A Comparative Study of the Tensile Behavior of Wrought 44W Steel, Monel 400, 304L Stainless Steel, and Arc-Directed Energy Deposited 308L Stainless Steel in Simulated Hydrogen Environments

Abstract

This study investigates the tensile behaviors of wrought 44W steel, Monel 400, 304L austenitic stainless steel, and arc-directed energy deposited (arc-DED) 308L austenitic stainless steel under simulated hydrogen environments to evaluate their endurance to hydrogen embrittlement (HE). The specimens were subjected to cathodic hydrogen charging in an alkaline solution, followed by uniaxial tensile testing at a strain rate of 0.2 min⁻¹. Based on measurements of elongation and toughness, the resistance to HE was ranked as follows: 304L stainless steel > Monel 400 > arc-DED 308L stainless steel > 44W steel. Notably, no significant changes were observed in the yield strengths, ultimate tensile strengths, or elastic modulus of 304L austenitic stainless steel, Monel 400, and 44W steel across all the levels of hydrogenation. However, the arc-DED 308L stainless steel exhibited a slight increase in these properties, attributed to its unique microstructural characteristics and strengthening mechanisms inherent to additive manufacturing processes. These outcomes contribute to a better understanding of the mechanical performance and suitability of these structural alloys in hydrogen-rich environments, highlighting the superior HE resistance of 304L stainless steel and Monel 400 for such applications.

1. Introduction

The mechanical performance of materials in hydrogen-rich environments is a critical area of research, particularly due to the detrimental effects of hydrogen embrittlement (HE) on structural integrity. HE is a phenomenon in which hydrogen atoms diffuse into a material, leading to a significant reduction in ductility and toughness [1,2]. It remains one of the most challenging issues, as it can lead to unexpected and disastrous failures in components that are otherwise considered structurally sound. The mechanisms underlying HE involve several complex processes that occur at the microstructural level, with their effects and traces often visible at both micro and macro scales. These processes can be broadly categorized into hydrogen entry (permeation), diffusion, interaction with microstructural features, and eventual embrittlement of the material [3,4,5].

Hydrogen can enter metals through various sources, including cathodic processes such as electroplating, corrosion, or cathodic protection, as well as from gaseous hydrogen sources like high-pressure hydrogen gas environments [6,7]. The solubility and degree of mobility of hydrogen in a material are influenced by several factors, including the material’s intrinsic properties, environmental conditions, and the specific method of hydrogen exposure [8].

Due to their small size, hydrogen atoms typically enter materials in atomic form (H) after molecular hydrogen (H₂) dissociates on the surface. These atoms then diffuse into the lattice of the material, occupying interstitial sites [9], where their mobility and potential for damage depend on the crystal structure. Hydrogen atoms can become trapped at various microstructural defects, such as dislocations, grain boundaries, vacancies, voids, or inclusions. These “trap sites” [10,11] can reduce hydrogen mobility but contribute to localized embrittlement over time.

In some metals, particularly titanium, zirconium, and their alloys, hydrogen can react with the metal matrix to form brittle metal hydrides (e.g., TiH₂, ZrH₂) [12,13]. These hydrides often precipitate along grain boundaries or in regions of high hydrogen concentration, significantly reducing the ductility and strength of the base metal. The brittle nature of these hydride phases makes them prone to cracking under stress, leading to failure via brittle fracture. Fracture surfaces affected by hydride formation often show features characteristic of brittle cleavage.

According to earlier studies [14,15,16], hydrogen has higher mobility in Body-Centered Cubic (BCC) structures than Face-Centered Cubic (FCC) structures. This is due to the relatively larger interstitial sites in BCC structures, such as octahedral and tetrahedral voids, compared to FCC structures. This, combined with the lower atomic packing factor (68% for BCC vs. 74% for FCC), creates a greater free volume within the lattice. In BCC materials, the relatively larger interstitial spaces and the atomic size mismatch make it difficult for permeated hydrogen atoms to remain in stable interstitial positions. Instead, these hydrogen atoms become highly mobile, leading to localized stress concentrations as they interact with the surrounding lattice. This mobility allows the hydrogen atoms to migrate from one point of damage to another, worsening the degradation of the material by creating pathways for crack initiation and propagation. This phenomenon is a key factor in HE observed in BCC materials, as it promotes widespread damage throughout the material.

In contrast, FCC materials possess smaller interstitial sites that can better confine hydrogen atoms. This restricts hydrogen diffusivity, limiting the mobility of hydrogen atoms within the lattice. As a result, the damage caused by hydrogen is more localized and confined to specific regions within the microstructure. The relatively high ductility and slip characteristics of FCC materials also enhance their ability to accommodate the stresses induced by hydrogen, thereby reducing the extent of embrittlement compared to BCC materials. This distinction in hydrogen behavior between BCC and FCC materials highlights their opposing proneness to hydrogen-induced degradation and is an important consideration in material selection for hydrogen-related applications.

Furthermore, grain size, dislocation density, and the presence of defects or precipitates can influence hydrogen solubility and distribution in metallic materials [17,18,19,20]. Finer-grain structures and higher dislocation densities provide more sites for hydrogen to occupy within a material. In materials with smaller grains, the increased grain boundary area provides more locations for hydrogen entrapment. This increased trapping capacity at the boundaries can lead to a reduction in the amount of free hydrogen within the grain interiors, as hydrogen tends to accumulate at the boundaries. However, although this may limit free hydrogen diffusion within grains, it also results in a higher overall hydrogen content in the material due to extensive boundary trapping.

This dual effect presents a trade-off: while refined grain structures can slow hydrogen migration through the lattice, they may simultaneously promote localized embrittlement at the grain boundaries where hydrogen accumulates. As a result, despite the apparent benefit of reduced intragranular hydrogen, the increased susceptibility of grain boundaries to hydrogen-induced cracking can still contribute to HE. Understanding this balance is essential in tailoring microstructures to improve resistance to HE in critical applications.

Alloying elements significantly affect hydrogen solubility in metals, either promoting or inhibiting uptake based on their interaction with the host matrix. For instance, the presence of nickel reduces hydrogen solubility in steels, while elements like manganese or silicon tend to increase it [21,22]. With respect to environmental conditions, it has been reported that higher partial pressures of hydrogen increase its solubility in a material [23], particularly under exposure to high-pressure gaseous hydrogen environments. Generally, hydrogen solubility increases with temperature [24], but the relationship can be complex due to competing factors like enhanced diffusion rates and potential microstructural transformations that may alter trapping behavior and solubility limits.

In aqueous environments, the pH of the electrolyte can affect the concentration of hydrogen atoms at the material’s surface. Cathodic conditions tend to promote higher hydrogen uptake due to the abundance of H⁺ ions and the enhanced hydrogen evolution reaction (HER) at materials’ surfaces. The high concentration of protons accelerates the HER, leading to increased hydrogen diffusion into the metal. This makes materials more susceptible to HE and corrosion under acidic conditions [25,26].

HE is a multifaceted phenomenon driven by multiple interacting mechanisms, with several key theories proposed to explain how hydrogen contributes to the loss of ductility and strength in metals, including hydrogen-enhanced decohesion (HEDE), hydrogen-enhanced localized plasticity (HELP), hydrogen-enhanced strain-induced vacancy (HESIV), and hydride formation [27,28]. The dominant failure mechanism varies depending on the material system, environmental conditions, and the nature of applied stress. Despite the fact that HE has been extensively researched, its underlying mechanisms remain complex and often vary from one material to the other.

HEDE theory proposes that hydrogen reduces the cohesive forces and the bonding strength between atoms in a metal lattice, especially at grain boundaries and interfaces [9,28,29]. Hydrogen atoms tend to diffuse into areas of high stress concentration, such as crack tips, grain boundaries, or inclusions. By accumulating at these points, hydrogen weakens the atomic bonds, lowering the critical stress required for crack initiation and subsequent propagation. As a result, cracks initiate and propagate more easily under tensile and fatigue stresses. This leads to brittle fracture as the crack propagates without significant plastic deformation. Fracture surfaces often appear intergranular or cleavage-like due to the weakening of grain boundaries.

The HELP mechanism [30,31,32], typically observed in hydrogen-charged materials subjected to cyclic or tensile loading, describes how hydrogen atoms interact with dislocations and reduce the barriers to their movement, thereby increasing their mobility. This results in localized plastic deformation around crack tips, creating regions of high strain that are more susceptible to crack growth. Although plastic deformation occurs, the ductility of the material generally decreases because the deformation is confined to small regions, which can result in microvoid formation and crack initiation, leading to premature failure [28]. It must be noted that crack propagation may appear ductile at a microscopic level but leads to brittle behavior at the macro scale.

In the HESIV formation theory [33,34], it is proposed that hydrogen facilitates the formation of vacancies due to strain, which eventually leads to void formation and crack growth. Under applied stress, hydrogen assists in the formation of vacancies by interacting with dislocations or other crystal defects. As these vacancies accumulate, voids or microcracks form, which eventually coalesce and lead to macroscopic crack propagation. This process typically results in ductile fracture that is influenced by hydrogen-assisted void formation, often envisioned under tensile loading. Fracture surfaces of materials undergoing HESIV often show porosity or voids.

