Improving hydrogen embrittlement resistance of
martensitic steel via interface B segregation

Hao Shi, Guillaume Hachet, Huijie Cheng, Prithiv Thoudden Sukumar, Dirk Ponge, Binhan Sun

► To cite this version:

Hao Shi, Guillaume Hachet, Huijie Cheng, Prithiv Thoudden Sukumar, Dirk Ponge, et al.. Improving hydrogen embrittlement resistance of martensitic steel via interface B segregation. International Journal of Hydrogen Energy, 2025, 164, pp.150954. 10.1016/j.ijhydene.2025.150954. hal-05244845

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# Improving hydrogen embrittlement resistance of martensitic steel via interface B segregation

Hao Shi<sup>a, *</sup> <sup>●</sup>, Guillaume Hachet<sup>a</sup>, Huijie Cheng<sup>b</sup>, Prithiv Thoudden Sukumar<sup>a</sup>, Dirk Ponge<sup>a, ●</sup>,
Binhan Sun<sup>b,c,**</sup>

## ARTICLE INFO

Handling editor: Mehran Rezaei

## Keywords:

Hydrogen embrittlement (HE)

Boron segregation

Lath martensitic steel

Intergranular (IG) fracture

Quasi-cleavage (QC) fracture

## ABSTRACT

In this work, we find that the engineering grain boundary boron (B) segregation can effectively enhance the resistance to hydrogen embrittlement (HE) of a low-carbon lath martensitic steel. The B segregation in prior austenite grain boundaries (PAGBs) strongly suppresses H-induced intergranular cracking. Such a beneficial role, associated with B segregation, is interpreted as the suppression of H-trapping at PAGBs and the enhancement of the atomic bonding strength of grain boundaries. Despite the improvement in HE resistance, an H-induced quasi-cleavage fracture still occurs in the B-doped samples. Further detailed fractography analysis, conducted using electron backscatter diffraction and transmission electron microscopy, reveals that the primary H-induced damage mode in the B-doped samples is {110} plane cracking, resulting from a combined effect of H-enhanced decohesion and H-enhanced localized plasticity. The obtained results advance the understanding of HE in lath martensitic steel with solute segregation and provide a strategy for mitigating such an issue.

## 1. Introduction

High-strength lath martensitic steels possess excellent mechanical properties, making them attractive materials for multiple applications (automotive, fastener industries, etc.) [1,2]. Low- and medium-carbon steels are materials with excellent mechanical properties (strength, ductility, toughness) which can be achieved via optimizing chemical composition and heat treatment [3,4]. Lath martensite is a typical microstructure obtained by quenching these types of steels [5]. Generally, four types of hierarchical substructures, depending on the length scales, exist in a lath martensitic structure, including laths, blocks, packets, and prior-austenite grains [6]. The lath is the smallest unit with a thickness varying from 50 to 800 nm [5,7]. A block is composed of a number of laths with nearly identical crystallographic orientation, and a packet contains blocks having the same habit plane. A prior austenite grain represents the grain inherited from the parent austenite phase.

However, it is well known that this type of steel often suffers from severe hydrogen embrittlement (HE) issues even with a few parts-per-

million in weight of H introduced into the material [8–11]. Among the different forms of H-induced failure, two types of damage behavior, namely intergranular (IG) and quasi-cleavage (QC) fracture, are often observed in lath martensitic steel [12]. Intergranular fracture is deduced from H-induced crack propagation along grain boundaries (e.g., prior-austenite grain boundaries (PAGBs) in martensitic steel). Several mechanisms have been proposed to interpret such a cracking behavior [13–18]. One interpretation is related to the decohesion of grain boundaries (GBs) induced by the co-segregation of H with other detrimental impurities such as P or S [13]. Preferential trapping of H at GBs was believed to be caused by a relatively high H binding energy (BE) of GBs (i.e., 30–45 kJ/mol) [14]. Additionally, mobile dislocations, when blocked by interfaces, can further transport H to these obstacles (i.e., GBs) [19]. The associated accumulation of H weakens the atomic bonding of GBs and increases their susceptibility to decohesion [15–17]. Another proposed explanation is that H can enhance dislocation pileups at GBs, which increases the stress concentration at GBs and causes it to exceed the cohesive strength of GBs [17,18]. On the other hand,

<footnote><sup>a</sup> Microstructure Physics and Alloys Design Department, Max Planck Institute for Sustainable Materials, 40237, Düsseldorf, Germany</footnote>

<footnote><sup>b</sup> Key Laboratory of Pressure Systems and Safety, Ministry of Education, School of Mechanical and Power Engineering, East China University of Science and Technology, Shanghai, 200237, China</footnote>

<footnote><sup>c</sup> State Key Laboratory of Chemical Safety, East China University of Science and Technology, Shanghai, 200237, China</footnote>

<footnote>* Corresponding author.</footnote>

<footnote>** Corresponding author. Key Laboratory of Pressure Systems and Safety, Ministry of Education, School of Mechanical and Power Engineering, East China University of Science and Technology, Shanghai, 200237, China.</footnote>

<footnote>E-mail addresses: shisjtu@outlook.com (H. Shi), binhan.sun@ecust.edu.cn (B. Sun).</footnote>
quasi-cleavage fracture describes a type of transgranular brittle fracture behavior that occurs on non-typical cleavage planes accompanied by fine facets (i.e., tear ridges/serrated markings) [8,19,20]. Previous studies of lath martensitic steels via TEM indicated that H-induced transgranular micro-cracking of the associated QC fracture propagated along lath boundaries [8,16,19]. Recent publications indicated that the QC fracture plane of lath martensitic steel was along {110}α slip planes [21–24]. However, it is yet unclear whether such microcracks propagate along the lath boundaries or within the laths since the longitudinal direction of some laths is also along {110} planes.

