of the bulk plate of tempered martensitic steel. Other researchers have suggested that HE is suppressed by the precipitation of various carbides in tempered martensitic steels (Akiyama, 2012; Kang et al., 2012; Li et al., 2010; Liu, Wang, & Liu, 2014; Nagao, Martin, Dadfarnia, Sofronis, & Robertson, 2014). In particular, Li et al. (2010) studied NIMS17 steel and reported that it offered high strength and strong HE resistance in comparison to commercial steels by utilizing carbide precipitation.

Although the precipitation of carbide is known as a useful method to enhance HE resistance, few studies have focused on the effects of carbide characteristics. The authors focused on Mo carbide and V carbide in the present study because they can be easily controlled to improve HE resistance, in contrast to other types of carbide, such as Ti carbide and Nb carbide. Thus, the present study clarified the role of these carbides in the HE behavior of tempered martensitic steels.

2 Materials and methods

Four ingots melted in a vacuum arc furnace were hot-rolled to a thickness of 13 mm, during which the temperature was decreased from 1200°C to 1000°C. Table 1 shows the chemical composition of each plate determined after the hot-rolling process at POSCO. The plates possessed different compositions of Cr, Mo, and V to change the carbide characteristics, whereas the other elements were controlled at the same amounts. Cr-Mo steel followed the chemical composition of NIMS17 steel (Li et al., 2010). Cr steel was designed to remove carbide from Cr-Mo steel. Mo steel was prepared to clarify the role of Cr elements in the HE phenomenon of tempered martensitic steel. Finally, Cr-V steel exchanged Mo for V to compare the effect of Mo carbide and V carbide on HE behavior. It is noted that the amount of V elements was reduced by half to maintain the standard atomic weight.

The plates were austenitized for 30 min and quenched to 60°C in an oil solution to eliminate residual stress, which may generate microcracks. The materials were then tempered for 60 min, followed by water quenching. The temperatures for the austenitizing and tempering processes were determined on the basis of the conditions of NIMS17 steel (i.e. Cr-Mo steel in this work) (Li et al., 2010), as summarized in Table 2. Cr steel was tempered at lower temperature to obtain a similar UTS compared to Cr-Mo steel. Meanwhile, Cr-V steel was austenitized at higher temperature, as V carbides dissolved less readily than Mo carbides.

Tensile properties were determined at a strain rate of 5×10^-3 s^-1 at room temperature. Gauge length and diameter of the tensile specimen were 25 and 6 mm, respectively. The tests were repeated thrice for each material using a 25-mm extensometer for reproducibility of the data. Microstructural characterization was carried out by means of a scanning electron microscope (SEM) and transmission electron microscope (TEM). SEM samples, excluding those for fracture surfaces, were prepared by mechanical polishing and subsequent etching in a 2% nital solution. Fracture surfaces were directly observed immediately after the tensile test without any additional treatment but cleaning in acetone. TEM samples were prepared by two different methods. First, carbon extraction replicas were made using replicating tape and an evaporation carbon coater. Second, samples were prepared by focused ion beam method at the National Institute for Nanomaterials Technology.

The HE phenomenon can be clearly observed at the notch of a material as hydrogen atoms move to this location under the hydrostatic field created by external stress (Gerberich & Chen, 1975; Johnson & Troiano, 1957; Keijiro, Hideo, & Xiaolie, 1986; Krom, Koers, & Bakker, 1999; Lufrano & Sofronis, 1998; Oriani & Josephic, 1974; Sofronis & McMeeking, 1989; Toribio, 1993, 1997; Toribio & Elices, 1991; Yokobori, Chinda, Nemoto, Satoh, & Yamada, 2002). Accordingly, notch specimens with an intensity factor of 4.37 and a diameter of 6 mm were prepared to evaluate the HE resistance of the investigated steels. These samples were mechanically polished with 1200-grit SiC paper to eliminate surface flaws prior to hydrogen charging. Hydrogen was charged to a sample in 0.1-n NaOH solution at a charging current of 0.1–30 A/m² for 48–72 h. Thermal desorption spectroscopy (TDS) was carried out to measure the amount of hydrogen atoms in the materials by means of Q-mass and gas chromatography. During the analysis, samples were heated to 800°C with a heating rate of 100°C/h. Hydrogen-charged samples were coated with Cd at a current of 50 A/m² for 5 min to minimize hydrogen escape during a slow strain rate test (SSRT) (Akiyama, 2012; Chun, Kim, Park, Lee, & Lee, 2012a; Kang et al., 2012; Kim et al., 2009; Kim, Chun, Won, Kim, & Lee, 2013; Wang, Akiyama, & Tsuzaki, 2007). The SSRT was performed at a loading rate of 5×10^-3 mm/min at room temperature.