Additionally, HE can result from hydrogen adsorption onto a metal’s surface [35], weakening atomic bonds and leading to surface cracking. The adsorbed hydrogen diffuses inward under stress, reducing surface energy and increasing susceptibility to crack initiation. Once a crack forms, hydrogen continues to diffuse along it, promoting further growth. This form of surface cracking is commonly observed in high-strength alloys, such as steels exposed to hydrogen-rich or aqueous environments, where corrosion plays a significant role.

To reduce carbon emissions, governments are promoting hydrogen–natural gas blending as a cleaner energy solution [36,37,38,39]. While this lowers fuel carbon footprints, it raises concerns about existing pipeline reliability, as most are made from steel alloys prone to HE. Studies show that hydrogen blending reduces fracture toughness [40,41] and fatigue resistance [42,43,44,45], increasing the risk of premature failure in pipeline steels under stress and cyclic loading. Due to its higher diffusivity compared to natural gas, hydrogen can permeate steel even at low concentrations, causing significant embrittlement over time [46,47]. These effects are intensified in high-pressure environments typical of gas transport lines, where hydrogen diffusion accelerates crack initiation and propagation, particularly at welds, grain boundaries, and stress concentration sites [38].

To address these risks, current research emphasizes the need for comprehensive material evaluations and potential pipeline modifications. Hydrogen-compatible materials, advanced coating technologies, and enhanced monitoring systems are being explored to mitigate embrittlement risks. Careful consideration of these challenges is crucial for the safe and efficient integration of hydrogen into existing natural gas infrastructure. While previous studies, including those by the authors of [48,49,50,51], have established that steels undergo significant HE under varying hydrogenating conditions, research on effective mitigation strategies remains limited and underexplored. Despite Monel and 304L stainless steel’s known resistance to hydrogen-induced degradation, there is a noticeable lack of comparative studies evaluating their tensile performance relative to both structural and corrosion-resistant steels in hydrogen-rich environments. Also, the arc-DED fabricated 308L austenitic stainless steel was employed in this work to aid understanding of the said mechanical behavior introduced by the additional layer of complexity due to the unique microstructural characteristics imparted by the additive manufacturing (AM) process. Unlike conventionally manufactured austenitic stainless steels, AM versions may exhibit anisotropy, directional grain growth, increased dislocation density, residual stresses, and non-equilibrium phases that could influence their hydrogen permeability and susceptibility to embrittlement. Preliminary studies on AM metals suggest that the layer-by-layer construction characteristic in AM can lead to a heterogeneous microstructure, potentially impacting hydrogen diffusion pathways and trapping sites, thereby altering the tensile properties under hydrogenating conditions. However, the extent to which these microstructural differences affect the hydrogen embrittlement resistance of arc-DED 308L austenitic stainless steel remains an open question in the field.

This study aims to envision how the tensile behavior of Monel compares with other types of structural and corrosion-resistant steels in hydrogen-rich environments, a topic that remains underexplored despite its potential applications in industries facing hydrogen-induced degradation challenges.

2. Materials

In this study, the tensile behavior of four materials, 44W steel, Monel 400, 304L austenitic stainless steel, and arc-DED 308L austenitic stainless steel, with dimensions of tensile test coupons as given in Figure 1a, was analyzed comparatively under tensile loading at varying hydrogen charging conditions. The chemical compositions of the studied alloys were measured using the inductively coupled plasma optical emission spectroscopy (ICP-OES) technique, and the results are presented in

Table 1. The measured chemical compositions of specimens used in this study (in wt.%).

	Elemental Composition (wt.%)
Material	Fe	C	Cr	Cu	Mg	Mn	Ni	S	Si
44W Steel	99.20	0.16	0.05	0.26	6 × 10−4	0.80	0.15	0.005	0.23
Monel 400	3.35	0.17	0.07	32.28	0.03	1.05	66.43	0.002	0.33
304L SS	69.26	0.03	17.79	0.42	5 × 10−4	1.74	7.90	0.019	0.35
Arc-DED 308L SS	64.88	0.03	19.39	0.14	4 × 10−4	1.88	10.12	0.023	0.54

.

Among them, 44W steel (Figure 1b, Table 1), though widely used due to its excellent balance of strength, ductility, and cost-effectiveness, is markedly susceptible to HE, especially under conditions that promote hydrogen ingress. As previously discussed, its BCC structure provides larger interstitial sites between iron atoms, allowing greater hydrogen solubility and diffusivity. Previous studies have shown that this increased hydrogen permeability, coupled with the presence of microstructural features like dislocations, grain boundaries, and inclusions, significantly enhances the susceptibility of low-carbon steels like 44W to hydrogen-induced cracking [52].

Monel 400 (Figure 1b), an alloy composed primarily of nickel and copper (Table 1), often containing minor additions of other elements like iron and manganese, is known for its exceptional resistance to corrosion and HE. The high nickel content in Monel significantly reduces hydrogen permeation, as nickel has a low affinity for hydrogen [53,54]. Copper also contributes to lowering hydrogen permeation by acting as a barrier to hydrogen diffusion. The FCC structure of the major constituents of Monel generally results in lower hydrogen diffusivity. As a result, Monel 400 exhibits lower hydrogen solubility and diffusivity than BCC-structured steels, making it substantially less permeable to hydrogen. Studies have consistently found that Monel exhibits good resistance to HE, even under harsh environmental conditions [55,56]. This resistance is largely attributed to its high nickel content. The results obtained from ICP analysis for Monel align closely with standard composition values reported in the literature [57].

It is worth noting that the wrought equivalent of 308L stainless steel feedstock wire, when considering mechanical properties and corrosion behavior, is typically 304L stainless steel (UNS S30403), which is commonly welded using 308L filler despite minor compositional differences. Therefore, when comparing a wrought base metal to arc-DED fabricated 308L, 304L stainless steel serves as the most practical and widely available reference material.

Arc-DED, also known as the wire arc additive manufacturing (WAAM) process, using the Cold Metal Transfer (CMT) mode, a highly controlled process known for its low-heat input and accurate material deposition, was employed to fabricate the 308L austenitic stainless-steel wall shown in Figure 1d. The CMT method is particularly advantageous for minimizing distortion and residual stresses, making it ideal for fabricating intricate geometries. The feedstock material was a 0.9 mm diameter 308L stainless steel wire fed at a speed of 4 m/min, ensuring consistent material flow during the deposition process. The arc was maintained using a steady current of 78 A and a voltage of 11.2 V, providing optimal energy for melting the filler material while preventing extreme heat that could compromise the material’s microstructure. Deposition was carried out using a zigzag printing strategy to enhance layer uniformity and structural integrity. To protect the molten pool from atmospheric contamination, pure argon was used as shielding gas at a flow rate of 15 L/min, maintaining an inert environment and preventing oxidation during deposition. The fabricated wall produced through this process is shown in Figure 1d, from which test specimens were sectioned and subsequently machined in the indicated direction. This specific orientation was strategically chosen to cut across the layering or deposition direction, rather than parallel to the deposited layers. This approach ensured that the machined specimens incorporated the full cross-section of the deposited layers, thereby providing a more comprehensive and representative evaluation of the material’s structure and properties.

Microstructure

The samples for microstructural analysis were prepared by sectioning, mounting, grinding, and polishing in accordance with standard practices [58] to achieve a smooth, mirror-like finish. This was followed by etching with appropriate etchants, as specified by ASTM standards for each material, to reveal microstructural features. For 44W steel, a Nital solution was applied [59], while Monel 400 was etched using a 1:1 mixture of HNO₃ and H₂O solution [60]. For the 304L and arc-DED 308L austenitic stainless steels, etching was conducted using an acid solution composed of hydrochloric acid (HCl) and nitric acid (HNO₃) in a 3:1 ratio [61,62,63]. The 304L austenitic stainless-steel sample was immersed in the solution for 80 s, while the arc-DED 308L sample required only 40 s of immersion due to differences in surface characteristics and microstructural response. The etched surfaces were then rinsed with distilled water, cleaned with ethanol, and dried to prevent contamination. The samples were then analyzed under a confocal laser scanning microscope (CLSM) (Keyence VK-X1000, Chicago, IL, USA) to allow comparison of microstructures, including phase distribution, and surface features.

X-ray diffraction (XRD) analysis was conducted to investigate the crystal structure and phase identification of the four materials under study. The evaluation was performed using the advanced Bruker D8 Advance XRD system (Bruker, Billerica, MA, USA), equipped with a high-speed LynxEye™ detector and a copper X-ray tube. The system was operated at 40 kV and 40 mA, utilizing Cu Kα radiation of 1.5418Å wavelength. The samples were scanned over a 2θ range of 20° to 100° with a step size of 0.049°. The diffraction patterns obtained were analyzed using Bruker’s EVA software (DIFFRAC.EVA V7, DIFFRAC.Part 11 V8) and compared against the International Centre for Diffraction Data (ICDD) Powder Diffraction File (PDF) database for phase identification and pattern matching.

The hardness of the samples was measured using the Rockwell hardness testing method in accordance with standards [64,65], specifically employing the B scale. Prior to testing, the sample surfaces were carefully prepared to ensure they were smooth and free from contaminants that could affect the results. Each specimen was securely mounted on the platform of the Rockwell hardness tester (Instron Wilson Hardness^®, Rockwell 2000, Rolling Meadows, IL, USA). A 1/16-inch steel ball indenter was brought into contact with the surface, and a preliminary (minor) load of 10 kgf was applied to establish an initial indentation and ensure proper seating of the indenter. Once the minor load was stabilized, a major load of 100 kgf was applied, causing deeper penetration of the indenter into the material. The equipment automatically measured the depth of this indentation after maintaining the major load for a specified dwell time, thereby determining the Rockwell B hardness value.