Several strategies were proposed to improve the HE tolerance of martensitic steels. These approaches are mainly focused on: (i) grain refinement [25], (ii) introducing H-trapping nanoprecipitates [26–30], (iii) introducing nano-films of inter-lath austenite [31,32], and (iv) carbon segregation at PAGBs [33]. Recently, B segregation at grain boundaries was considered an effective method to enhance the environmental damage tolerance of different types of materials [34–36]. One study performed by Guo et al. [34] indicated that B addition enhanced the HE resistance of Fe–Ni-based austenitic steel via suppressing the precipitation of η-precipitates (Ni₃Ti) at GBs, namely, 52.4 % loss of reduction of the sample’s cross-sectional area (B-free sample, after tensile test) vs 38.5 % loss of reduction of cross-sectional area (60 wt ppm of B). Boron addition was also reported to improve the HE resistance of the CrCoNi alloy [35]. The fracture mode changed from a brittle IG type (B-free) to a ductile transgranular (TG) type (B-doped). Relevant explanations include B segregation increasing GB’s cohesivity and strength, increasing the “Griffith work” of brittle boundary decohesion, preventing H diffusivity along GBs or-/and-relaxing dislocation pileups because of the easy transfer between neighbor grains [34, 36–38]. For martensitic steel, recent work showed a repulsive interaction between B and H at PAGBs [39]. However, considering the multiple interfaces existing in the microstructure and the HE susceptibility of martensite steel [6,8,12], experimental evidence of the HE mitigation from B addition in this system is still lacking in the literature.

The motivation of the current study is thus to understand the role of B segregation (at PAGBs) on HE resistance and H trapping behavior in a lath martensitic steel. Two low-carbon (Fe-0.2 wt% C) martensitic steels, one doped with B and the other devoid of B, were prepared for the study. The B-segregation at the grain boundary was characterized by atom probe tomography (APT). The HE effects of both steels were evaluated by slow-strain-rate tensile tests. Moreover, H-trapping behavior of both materials was quantified using thermal desorption spectroscopy (TDS) to further elucidate the effect of B-segregation on H-trapping behavior on PAGBs. Then, the fracture surface morphologies and secondary microcracks (beneath the fracture surface) were characterized by electron backscatter diffraction (EBSD) and transmission Electron Microscopy (TEM) to unveil the mechanism of the H-induced brittle fracture.

## 2. Materials and experiments

### 2.1. Materials and processing

Two low-carbon steel compositions, one without B (referred to as LC) and the other doped with B (referred to as B-LC), were cast using a levitation melting method to minimize contamination. Table 1 presents the chemical compositions of both materials, measured by wet-chemistry analysis. The primary difference between the two steels is the concentration of B and Ti. For the B-LC, Ti was added to protect B from forming boron-nitrides [40]. Both samples were heat-treated using

a Bähr DIL805 dilatometer. The materials were first heated to 1100 °C at a heating rate of 10 K/s and then annealed for 30 s, followed by cooling in a He gas atmosphere. The martensite starting temperature (M₅) of LC and B-LC steels, as determined from the recorded dilatation curves (shown in Fig. 1), is 440 °C and 460 °C, respectively. Adding B was reported to increase the hardenability of steels by delaying the ferrite nucleation [40].

### 2.2. H pre-charging, detection, and tensile tests

Hydrogen pre-charging was performed by an electrolytic charging cell with an aqueous solution of 0.1 M H₂SO₄ plus 0.3 wt% NH₄SCN. The charging procedure was performed in a three-electrode mode (with a silver chloride reference electrode), where the sample served as the working electrode, and a Pt foil (10*10 mm²) was used as a counter electrode. The current density was –12.5 A/m², and charging durations were 6 h and 15 h, respectively. The depth of the H-penetrated zone was estimated by the relation $\sqrt{D \cdot t}$ [41,42], where D is the H diffusion coefficient and t is the charging duration. According to the measured effective H diffusion coefficient [39], the H-penetrated zone could be around 3.5 mm (LC) and 4.6 mm (B-LC) with 6 h charging, which is larger than the sample thickness. All the sample surfaces were ground with a grit surface finish of 2400 before H charging. Charged samples were cleaned with alcohol immediately after being extracted from the electrolytic solution for other tests.

Thermal desorption spectroscopy (TDS) measurements were performed using a G8 GALILEO apparatus from Bruker®. The amount of H trapped in all steels was determined by temperature ramping with a mass spectrometer. The sample size is 10*10*1 mm³. After H pre-charging, the cleaned sample was inserted into the TDS analysis chamber. Then, the sample was heated in an infrared furnace from 298 K to 1173 K at a heating rate of 0.3 K s⁻¹. The time interval between the end of the H charging and the beginning of the temperature ramping was below 5 min for all samples.