3 Results

Figure 1 shows an SEM micrograph of the investigated steels presenting the typical microstructure of tempered martensite. The materials were composed of lath martensites, of which the average thickness was similar: 0.83 μm for Cr steel, 0.70 μm for Cr-Mo steel, 0.90 μm for Mo steel, and 0.75 μm for Cr-V steel. Figure 2 shows a TEM image and EDS pattern of the replica samples. Cr carbides, such as Cr₃C₂ and Cr₇C₃, were rarely observed in the investigated tempered martensitic steels; the intensity of Cr peaks in the EDS pattern was small enough to be neglected. Hence, the effect of Cr carbide was excluded to evaluate the HE behavior in the present work.

Figure 1:

SEM micrograph of (A) Cr steel, (B) Cr-Mo steel, (C) Mo steel, and (D) Cr-V steel.

Figure 2:

TEM bright-field image and EDS results of the replica specimens: (A) Cr steel, (B) Cr-Mo steel, (C) Mo steel, and (D) Cr-V steel.

The red arrows indicate the points for EDS analysis.

The EDS results show the absence of carbides in Cr steel except for cementite; Cu peaks are ascribed to the Cu grid used for the analysis. This is attributed to the tempering temperature of Cr steel, which was too low to induce carbide precipitation (Bhadeshia & Honeycombe, 2011). Gojic, Kosec, and Matkovic (1998) reported the similar results that Cr-added steel contained only cementites after a tempering process. It is also noted that cementite particles precipitated along the lath-grain boundaries, as confirmed in Figure 3 where the carbon intensity increased in the vicinity of these boundaries. Meanwhile, Cr-Mo steel and Mo steel exhibited Mo carbides, whereas Cr-V steel possessed V carbides. These carbides were not limited within lath boundaries but were distributed more uniformly than cementites in Cr steel. The tempering at the higher temperature also gave rise to the globularization of carbides. The TEM micrograph demonstrates undissolved carbides with a diameter of ~100 nm. In addition, carbides with a finer particle size (10–50 nm) were also confirmed, which precipitated during the tempering process as reported in the literature (Nagao et al., 2014; Porter, Easterling, & Sherif, 2009).

Figure 3:

TEM bright-field image and line-mapping results of Cr steel investigating (A) prior austenite boundaries and (B) lath boundaries.

Figure 4 shows the SSRT net fracture stress of the investigated steels depending on the amount of diffusible hydrogen. All samples exhibited a decrease in SSRT net fracture stress with increasing hydrogen content due to the HE phenomenon. Nevertheless, it is noted that Cr-Mo steel and Mo steel demonstrated the high SSRT net fracture stress at the same amount of diffusible hydrogen, suggestive of their strong HE resistance. In contrast, Cr steel recorded the weakest HE resistance, with the SSRT net fracture stress decreasing from 2000 to 500 MPa after the hydrogen charging. Cr-V steel exhibited significantly scattered data. The SSRT results also include the data after 24 h from the hydrogen charging. These results are discussed further in Section 4.2.

Figure 4:

SSRT results of the investigated steels with various amounts of diffusible hydrogen.

The results also include the data of Cr-Mo steel and Cr-V steel obtained after 24 h from the hydrogen charging.

The tendency of the HE resistance depending on the material was also confirmed by the fracture surface of the hydrogen-charged samples, as shown in Figure 5. It has been widely reported that intergranular fracture occurs in a hydrogen-charged tempered martensitic steel (Kim et al., 2009). The high yield strength of the lath structure with a high dislocation density and precipitate structure can promote hydrogen concentration at the notch tip due to the enhanced triaxiality. Cr steel exhibited the typical fractograph of an intergranular-fractured surface, implying its vulnerability to the HE phenomenon. Meanwhile, the other samples demonstrated surfaces with quasi-cleavage fracture and a partial area of ductile fracture composed of dimples.

Figure 5:

SEM image of fracture surface of hydrogen-charged tempered martensitic steels: (A) Cr steel, (B) Cr-Mo steel, (C) Mo steel, and (D) Cr-V steel.

Figure 6 shows the TDS results obtained by Q-mass from hydrogen-charged samples. A single peak was confirmed for all investigated steels regardless of the chemical composition. The peak of Cr-V steel was shifted to higher temperature than the other materials. Yokota and Shiraga (2003) ascribed this phenomenon to the precipitation of V carbides by comparing TDS peaks of V-free and V-bearing steel. Hydrogen-trapping sites are classified into two categories in general: diffusible and non-diffusible trapping sites. The criterion is based on the activation energy required for hydrogen to escape from the trapping sites (Chun et al., 2012a). Diffusible hydrogen-trapping sites are characterized by lower activation energy than non-diffusible trapping sites, and thus, hydrogen at diffusible trapping sites is released at lower temperature (<300°C) (Takai & Watanuki, 2003). Therefore, the TDS peaks at 100°C–270°C suggest that hydrogen elements were mainly trapped at diffusible trapping sites in the present tempered martensitic steels.