To ensure accuracy and repeatability, hardness measurements were taken at five different points across the surface of each specimen. The indentation points were spaced sufficiently far apart to minimize any potential influence of localized work hardening from previous indentations on subsequent measurements. The average of these five values was calculated and reported as the final hardness of each specimen.

3. Experimental Procedures

For the purpose of this study, two primary experiments were conducted under varying hydrogenating conditions: electrochemical hydrogen charging and uniaxial tensile testing. Each experiment was performed a minimum of four times, and the resulting data were averaged. The variation in the measured values was within a range of 5–8%. Detailed descriptions of these procedures are provided in the subsequent sections.

3.1. Hydrogen Charging

Electrochemical hydrogen charging was conducted on four different materials, 44W steel, Monel 400, 304L austenitic stainless steel, and arc-DED 308L austenitic stainless steel, to explore how hydrogen permeation influences their respective tensile properties. In this process, each specimen served as the working electrode in a two-electrode setup (Figure 2a,b), while a platinum counter electrode completed the circuit. A potentiostat/galvanostat (Figure 2c) system (Pine Research Instrumentation, Durham, NC, USA), capable of delivering a constant current of up to 1 A, was used to regulate the charging process. Hydrogen charging for each sample was maintained for three hours, a duration aligned with the time required to reach steady-state hydrogen absorption, as determined in previous studies [27,50].

The electrolyte solution used for the hydrogen charging process was an alkaline 0.1 M NaOH solution with an average pH of 12.5 and a controlled temperature of 23.3 °C. To inhibit the recombination of absorbed hydrogen atoms into hydrogen molecules (H₂), ammonium thiocyanate (NH₄SCN) was introduced as a recombination poison. To ensure uniformity of the electrolyte and minimize oxidation, the solution was continuously stirred using a magnetic stirrer and purged with argon gas to remove dissolved oxygen.

Prior to hydrogen charging, the gauge regions of each specimen were meticulously cleaned with acetone to remove any residual oil and surface debris possibly arising from processing and transportation. To restrict hydrogen exposure exclusively to the gauge section, chemical-resistant tape was used to mask the other areas of the specimens. The charging system was controlled via Aftermath software (Durham, NC, USA, 2.1.13189), which enabled precise management and monitoring of charging parameters such as current, charging time, and data logging intervals. For consistency, all samples were positioned in the same orientation within the charging cell.

The hydrogen concentration absorbed by each specimen was indirectly determined through the charging current, leveraging the principles established by Devanathan and Stachurski’s permeation method [66], in compliance with ISO 17081 standards [67], as detailed in the authors’ preceding work [27]. In this method, the specimen was mechanically polished to remove surface irregularities and subsequently cleaned to eliminate contaminants before being mounted in a Devanathan–Stachurski-type (double electrolyte) hydrogen permeation cell. The specimen, which served as the working electrode, was clamped between the charging and oxidation cells connected to the power supply unit (PSU) and electrochemical workstation, respectively. Both cells were filled with 0.1 M NaOH, pH ≈ 12.5, to provide a conductive alkaline medium, while NH₄SCN was added to the charging cell as a recombination poison to suppress the formation of H₂ on the entry surface.

Prior to initiating hydrogen charging, the electrolyte was first introduced into the oxidation cell, followed by purging with Ar gas to displace any dissolved oxygen. The charging cell remained idle while the system was monitored until the oxidation current dropped below 0.1 μA/cm², indicating the background current [68]. This value was later subtracted during data analysis to isolate the true hydrogen permeation signal and eliminate the effect of any residual hydrogen from prior processing.

Hydrogen was generated in the charging cell by applying a constant current from the PSU, with the positive terminal connected to a platinum counter electrode (anode) and the negative terminal to the specimen (cathode). The applied current density governed the electrochemical reduction of H₂O molecules, producing atomic hydrogen. These atoms diffused through the specimen and were subsequently oxidized at the exit surface in the oxidation cell, where the resulting anodic current was continuously monitored by the electrochemical workstation. The permeation current served as a direct quantitative measure of hydrogen transport through the metal. All experiments were conducted at room temperature.

This methodology provided critical data (

Table 2. Obtained data from hydrogen permeation tests.

	Current Density (mA/cm2)	CH (mol/m3)	Deff (m2/s)	tss (min)
Material	10	3.14	3.38 × 10−11	120

), including the effective diffusivity (D_eff), hydrogen concentration within interstitial lattice and trap sites (C_H), and the time to reach steady-state hydrogen absorption (t_ss). Using Sieverts’ law (Equation (1)) [69,70,71], a constant (K) was calculated, correlating the charging current density (i), a function of specimens’ geometry, as summarized in

Table 3. Selected hydrogenating conditions.

Hydrogen Concentration (wppm)	Steady-State Charging Time, tss (hours)	Charging Current Density (mA/cm2)	Charging Current (mA)
0.00	3	0.00	0.00
0.20		2.50	20.99
0.50		15.62	131.22
1.00		62.48	524.86
1.20		89.98	755.80
1.30		105.60	887.02

, to the hydrogen concentration (C_H) within the material. This constant enabled the quantification of hydrogen concentrations under different charging conditions, facilitating a comprehensive comparison of how hydrogen permeation affects the mechanical properties of each material.

$${\mathrm{C}}_{\mathrm{H}} =\mathrm{K}\sqrt{\mathrm{i}}$$

(1)

where C_H signifies the hydrogen concentration, K is a constant, and i is the current density.

The selection of the above hydrogenating conditions was guided by the geometric characteristics of the specimens, specifically, the 8.4 cm² surface area of the gauge section. This dimension was critical in determining the current density required for effective hydrogen charging as well as ensuring that the applied current remained within the potentiostat’s 1A operational range. Also, charging time was increased to 1.5 times that employed in the permeation tests to ensure full hydrogen saturation. These parameters helped maintain controlled and uniform hydrogen charging across the specimens’ surfaces, optimizing conditions for subsequent analysis.

It is recognized that the electrochemical hydrogen charging technique employed in this study differs fundamentally from hydrogen exposure under high-pressure gaseous environments. In particular, the dissociation and absorption mechanisms vary significantly: while gaseous hydrogen requires catalytic surface dissociation of molecular H₂ into atomic hydrogen prior to absorption, electrochemical methods generate atomic hydrogen in situ at the metal–electrolyte interface during cathodic polarization [72,73]. These mechanistic differences inherently affect both the kinetics and extent of hydrogen ingress.

Despite this distinction, electrochemical charging was selected due to laboratory safety considerations and its widespread acceptance as a controlled and reproducible method for introducing hydrogen into metallic systems. Furthermore, the present investigation builds directly upon our previous work by utilizing the same charging technique, thereby ensuring methodological consistency and enabling meaningful comparison across studies. It is, however, acknowledged that the embrittlement behavior reported herein is specific to the electrochemical charging route and may not fully replicate that induced by gaseous hydrogen exposure. All interpretations have, therefore, been made within the framework of the electrochemical hydrogen charging methodology.

Also, it is important to acknowledge that the hydrogen charging in this study was conducted ex situ, prior to mechanical testing. While this method effectively simulates residual hydrogen uptake from prior exposure or processing, it does not fully capture the dynamic interactions that may occur during simultaneous hydrogen ingress and mechanical loading. This is particularly relevant for passive materials such as austenitic stainless steels, where the integrity of the passive film plays a significant role in controlling hydrogen entry and where mechanical deformation can locally disrupt this film [74,75,76], thereby enhancing hydrogen absorption. As such, the embrittlement behavior observed in this study may differ in magnitude or character from that which would occur under in situ conditions.

However, the methodology in this study enabled a controlled assessment of HE susceptibility as a function of pre-charging levels, along with the associated microstructural and fractographic characteristics. By deliberately excluding mechanical loading during the hydrogen charging phase, the setup effectively isolated the effects of hydrogen exposure alone. This simplification enhances the clarity of interpretation by removing the confounding variable of stress-assisted hydrogen ingress. In contrast, in situ charging approaches, where mechanical loading is applied simultaneously with hydrogen exposure, can accelerate hydrogen uptake due to film rupture or enhanced diffusion under stress but make it difficult to quantify the actual hydrogen content during deformation.

It is noteworthy that all instruments used in this study were properly calibrated to ensure high measurement accuracy. For the electrochemical hydrogen charging, central to both previous and current investigations, the simulated alkaline environments were rigorously maintained. pH levels were consistently monitored, temperature was controlled throughout the charging duration, and the solution was uniformly stirred at speeds ranging between 350 rpm and 400 rpm. To minimize the risk of hydrogen ion depletion, a fresh solution was prepared after each specimen was charged. These measures collectively reinforce the reliability and reproducibility of the experimental data obtained.

3.2. Tensile Testing

Each of the uncharged and hydrogenated specimens underwent uniaxial tensile testing (Figure 3) to assess the mechanical properties before and after hydrogen charging, aiming to identify how hydrogen diffusion influenced tensile behavior. Specimens had surface and cross-sectional areas of 840 mm² and 40 mm², respectively. To ensure a comprehensive analysis of the materials’ behavior under varying hydrogen concentrations, six samples were prepared for each hydrogen concentration level: 0.00, 0.20, 0.50, 1.00, 1.20, and 1.30 weight parts per million (wppm). These specific conditions were carefully selected to ensure a well-distributed range of hydrogen concentrations, while considering the potentiostat’s capacity to deliver a stable and constant charging current, allowing identification of clear trends across hydrogen concentrations.