The HE susceptibility of the two materials was evaluated using slow

<figure><img src="image_4.png" /><figcaption>Fig. 1. Dilatation curves of LC and B-LC steels obtained by dilatometry during the heat treatment.</figcaption></figure>

<table><caption>**Table 1**

Chemical compositions of both martensitic steels measured by a wet-chemistry analysis (wt.%).</caption><thead><tr><th>Sample</th><th>Fe</th><th>B</th><th>C</th><th>Mn</th><th>N</th><th>P</th><th>S</th><th>Si</th><th>Ti</th></tr></thead><tbody><tr><td>LC</td><td>Bal</td><td>&lt;0.0005</td><td>0.149</td><td>1.50</td><td>0.002</td><td>&lt;0.001</td><td>0.0033</td><td>0.374</td><td>&lt;0.0005</td></tr><tr><td>B-LC</td><td>Bal</td><td>0.0024</td><td>0.147</td><td>1.46</td><td>0.0017</td><td>&lt;0.001</td><td>0.0017</td><td>0.400</td><td>0.0267</td></tr></tbody></table>
strain rate tensile (SSRT) tests with a strain rate of $8 \times 10^{-5} s^{-1}$. It was conducted using a Kammrath & Weiss stage coupled with the digital image correlation (DIC) technique. The gauge part of the tensile samples has a length of 4 mm, a width of 2 mm, and a thickness of around 1 mm. The interval between the end of H pre-charging and the start of tensile testing was less than 15 min. The elongation loss (EI) value, calculated by the following equation (1), is introduced to reveal the brittleness of the H-charged sample.

$$EI(\%) = \frac{\varepsilon_{unch \ arg \ ed} - \varepsilon_{ch \ arg \ ed}}{\varepsilon_{unch \ arg \ ed}} \times 100 \quad (1)$$

Where $\varepsilon_{unch \ arg \ ed}$ and $\varepsilon_{ch \ arg \ ed}$ are engineering strains of uncharged and H pre-charged samples at failure, respectively.

## 2.3. Microstructure characterization

The microstructure and fracture surfaces of steels with and without H were characterized by a ZEISS Sigma scanning electron microscope (SEM). Quantification of grain size, orientation, and martensite sub-microstructures was conducted with electron backscatter diffraction (EBSD) in a ZEISS Sigma SEM. It worked in a condition with an accelerated voltage of 15 kV, a beam current of 2 nA, and a step size of 50 nm. The acquired data was analyzed by the TSL-OIM v8.6 software package. Electron channeling contrast imaging (ECCI) was applied to characterize the microstructure of lath martensite by a Zeiss-Merlin high-resolution SEM equipped with a 4Q backscattered electron detector, operated in a voltage of 30 kV, a beam current of 2 nA and a working distance of 8 mm. Transmission Electron Microscopy (TEM) analysis of the fractured surface region was performed by a FEI Talos F200X. The thin lamella used for the TEM characterization was lifted out from the probed region by a FEI Helios 600 crossbeam focused ion beam (FIB) operated at 30 kV and an ion current of approximately 24 pA. The surface was protected by

depositing a Pt layer (2 µm thick). Atom probe tomography (APT) was also conducted to investigate boron segregation at grain boundaries. This was performed on a LEAP™ 5000 XR, operating at a base temperature of 60 K, with a detection rate of 1–1.5 % (10–15 ions per 1000 pulses), laser pulsing with a repetition rate of 200 kHz, and pulse energy of 30 pJ. X-Ray diffraction analysis (XRD) was performed by a Bruker D8 advance diffractometer with Cu Kα1 radiation ($\lambda$ = 0.15406 nm, 40 kV and 30 mA) in Bragg–Brentano geometry (θ-2θ) mode. The parameters used for the measurement were 2θ varying from 40° to 120° with a step size of 0.02° and 1 s collecting time. The heat-treated samples used for XRD measurements were ground to a 1200-grit surface finish.

# 3. Results

## 3.1. Material microstructure and B segregation

Fig. 2 (a, b) show IPF images of the two steels after the annealing treatment. Irrespective of B addition, both samples have similar prior-austenite grain size with an average value of 25 ± 2 µm. Detailed features related to the hierarchical martensitic structure are displayed in Fig. 2c. Multiple interfaces, including block boundaries (black line), packet boundaries (PBs, green line), and PAGBs (white line), are identified according to the reconstructed PAGs, as marked and shown in the image. Child variants distribution (Supplementary Fig. S1) and their orientation relationships are described elsewhere [5]. Since the orientation between neighboring laths within a block is less than 10°, the detailed structure of individual laths was further characterized by electron channeling contrast imaging (ECCI), as shown in Fig. 2d. Parallel lath boundaries and high density of dislocations close to laths were observed. The average width of the laths is around 0.5 µm.

Fig. 3a shows the distribution of misorientation angle of both samples, obtained by EBSD measurement. The major misoriented angles are located at 0–10° and 50–60°, where low angles are attributed to lath