Figure 6:

TDS results of the investigated steels after hydrogen charging at a current density of 10 A/m² for 48 h.

The data were obtained by the Q-mass method.

Dislocations and lath-grain boundaries have been revealed as diffusible hydrogen-trapping sites by a number of researchers (Brown, Hochmann, Slater, McCright, & Staehle, 1977; Choo & Lee, 1982; Hirth, 1980; Kim, Lee, Lee, Park, & Lee, 2009). In addition, Wei, Hara, and Tsuzaki (2011) recently reported that the type of hydrogen-trapping sites (i.e. diffusible or non-diffusible) depended on the type of carbides. It is of particular note that they proved that Mo carbide and V carbide were diffusible hydrogen-trapping sites. Consequently, the single TDS peak in this work resulted from diffusible hydrogen trapped at dislocations, lath-grain boundaries, cementites, and carbides.

4 Discussion

4.1 Enhanced HE resistance by carbide formation

The SSRT results and fractographs verified the significant vulnerability to HE phenomenon in the carbide-free steel (i.e. Cr steel). The poor performance of Cr steel originates from the characteristics of its hydrogen-trapping sites including lath-grain boundaries, dislocations, and cementite particles. Generally, dislocations accumulate at lath-grain boundaries in a tempered martensitic steel (Krauss, 1999). TEM observation also confirmed that cementite particles mainly precipitated at lath-grain boundaries in Cr steel, as shown in Figures 2A and 3. These results indicate the presence of a cluster of hydrogen-trapping sites along the lath-grain boundaries. The local hydrogen concentration on these sites decreased the interfacial strength between lath-grain boundaries, resulting in the intergranular fracture (Kameda & McMahon, 1983; Yamaguchi et al., 2011).

Meanwhile, the present carbide steels (i.e. Cr-Mo steel, Mo steel, and Cr-V steel) exhibited a clear improvement in HE resistance. The lath thickness was comparable among the investigated steels, as mentioned in Section 3, indicating that the extent of hydrogen trapping at lath-grain boundaries was similar regardless of the materials. It is also noted that the carbide steels were tempered at a higher temperature of 570°C than Cr steel, which gave rise to a reduction of dislocation density. Despite these factors, the carbide steels exhibited larger TDS peaks compared to Cr steel, suggesting that carbides trapped a considerable amount of diffusible hydrogen. This deduction is also supported by Figure 7, where Cr steel was hydrogen-charged at a current density three times higher than that of Cr-Mo steel to trap a similar amount of hydrogen for comparison. Most of the hydrogen elements were detrapped from Cr steel after heating the material at 150°C for 1 h. In contrast, half of the hydrogen remained in Cr-Mo steel after the same heating process. Therefore, it is reasonable to consider that the hydrogen elements were mainly trapped in Mo carbides (for Cr-Mo steel and Mo steel) or V carbides (for Cr-V steel). These carbides were widely distributed in the matrix, as shown in Figure 2B–D, which alleviated the local hydrogen concentration on lath-grain boundaries and thus inhibited the intergranular fracture in the carbide steels.

Figure 7:

TDS results of Cr steel and Cr-Mo steel obtained immediately after hydrogen charging and after heating at 150°C for 1 h.

Cr steel was hydrogen-charged at a current density of 30 A/m² for 72 h, whereas Cr-Mo steel was charged at 10 A/m² for 72 h.

4.2 Difference between Mo carbide and V carbide in terms of HE resistance

It is confirmed in Table 1 that Cr-Mo steel and Cr-V steel were composed of similar chemical compositions except for Mo and V. Accordingly, the difference between Mo carbide and V carbide can be revealed by comparing these two steels. Cr-V steel recorded weaker HE resistance in comparison to Cr-Mo steel, and this is characterized by the lower SSRT net fracture stress in Figure 4. The material also showed significantly scattered data, indicating the unstable HE resistance of Cr-V steel. These differences are rationalized by two factors. The first factor is the capability of hydrogen trapping of each steel. Cr-V steel demonstrated the largest peak of diffusible hydrogen among the investigated tempered martensitic steels in Figure 6, suggesting that V carbides trapped a higher amount of hydrogen elements than Mo carbides. This is ascribed to the high chemical affinity of V element with hydrogen. When vanadium and carbon make a covalent bond, an extra vanadium electron is left, which may attract hydrogen atom (Pressouyre, 1979). Quantitative TEM analysis provided the total surface of 0.61 μm² for Mo carbides and that of 0.50 μm² for V carbides in the area of 18.89 μm², assuming the spherical morphology of carbides. In other words, a higher amount of hydrogen was trapped at a smaller area of trapping sites in Cr-V steel, leading to a stronger hydrogen concentration compared to that in Cr-Mo steel. This reduced the HE resistance of the material by the mechanism explained in Section 4.1.