Uniaxial tensile testing was conducted immediately following the completion of electrochemical hydrogen charging. The average time interval between the removal of specimens from the charging cell and the initiation of tensile testing was approximately 2–3 min, which accounted for specimen handling and setup within the tensile testing grips. This brief time was consistent across all specimens and all hydrogen charging levels, ensuring uniformity in hydrogen exposure conditions prior to mechanical testing.

Moreover, preceding studies have demonstrated that diffusible hydrogen does not exit the metal matrix instantaneously after charging ceases [27,77,78]. In fact, several studies report that, even after exposure to air or inert environments, significant quantities of hydrogen, particularly when trapped at microstructural defects or interfaces, can remain within the specimen for extended periods, depending on the alloy type, charging conditions, and microstructural trapping capacity [79,80,81]. Therefore, the short handling time used in this study is not expected to have led to significant hydrogen loss or variability in embrittlement behavior across specimens.

Prior to testing, the specimens were carefully mounted in the upper and lower grips of the Instron 1332 tensile testing machine (Figure 3a,b), ensuring proper alignment using a straight edge to maintain perpendicularity to the grips. A well-calibrated extensometer (Figure 3a) with a 25 mm gauge length was attached to the gauge section of each specimen to accurately monitor and record displacement during testing.

The tensile tests were conducted at a constant crosshead speed of 5 mm/min, corresponding to a strain rate of 0.2 min⁻¹. The specimens were subjected to this strain rate until failure, ensuring uniform testing conditions for comparative analysis. Analysis of the collected data described the variations in tensile strength, elongation, and fracture behavior due to the presence of hydrogen, facilitating an in-depth understanding of HE mechanisms under different hydrogen charging levels.

4. Results and Discussion

It should be noted that, within the context of this study, the term “hydrogen diffusion” is used to describe the overall mobility of hydrogen atoms within the metallic microstructure. This includes both lattice diffusion through interstitial sites and the transient interactions of hydrogen with reversible trapping sites such as dislocations, grain boundaries, and inclusions. As such, the use of the term does not imply purely free, untrapped lattice diffusion but rather encompasses the effective transport behavior of hydrogen in the presence of microstructural traps.

4.1. Metallography

The microstructural analysis of etched 44W steel (Figure 4a) revealed a typical dual-phase structure characterized by ferrite and pearlite phases, with elongated grains indicative of a prior cold working process. Under a confocal microscope, the ferrite-rich regions appeared brighter, while the pearlite phase exhibited a darker contrast, aiding clear phase differentiation. The darker contrast of the pearlite phase is attributed to its alternating lamellae of ferrite and cementite (Fe₃C), which contributes to increased strength and hardness. In contrast, the ferrite phase, comprising nearly pure α-iron with a BCC crystal structure, appeared as smooth, bright regions and is associated with ductility and toughness.

Monel 400 (Figure 4b), a single-phase solid solution alloy, exhibited a uniform and homogenous matrix characterized by well-defined grain boundaries. The absence of secondary phases or precipitates is consistent with its metallurgical nature as a solid solution, where the copper atoms are substitutionally dissolved in the nickel matrix. This uniform microstructure contributes to Monel 400’s notable corrosion resistance and mechanical stability and provides a consistent baseline for evaluating hydrogen interaction mechanisms in comparison to more structurally complex alloys.

The 304L austenitic stainless steel (Figure 4c) revealed microstructural features including twin boundaries, equiaxed austenitic grains and ferrite stringers, and elongated ferrite regions dispersed within the austenitic matrix. The detection of the ferrite stringers suggests localized microstructural heterogeneities, likely originating from variations in chemical composition or thermal processing during fabrication. Microstructural analysis of the arc-DED fabricated austenitic stainless steel (Figure 4d) revealed a combination of lathy delta ferrite, skeletal delta ferrite, and columnar dendrites, features that reflect the interplay of rapid solidification dynamics, localized compositional variations, and steep thermal gradients inherent to the arc-DED process. Lathy delta ferrite (Figure 4d) refers to elongated, plate-like ferrite structures formed within the austenitic matrix due to rapid solidification and steep thermal gradients in AM processes. These features arise when ferrite-stabilizing elements like chromium promote delta ferrite formation during directional solidification. Skeletal delta ferrite appears as a branching network retained in the austenitic matrix, resulting from non-equilibrium conditions and localized compositional variations that prevent full transformation to austenite. Both morphologies reflect the rapid cooling and complex thermal profiles inherent in arc-DED fabrication [82,83].

Columnar dendrites (Figure 4d) are elongated, tree-like structures that form within the austenitic matrix, often aligned with the direction of heat dissipation during solidification. These structures are a quality of directional solidification processes driven by high thermal gradients and anisotropic heat flow during AM. As the molten pool cools, dendritic austenite grows preferentially along the thermal gradient, forming columnar morphology. The presence of these dendrites is closely tied to the rapid solidification rates and the unique thermal conditions of 3D printing. The columnar dendritic structure can also be influenced by the local chemical composition, with alloying elements like nickel stabilizing the austenitic phase during solidification [84,85].

4.2. X-Ray Diffraction Analysis and Hardness

Analysis showed that the diffraction pattern for 44W steel (Figure 5a) revealed a predominant BCC α-ferrite phase, accompanied by a significant presence of Fe₃C (cementite), consistent with the expected microstructure of low-carbon steel. Monel 400 (Figure 5b) exhibited an FCC crystal structure, as anticipated for this single-phase solid solution alloy. Both 304L and arc-DED 308L austenitic stainless steels (Figure 5c,d) displayed nearly identical diffraction patterns, confirming their FCC crystal structures and reinforcing the findings from the ICP analysis, which indicated comparable elemental compositions and material characteristics. These results collectively validate the structural similarities between conventionally manufactured 304L and AM arc-DED 308L austenitic stainless steels.

The 44W steel exhibited the highest hardness at 96.43 HRB (

Table 4. Hardness values of specimens.

Material	Average Rockwell Hardness (HRB)
44W Steel	96.43
Monel 400	82.41
304L Austenitic Stainless Steel	85.07
Arc-DED 308L Austenitic Stainless Steel	78.84

). 44W steel is a carbon steel with a predominantly ferritic–pearlitic microstructure, where the presence of carbon forms iron carbide (cementite) within the pearlite. This pearlite contributes to higher hardness due to its mixture of hard cementite and softer ferrite, providing a balanced strength-to-ductility ratio. Additionally, the cold-finished condition of the 44W steel promotes grain refinement and work hardening, further enhancing its overall hardness. Monel exhibited a hardness value of 82.41 HRB, which can be attributed to its solid solution strengthening mechanism. This atomic-scale distortion impedes dislocation movement, thereby enhancing hardness. Additionally, Monel’s FCC structure provides good resistance to deformation, while its relatively high nickel content contributes to its overall strength and corrosion resistance, further supporting the observed hardness value.

It was observed that 304L austenitic stainless steel exhibited a higher average hardness value of 85.07 HRB, while the arc-DED fabricated 308L stainless steel demonstrated a lower value of 78.84 HRB. This difference can be attributed to the combined effects of processing history and phase composition. The 304L sample, being cold-worked, had undergone significant plastic deformation, leading to an increased dislocation density and strain hardening, both of which contribute to enhanced hardness. In contrast, the 308L sample, produced via wire WAAM, was in an as-cast condition. The solidification dynamics and thermal cycling inherent in the WAAM process tend to produce a coarser microstructure with lower dislocation density, thereby reducing hardness relative to cold-worked material.

Additionally, the 308L alloy contained a higher volume fraction of delta ferrite (δ-ferrite), a BCC phase that is typically softer than the surrounding FCC austenitic matrix. The presence of δ-ferrite, which forms as a result of the thermal conditions during additive manufacturing, contributes to the reduction in overall hardness by introducing regions that deform more readily under load. Together, these microstructural and processing differences explain the lower hardness observed in the 308L sample compared to the 304L stainless steel. The influence of material hardness on hydrogen permeation is largely governed by its correlation with underlying microstructural features. While hardness itself is not a direct determinant of hydrogen transport, variations in hardness may reflect changes in grain size, phase distribution, or surface treatment, all of which can influence hydrogen uptake and mobility [86,87].

4.3. Stress–Strain Behavior

Comparative analysis revealed pointed differences in how each material responded to hydrogen charging under tensile loading. 44W steel (Figure 6a) underwent the most substantial reductions in both ductility and toughness, with %elongation dropping from 6.9% to 4.8% and toughness decreasing from 51.11 MJ/m³ to 36.96 MJ/m³, reflecting the highest vulnerability to HE among the studied materials. Monel 400 (Figure 6b) and 304L austenitic stainless steel (Figure 6c) demonstrated exceptional resilience, with Monel 400 showing minimal changes and 304L maintaining these properties. For Monel, %elongation decreased slightly from 29.3% to 28.69%, while the toughness declined marginally from 171.10 MJ/m³ to 165.69 MJ/m³.