<figure><img src="image_5.png" /><figcaption>Fig. 2. (a, b) Combined images of inverse pole figure (IPF) and image quality (IQ) of LC and B-LC martensitic steels, analyzed by OIM; (c) Magnified IPF image of lath martensitic structure, the black lines represent block boundaries, the green lines are packet boundaries (PBs), and the white lines are PAGBs; (d) ECC image of lath microstructure. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)</figcaption></figure>
<figure><img src="image_6.png" /><figcaption>Fig. 3. (a) Distribution of misorientation angle of LC and B-LC steels, the data was obtained from EBSD analysis of annealed samples, the red double-headed arrow marks the range of typical misorientation angles of PAGBs, and the green double-headed arrow marks the range of block boundaries (BGs) and packet boundaries (PGs); (b) APT characterization of a PAGB in B-LC sample. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)</figcaption></figure>

boundaries [43] and high angles are mainly from block, packet, and prior austenitic grain boundaries [5,43]. Regardless of B addition, both steels exhibit very similar distributions of misorientation angles, indicating that B doping does not influence the sub-microstructures of martensite. In addition, atom probe tomography (APT) characterization of a PAGB of the B-LC (Fig. 3b) confirmed the segregation of B with a

maximum concentration of 1.6 at.%. Considering the average matrix B composition (0.016 at.% B) near the GB, the grain boundary excess of B, called $\Gamma_B$, calculated based on the integral method [44,45], is equal to 8.1 ± 0.69 atoms/nm². Boron segregation is primarily induced by equilibrium- and non-equilibrium segregation [46,47]. The equilibrium-segregation dominates at low temperatures (<1000 °C),

<figure><img src="image_7.png" /><figcaption>Fig. 4. (a) Engineering stress-strain curves of uncharged and H pre-charged LC and B-LC steels, H charging duration: 6 h and 15 h; (b) Plot of elongation losses ($E_l$) of both materials as a function of charging duration; (c, d) Local engineering strain of LC and B-LC samples before the fracture point as a function of sample distance (the insets are corresponding strain maps).</figcaption></figure>
where it is assumed that solute atoms move from the bulk material to the loosely packed sites (e.g., GB, phase boundary, precipitates/matrix interface) due to the reduced grain boundary free energy [46,48]. Instead, non-equilibrium segregation plays a role at high temperatures (>1000 °C) due to the increased number of vacancies that drag B atoms to the GBs during fast cooling in terms of vacancy-B complex [46–50].

## 3.2. Evolution of the HE susceptibility with B incorporation

The tensile tests of both materials were conducted in H-free and H-pre-charged conditions. Fig. 4a plots the engineering stress-strain curves of LC and B-LC steel in both conditions. Without H, the total engineering strain of LC and B-LC samples is 20.0 % ± 0.3 % and 17.7 % ± 0.1 %, respectively. The yield stress (defined here as the 0.2 offsets ($\sigma_{0.2}$)) and ultimate tensile strength (UTS) of LC and B-LC steels are similar with and without B ($\sigma_{0.2}^{LC} = 1008 \pm 59$ MPa and $\sigma_{0.2}^{B-LC} = 1018 \pm 4$ MPa while UTS$^{LC} = 1245 \pm 12$ MPa and UTS$^{B-LC} = 1250 \pm 2$ MPa). These values indicate that the tensile properties in the absence of H are marginally influenced by B doping. When the two steels were pre-charged for 6 h, they exhibited a reduction in total elongation (i.e., ductility), with 8.8 % ± 1.6 % for LC and 12.5 % ± 1.7 % for B-LC. As the charging duration increases to 15 h, the total elongation of the LC is significantly reduced to 3.5 % ± 0.6 %. Although the elongation loss also exacerbates in the B-LC samples (7.6 % ± 2.3 %), it remains twice as large as that of LC. To reflect the HE susceptibility of both samples, the relation of EI values vs charging duration was plotted in Fig. 4b. The LC samples always show a larger EI value compared to the B-LC samples under the same charging duration, indicating their higher HE susceptibility. Fig. 4c and d compare the local engineering strain in the necking area, just before the failure of both steels. The maximum local engineering strains are 15

% (LC, 6 h), 7 % (LC, 15 h), 51 % (B-LC, 6 h), and 18 % (B-LC, 15 h). The maximum strain of the LC samples is always lower than that of the B-LC samples, implying a significant loss of local plastic deformation ability in the B-free samples after H charging. The results obtained indicate that B addition and the associated segregation can significantly improve the HE resistance of low-C martensitic steels. The values of the tensile properties of both samples are summarized and listed in the Supplementary Table S1.

## 3.3. The influence of B segregation on H trapping behavior

The multiple interfaces existing in lath martensite can trap H and influence the material's damage tolerance capacity [23,51]. To study the H trapping behavior, both steels were subjected to TDS measurement after H pre-charging with a constant heating rate of 0.3 K s⁻¹. Fig. 5(a–d) present the H-desorption rate ($\Phi$) of both steels as a function of temperature. An intensified H desorption occurs in the temperature range from 300 K (the starting temperature of the measurement) to 573 K. The signals from TDS correspond to H desorption from different trapping sites, which depend on the segregation energy between H and the crystal defects. To separate the individual H trapping sites, a gradient-conjugate algorithm was employed to minimize the absolute difference between the spectra and a function that is the sum of several Gaussian distributions. Based on this approach, the main peak from LC (black color) can be further deconvoluted into four sub-peaks, while the spectrum from B-LC can be divided into three sub-peaks (purple color), in agreement with the previous work [39]. Each peak represents an H-trapping site with a different binding energy (BE). The H partitioned in each sub-peak was quantified based on the area underneath the curve. For all measurements by temperature ramping, the H integration detected by peak integration is similar (2.2 ± 0.6 wppm), shown in Fig. 5e. The amount of