The second factor is the fraction of undissolved and precipitated carbides. To continue the discussion, it is required to determine which type of carbides (i.e. undissolved or precipitated) has higher activation energy for the diffusion of hydrogen. As aforementioned with regard to the work of Yokota and Shiraga (2003), the precipitated carbides move the TDS curve toward the region of higher temperatures. It is thus concluded that these carbides have higher activation energy for the hydrogen diffusion in comparison to undissolved carbides. Solubility of Mo in austenite is always higher than that of V at 800°C–1000°C on the basis of thermodynamic calculation using Thermo-Calc Software (Stockholm, Sweden) and TCS Steel and Fe-alloys Database Version 7 (Andersson, Helander, Höglund, Shi, & Sundman, 2002). In particular, the solubility of Mo (2.65×10^-3 in mole fraction) at 920°C is still higher than that of V (1.19×10^-3) at 980°C (note that these are the austenitizing temperatures for each material). This indicates the higher amount of undissolved V carbides in Cr-V steel compared to the amount of undissolved Mo carbides in Cr-Mo steel, despite the higher austenitizing temperature for Cr-V steel.

Hydrogen elements readily accumulated in the vicinity of undissolved carbides due to the low activation energy for diffusion, which contributed to the reduced HE resistance of Cr-V steel as the material contained a considerable amount of undissolved V carbides. This is also supported by the SSRT results obtained after 24 h from the hydrogen charging (Figure 4) and corresponding TDS results (Figure 8). In Figure 4, Cr-Mo steel and Cr-V steel exhibited comparable SSRT net fracture stress at this point, in contrast to the large difference observed immediately after the charging process. In Figure 8, the TDS peaks were decreased and shifted to higher temperature for both steels after 24 h because hydrogen was detrapped from trapping sites with low activation energy including dislocations, lath-grain boundaries, and undissolved carbides (Yokota & Shiraga, 2003). This is consistent with the report of Asahi, Hirakami, and Yamasaki (2003), which noted that hydrogen still remained at precipitated carbides in V-added steel 28 h after the charging process. Considering the similar distribution of dislocations and lath-grain boundaries in the investigated carbide steels, the higher recovery of SSRT net fracture stress after 24 h in Cr-V steel is attributed to hydrogen detrapping from the higher amount of undissolved carbides.

Figure 8:

TDS results of Cr-Mo steel and Cr-V steel obtained immediately after hydrogen charging and 24 h after charging.

All steels were hydrogen-charged at a current density of 10 A/m² for 72 h.

5 Conclusions

The present study investigated four different steels whose chemical composition had been carefully controlled. All materials demonstrated the typical microstructure of tempered martensite with similar lath thickness. Cr steel was a carbide-free alloy. Cr-Mo steel and Mo steel exhibited Mo carbides, whereas Cr-V steel possessed V carbides. Undissolved carbides with a diameter of ~100 nm and precipitated carbides of 10–50 nm were confirmed in these steels. For the investigated alloys, hydrogen was trapped at diffusible hydrogen-trapping sites, such as dislocations, lath-grain boundaries, cementites, and carbides. Cr steel recorded the weakest HE resistance, which is attributed to the local hydrogen concentration on the clusters of hydrogen-trapping sites along the lath-grain boundaries. This reduction of interfacial strength gave rise to intergranular fracture in the alloy. The unstable and low HE resistance of Cr-V steel is due to the following two reasons. First, V carbides trapped a higher amount of hydrogen elements than Mo carbides due to their high chemical affinity, resulting in stronger local hydrogen concentration. Second, Cr-V steel included a higher fraction of undissolved carbides due to the high melting point of V elements. This work showed that undissolved carbides had lower activation energy for hydrogen diffusion in comparison to precipitated carbides. Meanwhile, Cr-Mo steel exhibited the highest UTS as well as strong HE resistance in this study. This originated from the Cr element, taking into consideration that Cr-Mo steel and Mo steel possessed similar characteristics with regard to Mo carbides. Cr elements increased the UTS of Cr-Mo steel by retarding the softening of martensite.