The behavior of 304L austenitic stainless steel (Figure 6c) under hydrogen charging was particularly noteworthy, as its elongation and toughness remained virtually unchanged, with the former changing slightly from 28.5% to 28.9% and the latter from 202.51 MJ/m³ to 203.56 MJ/m³. This stability suggests that hydrogen may play a role in stabilizing the FCC lattice or relieving internal stresses, thereby preserving both ductility and toughness. These findings indicate that 304L maintains its mechanical integrity across all levels of hydrogenating conditions, showing performance consistent with its uncharged counterparts. On the other hand, the arc-DED 308L austenitic stainless steel (Figure 6d) showed a moderate reduction in performance upon hydrogen charging. Its elongation decreased from 33.6% to 27.4%, and its toughness dropped from 156.74 MJ/m³ to 149.92 MJ/m³. These reductions, though less severe than those observed in 44W steel, are more pronounced than in conventional 304L austenitic stainless steel. The increased susceptibility of the arc-DED fabricated material to the HE is likely due to microstructural characteristics inherent to AM, such as residual stresses, certain defects, and structural discontinuities, which can act as hydrogen trapping sites and intensify embrittlement.

These findings suggest a practical conception for material selection and design for hydrogen infrastructure. The observed unique tensile performances of materials under hydrogen charging highlight the critical importance of selecting hydrogen-compatible alloys for structural components in natural gas pipelines, storage tanks, and pressure vessels within hydrogen-rich environments.

The significant degradation in ductility and toughness observed in 44W steel indicates that it is poorly suited for hydrogen transport or storage infrastructure without additional mitigation measures. Its high susceptibility to HE suggests a greater risk of premature failure under service conditions involving fluctuating temperatures, pressures, or mechanical loading, conditions typical of pipeline systems and fittings.

In contrast, the considerable retention of mechanical integrity in Monel 400 and 304L stainless steel after hydrogen exposure is suggestive that these materials offer promising alternatives for critical components in hydrogen infrastructure due to their reduced hydrogen diffusivity and inherent microstructural resistance to cracking. This is particularly valuable in areas where localized hydrogen accumulation and stress concentrations are expected, such as weld zones or bends in piping systems.

The behavior of arc-DED 308L stainless steel highlights the emerging role of AM in customizing materials for specific hydrogen service applications. Its discussed unique microstructural characteristics may influence hydrogen trapping and mobility in ways that differ from wrought alloys. Understanding and tailoring these features can guide the development of next-generation components optimized for hydrogen exposure, particularly in high-pressure or cyclic loading environments. The results of this study can inform standards development, material qualification, and design strategies as industries transition toward cleaner energy systems involving hydrogen.

4.4. %Elongation and Toughness vs. Hydrogen Content

Toughness was calculated as the area under the engineering stress–strain curve up to the point of fracture. To compute this area, a custom function based on the trapezoidal numerical integration method was implemented. The %elongation and toughness of the four materials under hydrogenating conditions exhibited varying behaviors: 44W steel experienced significant reductions, Monel 400 showed a relatively smaller reduction, 304L austenitic stainless steel demonstrated the highest resilience to hydrogen diffusion, and arc-DED 308L austenitic stainless steel displayed an initial reduction followed by a subsequent improvement in these properties.

The observed significant reduction in the %elongation and toughness of 44W steel (Figure 7a) with increasing hydrogen content is ascribed to the HEDE mechanism, which weakens interatomic bonds at critical microstructural sites such as grain boundaries, inclusions, and dislocation pileups [28,29]. This weakening, particularly in high-stress regions like crack tips, promotes brittle fracture and reduces the material’s ductility and toughness. 44W steel’s ferritic–pearlitic microstructure, coupled with its BCC crystal structure, intensifies its susceptibility to the HEDE mechanism. Grain boundaries and ferrite–pearlite interfaces serve as hydrogen trapping sites where hydrogen accumulates and weakens atomic cohesion, making these regions vulnerable to crack initiation and propagation. Additionally, the alternating ferrite and cementite layers in pearlite may provide numerous interfaces for hydrogen localization, further intensifying decohesion.

The BCC crystal structure, with its lower APF, facilitates easier hydrogen diffusion; thus, an infiltrated hydrogen atom can cause damage at various areas within the material’s microstructure, mechanistically termed as reversible hydrogen traps. This increased diffusivity enables hydrogen to accumulate at stress-concentrated areas, amplifying its detrimental effects on cohesion. Consequently, the ease of hydrogen diffusion, combined with abundant trapping sites, leads to a transition from ductile to brittle failure, resulting in significant reductions in the elongation and toughness of 44W steel under hydrogenated conditions.

The relatively minimal impact of hydrogen diffusion on the %elongation and toughness of Monel 400 (Figure 7b), as compared to 44W steel, can be attributed to the interplay between the HEDE mechanism and Monel 400’s distinct microstructural characteristics. Unlike 44W steel, Monel 400 possesses an FCC crystal structure, which significantly influences its interaction with hydrogen. In the FCC structure of Monel 400, hydrogen trapping sites are predominantly irreversible. This means that hydrogen atoms, once trapped, are unable to diffuse further or cause widespread embrittlement across the material [88,89]. Accordingly, the embrittling effects of dissolved hydrogen are largely localized to isolated microstructural features, such as grain boundaries or discrete defects, without significantly compromising the overall mechanical integrity of the material. This limited hydrogen mobility contrasts sharply with the behavior observed in BCC materials like 44W steel, where hydrogen diffuses readily and accumulates extensively at stress-concentrated regions, augmenting the HEDE mechanism and leading to significant embrittlement.

Furthermore, the higher APF of Monel 400’s FCC structure further restricts hydrogen diffusivity, reducing the extent of hydrogen-induced bond weakening. This chemistry between the HEDE mechanism and the microstructural features of Monel 400 effectively mitigates the embrittling damages caused by hydrogen, as evidenced by the relatively minor reductions in %elongation and toughness observed under hydrogenated conditions.

For 304L austenitic stainless steel (Figure 7c), the negligible impact on %elongation and toughness focuses on its notable resistance to HE. This resilience is attributed to its FCC crystal structure and notably uniform microstructure, which inherently offers high ductility and effectively mitigates hydrogen-induced degradation. These attributes not only minimize hydrogen diffusion but also inhibit its localization, thereby reducing the susceptibility to embrittlement and preserving the material’s mechanical integrity. This behavior aligns with the reported outcomes in sour service conditions [90], which confirm the alloy’s stability in hydrogen-rich environments.

In contrast, the arc-DED 308L austenitic stainless steel (Figure 7d) exhibited a more complex response. The observed initial reduction in %elongation and toughness, once again, can be attributed to the HEDE mechanism. This mechanism dominated at relatively lower hydrogen concentrations due to the material’s higher defect density, including porosities, layered interfaces, and residual stresses introduced during the AM process. These features create favorable sites for hydrogen-induced embrittlement, leading to relatively less ductile fracture and reduced ductility in comparison to Monel 400 and 304L SS.

However, as the hydrogen concentration increased further, the arc-DED 308L austenitic stainless steel exhibited a noticeable rise in %elongation and toughness, a behavior driven by the activation of the HELP mechanism. Under HELP, hydrogen lowers the energy barriers for dislocation motion by weakening atomic bonds near dislocations [91,92], reducing lattice friction between atoms, and diminishing dislocation interaction stresses. This facilitates more uniform and localized plastic deformation, which can temporarily offset embrittlement effects. This results in a partial recovery of ductility and toughness at higher hydrogen levels, reflecting the dynamic interplay between competing HE mechanisms in the arc-DED microstructure.

Additionally, the unique microstructure of the arc-DED 308L austenitic stainless steel plays a pivotal role in this transition. Its columnar grain structures create numerous sites for dislocation activity, rendering it more responsive to HELP-driven plasticity at elevated hydrogen concentrations. Despite its initial susceptibility to embrittlement due to its heterogeneous microstructure, the material benefits significantly from the enhanced dislocation motion induced by hydrogen. This increased plasticity delays crack initiation, thereby improving the material’s mechanical performance under high hydrogen levels.

4.5. Yield Strength and Ultimate Tensile Strength vs. Hydrogen Content

No significant changes were observed in the average yield strength (YS) and ultimate tensile strength (UTS) of cathodically hydrogen-charged 44W steel, Monel 400, and 304L austenitic stainless steel (Figure 8a,b) at all levels of hydrogenating conditions. This observation suggests that hydrogen did not significantly affect their elastic or uniform plastic deformation behavior under the applied stress conditions. This behavior is often characteristic of materials where HE mechanisms primarily affect post-yield phenomena such as ductility loss, premature fracture, or toughness degradation rather than the initial plastic flow. These findings are consistent with those reported in previous studies by the authors [50,51].

The observed gradual increase in the YS and UTS of hydrogen-charged arc-DED 308L austenitic stainless steel (Figure 8a,b) can be attributed to the interaction of hydrogen with the material’s unique microstructural features likely due to the combined effects of hydrogen trapping, solid solution strengthening, and the distinct microstructural characteristics of arc-DED austenitic stainless steel. These factors work together to improve the material’s tensile properties when exposed to hydrogen.