<figure><img src="image_8.png" /><figcaption>Fig. 5. TDS spectrum of LC and B-LC after 6 h and 15 h of H pre-charging, respectively, (a) LC sample with 6 h charging; (b) B-LC sample with 6 h charging; (c) LC sample with 15 h charging; (d) B-LC sample with 15 h charging; peak 1: blue curve, peak 2: red curve, peak 3: green curve, peak 4: orange curve; (e) The total amount of trapped H in each charged sample; (f) Hydrogen partitioning in different defects in LC and B-LC after 6 and 15 h of H pre-charging, respectively. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)</figcaption></figure>
H in each trapping site, deduced from peak integration, is summarized in Fig. 5f. The distributions of H in peak 1 to 3 (blue, red, and green curves in Fig. 5(a-c)) do not vary significantly in the LC samples with the different charging durations. However, H trapped in the trapping site related to peak 4 (orange curve) increases significantly as the charging duration increases. For the B-LC samples, a higher fraction of H was trapped in peak 1, with less H in trapping sites related to peaks 2 and 3 compared to the LC samples. Peak 4 is absent from the B-LC samples for both charging durations. As the charging duration reaches up to 15 h, the total amount of trapped H increases as well.

The identification of different traps associated with the different TDS peaks in steels was discussed based on information available from previous works [39,51–60], summarized in Table 2. The first three peaks (peak 1 to 3) are attributed to lattice diffusible H, H trapped in dislocations, and martensitic interfaces (i.e., lath/block/packet boundaries). An additional peak (peak 4) was only obtained for LC steel when the same TDS measurements were performed. Since boron repels the trapping of H in PAGBs [39,59], this peak should correspond to H trapped in PAGBs. The result aligns with previous work, which showed that the binding energy of H trapped in HAGBs is larger than that of H trapped in dislocations of steel and iron [54–56].

### 3.4. Fracture surface and sub-microcracks characterization

Further investigation of the fracture surfaces of the studied steels with and without H pre-charging was performed. Without H, both materials display a ductile fracture with dimples varying between 1 µm and 4 µm (the fracture surfaces are shown in the Supplementary Fig. S2). After charging, the fracture surfaces of both samples show some different features. Fig. 6a shows the fracture surface view of the LC sample after 6 h charging. Although most of the surfaces remain ductile, a brittle fracture is visible within the region outlined by the white dashed line. Fig. 6e shows the area that contains intergranular (IG, smooth and integral facets) and quasi-cleavage (QC, “cleavage-like” facets with tear ridges/serrated markings) fracture features. The brittle area fraction (BF), defined as the ratio of the brittle fractured area (IG and QC fractures) to the whole fracture surface area, is around 10% for this sample. When the charging time increases (i.e., 15 h charging), the brittle area also enlarges (Fig. 6b and f), as indicated by the BF value increasing to 27%. In contrast, the B-doped samples display distinct differences in surface morphologies (Fig. 6c and d). Almost the entire fracture surface of this sample after 6 h charging still presents a ductile dimpled behavior, with only a few exceptional regions displaying a QC-like feature (marked in Fig. 6g). When the charging duration increases to 15 h, a small area with QC facet, accounting for 11% of the whole surface area, is identified at the sample edge, as shown in Fig. 6h. Nevertheless, no evidence of IG fracture is identified in the B-doped samples for both charging conditions. The observed results indicate that

<caption>**Table 2**</caption>

The reported H trapping sites and corresponding TDS peak positions, along with the binding energies.

<table><thead><tr><th>H trapping site</th><th>TDS peak position (Temperature, K)</th><th>Binding energy (kJ/mol)</th></tr></thead><tbody><tr><td rowspan="3">Lattice interstitials</td><td>320 K [39];</td><td>15.4 [39];</td></tr><tr><td>303 K [57]</td><td>13.5 [51];</td></tr><tr><td>348 K [53];</td><td>32.8 [39];</td></tr><tr><td rowspan="2">Dislocations</td><td>433 K [57];</td><td>18.0 [53];</td></tr><tr><td>395 K in single crystal Fe [60];</td><td>25.7 [18];</td></tr><tr><td rowspan="2">Martensite boundaries (lath/packet/block)</td><td>488 K in pure Fe [58];</td><td></td></tr><tr><td>398 K [53];</td><td>25.1 [39]</td></tr><tr><td rowspan="4">Grain boundaries</td><td></td><td>35.0 [53];</td></tr><tr><td>385 K in pure Fe [58];</td><td>35.7 [39]</td></tr><tr><td>415 K in polycrystalline Fe [60];</td><td>47.4 [56];</td></tr><tr><td>373 K [57];</td><td>57.4 [16];</td></tr></tbody></table>

B-addition modifies the fracture mode of the H pre-charged lath martensitic steel and completely suppresses the H-induced IG fracture.

The secondary cracks induced by H, which are located underneath the fracture surface, were investigated to unravel the HE mechanisms of the different systems. Fig. 7 presents the sub-surface of the LC steel charged with hydrogen (6 h). A few microcracks, visible in dark contrast, are evident from the SEM image. To further illustrate the spatial distribution of these microcracks, an image of parent-to-child orientation (Fig. 7b), reconstructed by the MTEX program according to the Kurdjumov-Sachs (KS) orientation relationship [61], confirmed the nucleation and propagation of microcracks along the PAGBs. Besides, other high-angle boundaries, including block (black line) and packet boundaries (green line), are also identified and marked in the IPF image of Fig. 7c. These results further support the formation of IG fracture of the LC steel from the micro-crack nucleation and propagation perspective. In addition, the Kernel average misorientation (KAM) map (Fig. 7d) also unveils large misorientations (red color) in the PAGB region, which was quantified by the KAM line profile in Fig. 7e (following the red arrow in Fig. 7d). The result reveals a larger KAM value of the probed PAGB, nearly two times more than the other measured region.