Hydrogen trapping at defects and grain boundaries is a contributing factor. As discussed above, 3D-printed materials often contain microstructural imperfections such as porosities, residual stresses, and non-equilibrium grain structures [93,94]. These features act as effective hydrogen traps, increasing localized strain hardening and impeding dislocation motion. The interaction between hydrogen and these defects enhances YS. Moreover, the columnar grains and substructures inherent to 3D-printed materials further contribute to this strengthening effect. Also, diffused hydrogen atoms can induce a degree of solid solution strengthening. As an interstitial atom, hydrogen interacts with the material’s lattice, causing local distortion and making dislocation motion more difficult. While typically subtle, this strengthening effect is amplified in 3D-printed materials due to the rapid cooling and residual stresses introduced during the AM process. This phenomenon aligns with findings in studies such as reference [95].

Finally, the unique microstructure of arc-DED 308L austenitic stainless steel, characterized by features like fine cellular substructures, residual stresses, and anisotropic grain orientation, synergizes with hydrogen’s effects to enhance tensile properties. These heterogeneous microstructural features, combined with hydrogen-induced interactions, may increase resistance to plastic deformation, resulting in higher YS and UTS. This behavior aligns with findings reported in [96], which highlight how AM-induced microstructural characteristics govern the mechanical response of materials under hydrogen charging.

4.6. Modulus of Elasticity vs. Hydrogen Content

The results of this study revealed that the average elastic moduli of hydrogenated 44W steel, Monel 400, and 304L austenitic stainless steel remained consistent with those of their non-hydrogenated counterparts, as demonstrated by the near-linear trend in the data (Figure 9). This consistency indicates that the intrinsic stiffness of these materials and their ability to undergo elastic deformation under stress are largely unaffected by varying levels of hydrogen absorption. The negligible respective variation in elastic moduli across different C_H aligns with the fundamental understanding that elasticity, as an inherent material property governed by atomic bonding forces, is relatively insensitive to external factors such as hydrogen uptake. This behavior stresses the materials’ resilience to hydrogen-induced effects on their elastic behavior, accounting for the stability observed in their elastic moduli.

In the arc-DED 308L austenitic stainless steel, hydrogen’s interaction with the unique microstructural features tends to enhance the stiffness of the material, as observed in Figure 9. For instance, hydrogen trapping at dislocations and grain boundaries can impede localized deformation, effectively increasing the material’s resistance to elastic strain and resulting in a higher elastic modulus.

Also, hydrogen, as an interstitial element, has the potential to alter the stiffness of metallic materials by interacting with their atomic bonds, a phenomenon termed as hydrogen interstitial strengthening [9]. In FCC metals like austenitic stainless steels, dissolved hydrogen atoms can increase repulsive forces between atoms and reduce lattice compliance. This hydrogen-induced stiffening effect has been shown to produce a measurable increase in the elastic modulus, likely due to the constrained movement of atoms in the distorted lattice and the hydrogen’s role in stabilizing certain microstructural features [97,98,99]. Consequently, the presence of interstitial hydrogen can subtly but consistently raise the elastic stiffness of hydrogen-sensitive alloys.

While traditional strengthening mechanisms may primarily influence YS and hardness, certain atomic-scale interactions can subtly impact the elastic modulus. Solid solution strengthening, either substitutional or interstitial, can distort the host metal lattice and hinder dislocation motion. Although this mechanism typically has a negligible effect on the elastic modulus, interstitial solutes, such as hydrogen, can significantly alter the stiffness of atomic bonds within the lattice [100,101]. When these atoms occupy interstitial positions, they introduce localized lattice distortions and modify the interatomic force landscape, thereby slightly enhancing resistance to elastic deformation.

When hydrogen diffuses into the lattice of arc-DED 308L austenitic stainless steel, it can locally strengthen the atomic interactions, leading to an increase in stiffness, which is reflected in the elastic modulus. This phenomenon is more pronounced in materials with heterogeneous microstructures, such as 3D-printed alloys, where hydrogen diffusion can reinforce localized regions of the material.

5. Fractography Analysis

Fracture surfaces of uncharged specimens were analyzed and compared with those of their hydrogen-charged counterparts using a Thermo Scientific AXIA ChemiSEM scanning electron microscope (SEM) (Waltham, MA, USA) equipped with energy dispersive X-ray spectroscopy (EDX) capabilities. The findings have been discussed in the subsequent sections.

5.1. Uncharged Specimens

Analysis of the fracture surfaces of the uncharged 44W steel subjected to uniaxial tensile loading is presented in Figure 10. As shown in Figure 10a, the three distinct regions highlighted and color-coded showed various features. Region 1 (red) revealed the presence of inclusions along with evidence of ductile failure through dimpling. Region 2 (yellow) exhibited signs of directional cracking, indicative of localized stress concentration and crack propagation. Region 3 (blue) was characterized by stretched dimples and microvoids, features typically associated with plastic deformation and void coalescence mechanisms.

The dimples and microvoids (Figure 10b,d) are indicative of expected ductile fracture mechanisms, which typically occur when the material undergoes significant plastic deformation prior to fracture. This deformation initiates at stress concentration points, such as inclusions or second-phase particles within the microstructure, leading to microvoid nucleation. These microvoids grow and coalesce under tensile stress, forming the characteristic dimpled appearance on the fracture surface.

The directional cracks observed (Figure 10e) are a result of localized stress intensification within the material, where the primary cracks propagate along paths of least resistance. These paths are often aligned with areas of microstructural weakness, such as inclusions or interfaces between phases, which act as crack initiation and propagation sites.

The inclusions identified on the fracture surface (Figure 10b,c), through EDX analysis (

Table 5. Elemental composition of observed inclusions in uncharged 44W steel.

	Elemental Composition (wt.%)
Element	Al	Si	Cr	Mn	Fe	Co	Ni	Cu	Zn
Point 1	13.8	64.7	1.4	0.4	14.6	0.8	0.1	1.1	3.1
Point 2	0.1	0.3	0.4	1.1	97.4	0.0	0.7	0.0	0.0

), presented additional interpretations into their elemental composition and role in the fracture process. The high silicon content at point 1 (Figure 10c) suggests the presence of silicon-containing inclusions, such as silica (SiO₂) and silicates (e.g., MnSiO₃, Fe₂SiO₄). The former commonly forms during steelmaking as a result of oxidation reactions involving silicon when it reacts with oxygen through solidification or refining, and the latter forms when silicon reacts with other elements like manganese (Mn), iron (Fe), or calcium (Ca) during the steelmaking process [102,103,104]. Such inclusions are often brittle and act as stress concentrators, facilitating the nucleation of microvoids and crack initiation under tensile loading [50,51].

Point 2 (Figure 10c), characterized by high Fe content, is identified as a ferritic region within the matrix. This area represents an iron-rich phase with minimal alloying element segregation, likely where the matrix remains relatively pure. While such ferritic regions are generally less susceptible to fracture initiation, they may still influence crack propagation by interacting with surrounding inclusions.

The fracture surface of uncharged Monel 400 (Figure 11) subjected to uniaxial tensile testing was also analyzed. As shown in Figure 11a, region 1 (red) displayed coalesced microvoids and dimples, which are typical indicators of ductile fracture mechanisms. Region 2 (yellow) revealed the presence of microcracks and scattered microvoids, suggesting localized initiation sites for fracture propagation. Region 3 (blue) exhibited slip regions and fibrous dimples, reflecting plastic deformation and a fibrous fracture texture associated with the material’s toughness.

Shear lips (Figure 11a) form near the edges of the fracture surface due to shear stresses that dominate in these regions, resulting in localized plastic deformation. This is characteristic of a mixed-mode fracture where both tensile and shear forces contribute to failure. Dimples and fibrous dimples (Figure 11b,c), further indicative of ductile fracture, result from the nucleation, growth, and coalescence of microvoids, as explained above. However, as voids grow under tensile stress, they eventually merge, forming a fibrous appearance in the fracture surface, especially in regions of extensive plastic deformation.

Slip regions (Figure 11c) evolve from localized plastic deformation due to dislocation movement along specific crystallographic planes under the applied tensile stress. These regions appear relatively smoother and are associated with areas of intense deformation. Microcracks are initiated at stress concentrators such as inclusions, grain boundaries, or other defects within the material and propagate as the applied stress exceeds the material’s fracture toughness, and their growth is often influenced by the microstructural characteristics of Monel 400.

Coalesced microvoids (Figure 11b,d), as explained above, result from merging voids as the material encounters further plastic deformation. This coalescence marks the final stage of ductile fracture, where the material fails as a result of the complete loss of load-bearing capacity in the deformed region.

The fracture surface of the uncharged 304L austenitic stainless steel subjected to uniaxial tensile testing is displayed in Figure 12. As shown in Figure 12a, region 1 (red) revealed coalesced microvoids, micropores, elongated voids, dimples, and a visible crack direction, features indicative of ductile fracture accompanied by localized plastic deformation. Region 2 (yellow) showed a concentration of dimples and voids, further supporting the ductile nature of the fracture in that area. Region 3 (blue) identified the crack origin and the direction of its propagation, offering detailed comprehension of the initiation site and fracture path.

The presence of elongated dimples (Figure 12b) observed on the fracture surface of uncharged 304L austenitic stainless steel subjected to tensile loading, alongside other fracture characteristics (Figure 12c,d) already discussed, reflects the ductile nature of the material under the given loading conditions. Elongated dimples are indicative of crack initiation and microvoid coalescence, which are common features of ductile fracture, where the shape and orientation of the dimples correspond to the direction of plastic deformation experienced during tensile loading. Mechanistically, this phenomenon begins with the nucleation of microvoids at the aforementioned stress concentrators present in the material. As the tensile stress increases, these voids grow and elongate in response to the applied load, with the elongation direction aligning with the principal tensile axis. This morphology is a result of the high degree of plastic strain sustained by the material before the final fracture, indicative of significant energy absorption during deformation.