Though sample B-LC (6 h charging) shows a ductile fracture surface morphology, a 20.2% reduction of the total elongation was still observed. The H-induced ductility loss is assumed to be associated with the occurrence of QC fracture in a few local regions (see another example shown in Fig. 8a). To further study the mechanisms underlying such H-induced fracture behavior, the fracture surface along with the adjacent regions were characterized from the cross-sectional point of view (Fig. 8b). A Ni-layer (10–15 µm) was deposited on the fracture surface by an electrical chemical process to prevent the mechanical damage of the original fracture surface induced by sample polishing process. Most fracture surface regions exhibit a jagged cracking path observed from the cross-sectional view, corresponding well to the ductile dimpled fracture morphology [12,21]. However, a few small regions close to the sample edge are labeled by a relatively flat cracking trace, corresponding to the QC-typed facets observed from the fracture surface and displayed in Fig. 8b. A few microcracks were also observed beneath the brittle fractured surface, marked in Fig. 8c. Fig. 8(d–h) show the EBSD scans covering the brittle area and microcracks demonstrated in Fig. 8b and c. Generally, three features can be obtained from the EBSD characterization. First, the boundaries at some regions (i.e., Reg A, underneath the fracture surface) are nearly parallel to the fracture surface, as shown by the magnified image in Fig. 8e. In addition, a scan of the H-induced microcracks (marked in Fig. 8c and h) reveals that the curved crack traces follow martensitic boundaries. Similar features were also observed in the B-LC sample after 15 h charging (Supplementary Fig. S3). Additional investigations were performed using TEM due to the extremely fine size of the lath substructures (less than 100 nm in width), induced by severe plastic deformation during the tensile test, and the results are presented below. Secondly, the fracture surface in some regions (i.e., Reg B in Fig. 8d) is nearly perpendicular to the different boundaries. A local analysis of the fracture surface trace (see Fig. 8f) indicates that the fracture surface is parallel to the {100} plane according to the {100} pole figure analysis in Fig. 8g, as also observed by Cho et al. [22].

Then, a TEM lamella was lifted out from the QC region of B-LC steel with H pre-charged for 15 h (details are shown in Supplementary Fig. S4) to further investigate the H-induced brittle fracture behavior described above. Fig. 9 shows the fracture surface trace (blue dashed line) deposited with a Pt layer above and the lath microstructure beneath the fracture surface. In some regions, the fracture surface is parallel to the laths underneath (i.e., lath IV, VII), while in some areas the fracture surface also cuts through the laths (i.e., lath III, V). According to the selected area electron diffraction (SAED) performed on the neighboring laths (II-III, III-IV, VI-VII (Supplementary Fig. S5)) or individual lath (VII), most of the boundaries among the laths are low-angle boundaries (<10°), belonging to lath boundaries [5], which was
<figure><img src="image_9.png" /><figcaption>**Fig. 6.** SEM fracture surface morphologies of LC and B-LC samples after H pre-charging; (a) and (b) are overview images of LC steels after 6 h and 15 h pre-charging; (c) and (d) are overview images of B-LC steels after 6 h and 15 h charging; (e) and (f) are magnified images taken from (a) and (b), respectively; (g) and (h) are magnified images taken from (c) and (d), respectively.</figcaption></figure>

<figure><img src="image_10.png" /><figcaption>**Fig. 7.** (a) SEM Image quality (IQ) of microcracks formed beneath the fracture surface of LC sample after 6 h of H pre-charging; (b) Reconstructed PAGBs of (a), obtained from the MTEX program [61], the black lines represent the PAGBs, the white areas are un-identified regions due to microcracks and severe plastic deformation; (c) Inverse pole figure (IPF) of (a), analyzed by OIM, the black lines are block boundaries, the green lines are packet boundaries (PBs) and the white lines are PAGBs; (d) Combined Kernel average misorientation (KAM) and IQ image of (a); (e) The KAM line profile obtained from the red arrow marked in (d) across a PAGB. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)</figcaption></figure>

also confirmed by TKD analysis (Supplementary Fig. S6). In parallel regions, the fracture surface plane, determined by the SAEDs performed on the different laths underneath, is identified as {110} planes, as marked by the short red lines in the figure. In the inclined regions, the fracture surface does not follow the {110} planes. There are also two exceptional positions (marked with yellow arrows) where the trace line is almost perpendicular to other places. In addition, tear ridges (indicated by the red arrows) were also observed on the entire fracture surface trace, exactly located at the regions where the fracture surface trace encounters a martensite boundary. Such features were also reported in previous works [22,24]. It claimed that the formation of such features was due to crack deflection induced by martensitic boundaries [22].

Fig. 10 shows an example of the interaction between the fracture surface and a martensitic boundary where the crack cuts through the lath III and a tear ridge forms on the interface of lath III and lath IV. The crack line trace (blue dashed line) and the lath boundary (yellow dashed

line) are indicated. As the crack approaches and encounters a lath boundary during propagation, its original propagation direction is altered, and a tear ridge forms at the intersection point. After passing through the intersection point, the crack again propagates along the encountered lath boundary.

# 4. Discussion

## 4.1. Enhancement of the HE resistance with B addition

The SSRT tests of LC containing H show clear evidence of IG fracture along PAGBs. According to the literature, this type of fracture was either due to a weakening of the atomic bond strength, in agreement with the H-enhanced decohesion (HEDE) model [9,62–64] or/and an increase of the pile-up stress from an enhancement of dislocation motion, which is localized the plasticity, in agreement with the hydrogen-enhanced
<figure><img src="image_11.png" /><figcaption>**Fig. 8.** (a) SEM fracture surface of *B-LC* sample after 6 h of H pre-charging; (b) Cross-sectional image of the fracture surface, the flat fracture trace corresponds to the QC region in (a); (c) Microcracks underneath the fracture surface; (d) Combined IPF and IQ image of fracture surface trace, analyzed by OIM, PB (packet boundary); (e) Magnified image of Reg A in (d); (f) Magnified image of Reg B in (d); (g) {100} pole figure of lath martensite in Reg-B; (h) Combined IPF and IQ image of microcracks beneath the fracture surface.</figcaption></figure>

<figure><img src="image_12.png" /><figcaption>**Fig. 9.** TEM bright field (BF) micrograph of fracture graph formed on *B-LC* sample after 15 h of H pre-charging, the fracture surface trace (blue dashed line) and lath boundaries (yellow dashed line) were marked, the short red lines represent {110} planes trace, LB/BB/PB represents lath/block/packet boundary, respectively. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)</figcaption></figure>

localized plasticity (HELP) model [16,18,65]. This IG fracture was entirely suppressed in the presence of boron at PAGBs, and the disappearance of one peak was also observed in the H pre-charged *B-LC* steels during the H desorption analysis conducted using TDS.

According to trap models described in the literature [16,51,58], the total amount of H in a system corresponds to the sum of diffusible H in the lattice ($C_L$), H trapped into reversible traps ($C_{Tr}$) with a BE less than 40 kJ/mol, and H trapped in the irreversible traps ($C_{Ir}$) with a BE over 60 kJ/mol. When comparing the trapping behavior of H in *LC* and *B-LC*

steels, only the reversible trapping sites can be impacted, which are mostly residual dislocations, martensite boundaries, and PAGBs. Furthermore, the dislocation densities of *LC* and *B-LC* steels were estimated by XRD measurements (Fig. 11). Using the modified Williamson-Hall (WH) equation [66,67], the density of this line defect is $5.4*10^{15} \text{ m}^{-2}$ for *LC* and $3.8*10^{15} \text{ m}^{-2}$ for *B-LC*. The same level of dislocation density indicates both steels should trap a similar amount of H by dislocation. In addition, a small peak corresponding to the (111) plane of γ-Fe (austenite) was identified in both materials by the XRD
<figure><img src="image_13.png" /><figcaption>**Fig. 10.** TEM bright field (BF) micrograph of fracture graph and tear ridge formed on B-LC sample after 15 h of H pre-charging, the fracture surface trace (blue dashed line) and the lath boundaries (yellow dashed line) were marked in the figure. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)</figcaption></figure>

measurements. Since the volume fraction of γ-Fe is less than 2 % for both steels [39], its contribution to H-trapping can be neglected.

Moreover, the EBSD IPF images (Fig. 2(a and b)) reveal that both samples possess a similar PAG size, which precludes the impact of B segregation on martensite nucleation and growth, in agreement with the literature [68]. Moreover, the similar misorientation distribution of both steels (Fig. 3a) indicates that both materials have nearly the same ratio of each type of boundary. While having similar microstructures, we observed that the main difference between LC and B-LC is the segregation of B in PAGBs. Since the fracture surface of B-doped steel shows the apparition of secondary cracks at crystal planes/boundaries that are not PAGBs, it can be concluded that boron prevents the H-induced IG fracture via modifying the H-trapping behavior. Such a phenomenon could be interrupted by using *ab initio* calculations to determine the binding energy between H and B at GBs in α-Fe [39,59]. While both interstitials

are attracted to the core of the crystalline defect, boron has a stronger interaction with GB, and the interaction between H and the B-doped GBs becomes repulsive [39,59]. Besides, B-segregation was also reported to improve the cohesivity and bonding strength of GB, as B enhances the formation of covalent bonds [36,69].

## 4.2. Damage mechanisms of H pre-charged systems

As another type of brittle fracture, quasi-cleavage (QC) fracture still occurred in the B-doped steels. The fracture mechanism is different from that of the IG fracture. Previous studies have shown that QC propagated along {110} planes, as this plane is the dominant slip plane in the BCC structure. The hydrogen-enhanced dislocation motion (HELP mechanism) carried H and resulted in H accumulation on the slip plane, causing local slip plane decohesion or microcrack nucleation. On the other hand, local plastic strain analysis by Shibata et al. [23] clarified that the largest plastic strain along the habit plane or block boundary led to its decohesion, as the deformation bands observed were parallel to the habit plane. Another study conducted by Okada et al. [24] indicated that the resolved normal stress impinged on {110} planes was responsible for its decohesion. According to the work of Okada et al. [24], the microcracks originated from the micro-void/vacancy's coalescence induced by jog-dragging of screw dislocations along {110} planes. In our study, both fracture surfaces parallel to lath boundaries and cutting through the lath boundaries were observed via TEM characterizations. The parallel regions could be due to decohesion of lath boundaries or {110} planes as the laths within a packet are separated by highly dislocated, low-angle boundaries that tend to follow {110} planes [19]. However, the cutting regions neither follow the {110} planes nor the lath boundaries. In this case, a tear ridge formed at the intersection point between the boundary and crack trace, resulting in the deflection of the crack from the original propagation direction. The formation of the tear ridge is attributed to the barrier effect of martensitic boundaries, which hinders the propagation of microcracks. Additionally, the curved fracture crack (Fig. 10) and high dislocation density underneath the fracture surface also indicate the plastic deformation ability of the H-charged B-LC samples. Previous TEM observations also confirmed the H-enhanced dislocation mobility in QC fractured regions [70,71]. Therefore, the fracture mode in such a condition is {110} slip planes de-cohesion accompanied by hydrogen-enhanced plasticity.