The fracture surface of the uncharged arc-DED 308L austenitic stainless steel (Figure 13) subjected to tensile loading reveals a complex relationship of deformation and failure mechanisms, as evidenced by the fibrous dimples, very large voids, and stream-like markings. These features collectively highlight the material’s ductile response under tensile stress, while also pointing to microstructural inhomogeneities and anisotropic properties fundamental to the AM process. The mechanisms of some observed features, i.e., shear lips and fibrous dimples (Figure 13a,b), are as discussed above.

The very large voids observed (Figure 13c,d) on the fracture surface likely originate from the inherent characteristics peculiar to the AM process. In AM-manufactured materials, incomplete fusion, trapped gas, or uneven solidification can create larger voids within the microstructure [105,106]. Under tensile loading, these voids act as stress concentrators, growing disproportionately as the material deforms plastically. Their size and prominence suggest they play a significant role in the fracture process, as they reduce the load-bearing cross-sectional area and promote early void coalescence.

The stream-like markings reflect the directional nature of plastic flow during tensile deformation. These markings likely result from the alignment of microstructural features and localized stress gradients, guiding the flow of material as voids grow and coalesce. This directional texture is accentuated in AM materials due to the layer-by-layer nature of the deposition process, which introduces anisotropy and preferential deformation paths. Together, these features paint a comprehensive picture of the interplay between the AM microstructure and the ductile failure mechanisms activated under tensile stress.

5.2. Charged Specimens

To effectively evaluate and delineate the potential effects of hydrogen permeation on material behavior, a comparative analysis was conducted between the fracture surface features of uncharged specimens and those subjected to the highest levels of hydrogen charging (1.30 wppm). This approach enabled a clear distinction between the baseline fracture characteristics and the alterations induced by maximal hydrogen ingress, providing critical insights into the embrittlement mechanisms and microstructural responses associated with hydrogen exposure. The fracture surface of 44W steel subjected to 1.30 wppm hydrogen content cathodic hydrogen charging (Figure 14) exhibited distinct features: cleavages, quasi-cleavages, microcracks, river marks, transgranular cracks, and sharp edge ridges that are significantly influenced by the presence of diffused and dissolved hydrogen. In contrast to the ductile features like dimples, microvoids, and shear lips observed in uncharged specimens, these brittle fracture characteristics are a direct consequence of HE, where the diffusible hydrogen weakens the steel’s microstructure, making it more susceptible to brittle failure.

The formation of sharp edge ridges (Figure 14a) is indicative of severe localized strain near crack tips, where hydrogen has embrittled the material, reducing its ability to undergo plastic deformation and leading to the formation of sharp, brittle fracture features. Also, river marks (Figure 14b), typically associated with ductile fracture, were also present but less pronounced in the hydrogen-charged specimens. In the presence of hydrogen, these river marks can become irregular or even suppressed due to the brittleness of the material caused by HE. This transition from ductile to brittle behavior is a direct manifestation of the detrimental effects of hydrogen diffusion, where the material’s ability to undergo plastic deformation is significantly reduced.

The formation of a directional crack mechanism in hydrogen-charged 44W steel (Figure 14c) is primarily driven by HE. As diffused hydrogen accumulates, it weakens the atomic bonds within the steel, particularly along grain boundaries and slip planes, reducing the cohesive strength between atoms. This weakening can facilitate crack initiation either intergranularly along grain boundaries or transgranularly along specific crystallographic planes typical of ferritic steels like 44W [107,108,109]. The fracture process is further influenced by the aforementioned dominant mechanism, HEDE, wherein hydrogen reduces the bond strength between atoms, particularly in high-stress areas like crack tips, thereby promoting brittle failure. Under applied tensile stress, cracks tend to initiate at these hydrogen-weakened regions and propagate in a relatively straight, directional manner, following paths of least resistance. These paths are often perpendicular to the loading direction or specific crystallographic orientations, giving rise to a directional crack morphology. This behavior contrasts with the more random, equiaxed, and dimpled appearance of ductile fractures and is a clear indication of the embrittling effects of hydrogen in 44W steel under tensile loading.

Furthermore, HE causes an increase in the susceptibility of the material to cleavage fracture (Figure 14b), where cracks propagate along specific crystallographic planes, typically in a brittle manner. The quasi-cleavages (Figure 14d) seen on the fracture surface result from the material experiencing brittle failure that is not entirely aligned with perfect cleavage planes, but still exhibit characteristics of brittle fracture, often influenced by localized hydrogen-induced weakening [110,111]. The formation of microcracks (Figure 14d) is another consequence of HE. As hydrogen accumulates at grain boundaries, it weakens and embrittles these regions, causing the initiation of microcracks, which propagate transgranularly through the steel, leading to the observed transgranular cracks (Figure 14b) on the fracture surface.

The fracture surface of Monel 400 charged with 1.30 wppm hydrogen and subjected to uniaxial tensile testing is shown in Figure 15. As illustrated in Figure 15a, region 1 (orange) exhibited numerous coalesced voids and a macrocrack, characteristically implying mixed fracture modes in this region but predominantly indicative of ductile fracture behavior. Region 2 (red) showed noticeable elongation of dimples, indicating significant plastic deformation prior to final fracture. Overall, the fracture surface features reflect a combination of predominantly ductile and slightly brittle fracture mechanisms. This behavior is attributed to the limited permeation of hydrogen into the material; while most regions retained the inherent ductility of Monel 400, localized interactions between the infiltrated hydrogen atoms and the microstructure contributed to isolated brittle responses during deformation and failure.

The observed sharp-edge cracks, macrocracks, and cleavage (Figure 15a,b,d), all characteristic features of brittle failure, are direct results of HE. Macrocracks and sharp-edge cracks primarily form as hydrogen diffuses into the alloy and accumulates in areas of high stress, as previously discussed, leading to the weakening of atomic bonds, followed by crack nucleation and growth under applied tensile stresses. Observed elongated dimples (Figure 15c) are characteristic of crack initiation and microvoid coalescence, hallmark features of ductile fracture. Their shape and alignment reflect the direction of plastic deformation sustained by the material during tensile loading. While hydrogen can reduce fracture toughness and facilitate cleavage-like fracture along crystallographic planes (Figure 15d), its limited ingress in this case means these brittle features had a negligible impact on the overall toughness and elongation of the charged Monel 400.

Furthermore, the formation mechanisms of the observed ductile features, elongated dimples, coalesced voids, shear lips, fibrous dimples, and slip regions (Figure 15a–d), have been thoroughly discussed in earlier sections of this article. However, the wrinkle formations (Figure 15a), appearing near regions of plastic deformation, such as shear lips or areas adjacent to microvoid coalescence, occurred due to localized plastic instability during deformation. As the material was subjected to tensile stress, slip bands, which are zones of concentrated dislocation movement, formed within the microstructure. These slip bands interact with stress concentrators, such as grain boundaries, inclusions, or other microstructural irregularities, causing uneven deformation and resulting in localized folding or wrinkling on the fracture surface [112,113].

Examination of the fracture surface of 1.30 wppm charged 304L austenitic stainless steel (Figure 16) showed that the material maintained its inherent ductile properties, indicating a negligible to no effect of hydrogen on its mechanical behavior, in agreement with the preceding findings on %elongation and toughness. The fracture surface exhibited macrocracks, microvoids, dimples, shear lips, necking, and tear ridges. The formation mechanisms of the macrocracks, microvoids, dimples, and shear lips (Figure 16a–d) have been discussed in the preceding sections.

Necking (Figure 16a) occurs as the material undergoes localized deformation under tensile loading, leading to a reduction in cross-sectional area. This phenomenon is a characteristic of ductile materials, and in this case, the material maintained its ductility despite the hydrogen exposure. Tear ridges (Figure 16d) are typically observed in ductile fractures formed as a result of localized plastic deformation during fracture, where the material undergoes tearing along specific crystallographic planes [114,115]. These ridges develop as cracks propagate through the material, particularly in regions experiencing significant necking. The formation of tear ridges contributes to the roughness and irregularity of the fracture surface, reflecting the material’s resistance to fracture and its ability to undergo substantial plastic deformation before failure.

The fracture surface of the charged arc-DED 308L austenitic stainless steel (Figure 17) displayed transgranular cracks, secondary cracks, tear ridges, and sharp edges (Figure 17a–f), all of which resulted from the interaction of diffused hydrogen atoms with its microstructure. Transgranular cracks (Figure 17d) occur when hydrogen accumulates at both the grain boundaries and within the grains themselves, weakening the atomic bonds and facilitating crack propagation through the grain structure under tensile stress. Consequently, the fracture path follows a transgranular route, with hydrogen promoting brittle failure across the grains.

Secondary cracks (Figure 17e) form at stress concentrators, such as inclusions, defects, or grain boundaries. Hydrogen intensifies these cracks by damaging the material in localized regions, which advances crack initiation and propagation at points of high stress concentration. As earlier discussed, tear ridges arise from localized plastic deformation as cracks propagate through the material. The presence of hydrogen amplifies the material’s susceptibility to this deformation by lowering the energy required for dislocation motion, leading to an irregular, ridge-like fracture surface.

The EDX analysis of four points (Figure 17c) on the fracture surface of arc-DED 308L austenitic stainless steel revealed a high weight % of Fe, followed by Cr and Ni (

Table 6. Summary of EDX analysis results on the fracture surface of arc-DED 308L SS shown in Figure 17c.

	Weight %
Element	Cr	Fe	Ni	Cu
Point 1	32.3	52.8	12.7	2.2
Point 2	29.1	52.8	10.6	7.5
Point 3	24.6	60.0	15.4	0.0
Region 4	19.9	66.6	13.5	0.0

). This composition indicates the predominance of the austenite phase, stabilized by the high Ni content, which, together with Cr, enhances the material’s corrosion resistance and ductility.

Additionally, the fracture surface might also contain small regions of ferrite, even though the material is predominantly austenitic. This ferrite could form as a secondary phase due to the rapid cooling during the arc-DED process, which can induce microsegregation of alloying elements. Ferrite, being a BCC phase, can contain a higher iron content, which aligns with the high weight percent of iron detected in the EDX results.

The EDX analysis of the fracture surface of charged arc-DED 308L austenitic stainless steel, as shown in Figure 17f, reveals intriguing compositional variations across six distinct areas (

Table 7. Summary of EDX analysis results on the fracture surface of arc-DED 308L SS shown in Figure 17f.

Weight %
Element	Si	Cr	Mn	Fe	Ni
Point 1	0.8	23.3	2.6	66.4	6.9
Point 2	0.8	24.3	2.5	65.3	7.1
Point 3	4.7	14.5	50.2	26.4	4.2
Point 4	0.1	21.9	3.1	66.5	8.4
Region 5	0.5	24.4	2.8	65.4	6.9
Region 6	0.7	25.2	0.0	66.6	7.5

). Five of the analyzed locations, points 1, 2, 4, and regions 5 and 6, exhibited a relatively higher weight % of Fe, followed by Cr. This distribution strongly indicates the expected dominance of the austenite phase in these regions, with the high Fe content further confirming a principally uniform austenitic microstructure, consistent with the typical composition of stainless steels. Cr serves a dual purpose as it stabilizes the matrix while also facilitating the formation of a passive chromium oxide film, essential for the alloy’s remarkable corrosion resistance.

In contrast, point 3 displayed a prominent deviation, with its composition dominated by Mn and Fe. The manganese-rich region observed at point 3 may correspond to inclusions or secondary phases, such as manganese sulfides (MnS) or silicates (Tephroite and Rhodonite) [116,117], which occasionally develop in stainless steels during solidification. Tephroite (Mn₂Si0₄) and Rhodonite (MnSiO₄) phases may form under localized conditions where manganese and silicon segregate during rapid solidification. Tephroite, a simple silicate, and Rhodonite, a common inosilicate, are likely to appear in manganese-rich regions with sufficient silicon content, as shown in Table 7, particularly in areas with high thermal gradients, such as grain boundaries or inclusions [118]. These inclusions, although not part of the austenitic matrix, can have a significant influence on fracture behavior by acting as stress concentrators or initiating sites for cracks.

The compositional contrast between the Fe- and Cr-rich areas and the Mn-dominated region stresses the essential intricacy introduced by the 3D printing process. While the majority of the material is likely to retain a stable austenitic structure, the presence of localized Mn-enriched phases highlights the potential for microstructural heterogeneity. Such heterogeneity can have critical implications for the mechanical properties and fracture characteristics of the material, pointing out the need for careful consideration of these variations in the context of performance and reliability.

6. Conclusions

Despite ongoing advancements in hydrogen-related material research, the mitigation of HE in structural alloys remains underexplored, particularly in the context of comparing conventionally wrought and AM materials under identical simulated hydrogen environments for possible integration into gas pipeline infrastructure. In this article, a comprehensive study was conducted to comparatively analyze the tensile behavior of four materials, 44W steel, Monel 400, and 304L austenitic stainless steel, all in wrought form, as well as arc-DED fabricated 308L austenitic stainless steel, in simulated hydrogen environments to evaluate their resistance to HE. The materials demonstrated distinct responses to hydrogen permeation, with the major findings summarized as follows:

• Among the four materials studied, 44W steel exhibited the highest susceptibility to HE, as evidenced by significant reductions in both elongation and toughness. The elongation dropped markedly from 6.9% to 4.8%, representing a 30.4% reduction, while toughness decreased from 51.11 MJ/m³ to 36.96 MJ/m³, a 27.7% decline. Fractographical analysis further confirmed these observations, revealing dimpling and microvoid coalescence indicative of ductile fracture in uncharged specimens, in stark contrast to cleavages, microcracks, transgranular cracks, and sharp-edge cracks characteristic of brittle failure in hydrogen-charged samples. These pronounced reductions highlight the material’s vulnerability to hydrogen-induced degradation, primarily attributed to its BCC crystal structure, which facilitates hydrogen diffusion and promotes the HEDE mechanism.
• Monel 400 demonstrated relatively exceptional resistance to HE, with only minimal changes observed in both %elongation and toughness. The elongation showed a slight reduction from 29.3% to 28.69% (a 2.1% decrease), while the toughness decreased marginally from 171.10 MJ/m³ to 165.69 MJ/m³ (a 3.2% reduction). Fractographic analysis additionally supported these findings, revealing the retention of the material’s intrinsic ductile characteristics even after hydrogen charging. This outstanding resilience was primarily attributed to Monel 400’s FCC crystal structure and high nickel content, which significantly reduces hydrogen diffusion and mitigates its harmful effects.
• The behavior of 304L austenitic stainless steel under hydrogen charging conditions was particularly noteworthy, as its elongation was effectively stable from 28.5% to 28.9% (a 1.4% improvement), as well as its toughness, from 202.51 MJ/m³ to 203.56 MJ/m³ (a 0.5% enhancement). These results suggest that hydrogen may play a stabilizing role within the FCC lattice or alleviate internal stresses, leading to improved ductility and toughness through a mechanism known as HELP. These findings reaffirm 304L austenitic stainless steel’s exceptional suitability for applications in hydrogen-rich environments, where both mechanical integrity and durability are critical, as described in sour service conditions.
• The arc-DED 308L austenitic stainless steel showed a moderate reduction in performance post-hydrogen charging. Specifically, its elongation decreased from 33.6% to 27.4% (an 18.5% reduction), and its toughness dropped from 156.74 MJ/m³ to 149.92 MJ/m³ (a 4.3% reduction). These reductions, though less severe than those observed in 44W steel, are more pronounced than in conventional 304L austenitic stainless steel. The increased susceptibility of the arc-DED material may be attributed to microstructural features inherent to AM, including residual stresses, porosity, and anisotropy, creating hydrogen trapping sites that aggravate embrittlement effects. However, fragmentary analysis showed an initial reduction with increasing hydrogen permeation and a gradual increase at relatively higher hydrogenating conditions. This dual behavior was mechanistically linked to a combination of HEDE, which accounts for the initial reduction, and HELP, which explains the subsequent improvement. These findings highlight the complex interplay of hydrogen-related mechanisms and the microstructural features of arc-DED fabricated materials, emphasizing the need for careful consideration in hydrogen-rich applications.
• Other major tensile properties of the materials studied exhibited distinct behaviors under hydrogenating conditions. For 44W steel, Monel 400, and 304L austenitic stainless steel, the YS, UTS, and elastic moduli remained largely unchanged across all levels of hydrogen charging. This stability is attributed to their favorable crystal structures, chemical compositions, and microstructural features in agreement with preceding studies. Conversely, arc-DED 308L austenitic stainless steel exhibited a gradual increase in YS and UTS under hydrogen charging. Features such as high dislocation density, fine cellular substructures, anisotropic columnar grains, and residual stresses inherent to AM act as hydrogen traps, enhancing strain hardening and impeding dislocation motion. Additionally, hydrogen-induced solid solution strengthening and the interaction with residual stresses may further improve the material’s tensile properties. These combined effects synergize to enhance resistance to plastic deformation and increase stiffness.
• The elastic modulus of arc-DED 308L austenitic stainless steel also showed a slight increase under hydrogenation, unlike the behavior typically observed in its conventionally processed counterparts. This behavior was linked to hydrogen trapping at dislocations and grain boundaries, which impedes localized deformation and reinforces atomic interactions. The heterogeneous microstructure of the material, combined with hydrogen-induced lattice distortions and residual stresses, contributes to an apparent increment in stiffness. Also, the hydrogen interstitial strengthening phenomenon, which explains hydrogen atoms’ potential to alter the stiffness of metallic materials by interacting with their atomic bonds, was quoted as a possible prevailing condition. These results highlight the distinct behavior of arc-DED fabricated materials under hydrogen-rich environments and their potential for enhanced performance in such conditions.
• These findings propose appreciative considerations for material selection and design in hydrogen infrastructure. The pronounced susceptibility of 44W steel to hydrogen embrittlement suggests it is unsuitable for critical applications without mitigation, whereas Monel 400 and 304L stainless steel demonstrate superior resistance, making them strong candidates for components in pipelines, storage tanks, and pressure vessels. The performance of arc-DED 308L stainless steel also highlights the potential of AM to tailor materials for hydrogen service through controlled microstructural features. These results can inform industry standards, material qualification protocols, and engineering strategies as hydrogen energy systems continue to evolve.