Besides, both EBSD and TEM analyses reveal that the fracture surface in some regions is nearly perpendicular to the laths, corresponding to a {100} plane de-cohesion, which is the typical cleavage plane of BCC structures at cryogenic conditions [8,72]. Cottrell's dislocation-based model was often employed to understand the observed cleavage along the {100} plane [73,74]. In this theory, two intersecting {110}

<figure><img src="image_14.png" /><figcaption>**Fig. 11.** XRD results of LC and B-LC samples after the annealing treatment.</figcaption></figure>
<111>-type of dislocation systems move and form sessile dislocations on {100} planes. As the number of sessile dislocations increases, microcracks start to nucleate at {100} planes due to the reduced surface energy of {100} planes by brittle H media. Based on the aforementioned mechanisms, H-induced IG (for LC) and QC fracture mechanisms (for LC and B-LC) are schematically illustrated in the following Fig. 12.

The results presented above provide an efficient and cost-effective strategy to enhance the HE tolerance of lath martensitic steel through grain boundary segregation engineering. Hydrogen embrittlement tests confirmed that B-segregation can efficiently impede the H-induced grain boundary decohesion. However, the H-induced QC still occurred and resulted in a reduction of ductility due to the multiple interfaces existing in the martensitic structure, which are also potential trapping sites for H atoms. One possible solution to further inhibit the H effect on other interfaces could be B-segregation engineering on lath/block/packet boundaries. Further investigation is needed to understand the role of B segregation (i.e., at lath/block/packet boundaries) in inhibiting HE of lath martensitic steel.

## 5. Conclusions

The role of boron on HE resistance of lath martensitic steel was investigated in the present work. It was observed that the segregation of B on PAGBs significantly improved the HE resistance of Fe-0.2C (wt.%) lath martensitic steel. The total elongation of the B-alloyed sample (B-LC) is almost twice that of the B-free sample (LC) after H pre-charging. Micro-cracks analysis reveals that the brittle failure of the H-charged LC

steel is caused by intergranular and quasi-cleavage types of fracture. The B-LC steel shows a ductile fracture after 6 h of H pre-charging and a quasi-cleavage type hydrogen embrittlement fracture after 15 h of pre-charging. Further fractograph characterization of the B-LC steels unveils that the fracture mechanisms are {110} plane decohesion accompanied by hydrogen-enhanced plasticity and {100} plane cleavage. Thermal desorption spectroscopy (TDS) analysis of both steels reveals the absence of H in trapping sites dependent on HAGBs, the nature of PAGBs in B-LC steel, which is attributed to B segregating at these GBs. By comparing the hydrogen partitioning in the microstructure through peak integration and fracture mechanism analysis, it can be concluded that B-segregation repels H trapping on PAGBs, thereby mitigating the H-induced GB damage.

## CRediT authorship contribution statement

**Hao Shi:** Writing – review & editing, Writing – original draft, Methodology, Investigation, Conceptualization.
**Guillaume Hachet:** Writing – review & editing, Methodology, Investigation.
**Huijie Cheng:** Writing – review & editing, Methodology, Investigation.
**Prithiv Thoudden Sukumar:** Methodology, Investigation.
**Dirk Ponge:** Writing – review & editing, Validation, Methodology, Investigation, Conceptualization.
**Binhan Sun:** Writing – review & editing, Validation, Methodology, Investigation, Conceptualization.

<figure><img src="image_15.png" /><figcaption>**Fig. 12.** Schematic diagram of the mechanisms of H-induced damage, (a) IG fracture, (b) QC de-cohesion of {110} planes/lath boundary, (c) Cleavage along {100} planes.</figcaption></figure>
## Data availability

Data will be made available on request.

## Declaration of competing interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

## Acknowledgements

The authors gratefully acknowledge Prof. D. Raabe for the scientific support and fruitful discussions. Binhan Sun acknowledges the financial support from the National Key Research and Development Project (No. 2023YFB3712103), National Natural Science Foundation of China (Grant No. 52275147), Key Research and Development Program of Shandong Province (No. 2024CXPT064), State Key Laboratory of Long-life High Temperature Materials (DEC8300CG202210288EE280293) and Shanghai Gaofeng Project for University Academic Program Development. H. Shi acknowledges the funding from the National Natural Science Foundation of China (52401090). G.H. acknowledges the financial support from the Alexander von Humboldt Foundation, Germany, through grant no 3.3-FRA-1227460-HFST-E.

## Appendix. A Supplementary data

Supplementary data to this article can be found online at https://doi.org/10.1016/j.ijhydene.2025.150954.

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