Effect of Electrochemical Hydrogen Charging on the Notch Tensile Properties of Natural Gas Transportation Pipeline Steel with Electroless-Plated Coatings and Their Adhesiveness Characterization

Abstract

Traditional natural gas transportation pipeline steels, such as API 5L X42 grade and the higher grades, are currently receiving a lot of attention in terms of their potential implementation in hydrogen transmission infrastructure. However, the microstructural constitution of steels with a ferrite phase and the presence of welds, with their non-polyhedral “sharp” microstructures acting as structural notches, make these steels prone to hydrogen embrittlement (HE). In this work, the notch tensile properties of copper- or nickel–phosphorus-coated API 5L X42 grade pipeline steel were studied in both the non-hydrogenated and electrochemically hydrogen-charged conditions in order to estimate anticipated protective effects of the coatings against HE. Both the Cu and Ni–P coatings were produced using conventional coating solutions for electroless plating. To study the material systems’ HE sensitivity, electrochemical hydrogenation of cylindrical, circumferentially V-notched tensile specimens was performed in a solution of hydrochloric acid with the addition of hydrazine sulfate. Notch tensile tests were carried out for the uncoated steel, Cu-coated steel, and Ni–P-coated steel at room temperature. The HE resistance was evaluated by determination of the hydrogen embrittlement index (HEI) in terms of relative changes in notch tensile properties related to the non-hydrogenated and hydrogen-charged material conditions. The results showed that pure electroless deposition of both coatings induced some degree of HE, likely due to the presence of hydrogen ions in the coating solutions used and the lower surface quality of the coatings. However, after the electrochemical hydrogen charging, the coated systems showed improved HE resistance (lower HEI_RA values) compared with the uncoated material. This behavior was accompanied by the hydrogen-induced coatings’ deterioration, including the occurrence of superficial defects, such as bubbling, flocks, and spallation. Thus, further continuing research is needed to improve the coatings’ surface quality and long-term durability, including examination of their performance under pressurized hydrogen gas charging conditions.

1. Introduction

The feasibility of plans for carbon neutrality and global environmental improvement will strongly depend on the implementation possibilities of green hydrogen technologies, which will enable the transformation of fossil fuel-powered manufacturing sectors into the so-called hydrogen economy [1,2,3]. Currently, one of the main obstacles for the massive introduction of hydrogen technologies into the application sphere is directly related to the degradation effect of hydrogen on the mechanical properties of structural metallic materials. The phenomenon known as hydrogen embrittlement (HE) or hydrogen-induced cracking (HIC) has been known for over 100 years, starting with the pioneering work of Johnson [4]. It has been adopted that HE in metals is caused by the action of free diffusible atomic hydrogen, more specifically the hydrogen cations H+ (in principle, the protons), capable of diffusion into metallic crystal structures even at room temperature. Since then, numerous experimental and theoretical studies have been published on this degradation phenomenon in various types of metallic materials [5,6,7,8,9,10,11,12,13]. In general, the metals and alloys with a face-centered cubic (FCC) crystal structure exhibit a lower HE susceptibility due to their higher hydrogen solubility and lower hydrogen diffusivity compared with the metallic materials, with a body-centered cubic (BCC) crystal structure showing the opposite hydrogen characteristics, i.e., lower hydrogen solubility and higher hydrogen diffusivity [14,15]. At tensile loading, a typical HE is manifested by a reduction in deformation properties, i.e., total elongation and reduction of area [16,17,18,19]. However, a reduction in strength properties, i.e., yield stress and ultimate tensile strength, may also occur in tensile-tested metallic materials as a consequence of the HE [20,21,22,23]. Fracture mechanics studies, focused on various fracture toughness determination approaches, e.g., stress intensity factor, elastic energy release rate, crack-tip opening displacement, J-integral, etc., have also generally confirmed the deleterious effects of hydrogen exposure on resulting values of fracture toughness [24,25,26,27]. The notch tensile strength and reduction of area evaluated from the notch tensile tests also represent a commonly used approach for estimation of the HE resistance of metallic materials with the presence of a stress concentrator. The notch tensile properties also generally decrease as a result of material hydrogenation [11,28,29,30].

The generally known and accepted fundamental HE mechanisms, such as hydrogen-enhanced decohesion (HEDE), hydrogen-enhanced localized plasticity (HELP), adsorption-induced dislocation emission (AIDE), hydrogen-enhanced strain-induced vacancy (HESIV), etc., have been intensively studied by numerous researchers, e.g., [31,32,33,34], on various metallic materials using both the experimental and computational approaches. Accordingly, two main HE model types have been distinguished, namely the “plasticity-driven” HE models (i.e., the models based on the HELP, AIDE, or HESIV mechanisms) and the “decohesion-based” model (i.e., the model based on the HEDE mechanism). Some studies, e.g., [35,36], indicated that the multiple HE mechanisms can occur and be active in the same material. Djukic et al. [37] reviewed the synergistic action and interplay of HELP and HEDE mechanisms in steels and iron. They also proposed a novel and unified HELP+HEDE model based on the specific microstructural mapping of the dominant HE mechanisms with implications on the fracture process and resulting hydrogen-assisted fracture modes [37].

Nowadays, HE represents a highly topical issue, since the idea of blending natural gas with hydrogen in transportation pipeline infrastructure is receiving a constantly increasing amount of attention [38,39,40,41,42,43,44,45]. Although several noble metals and their derived alloys possess a relatively good resistance against HE, from an economic point of view, a realistic material solution for long-distance gas transportation pipelines can only be feasible using lower-cost carbon steels or low-alloy steels [22,46,47,48,49,50]. However, these materials are known to be vulnerable to HE, and thus they need an effective surface protection against hydrogen permeation. Numerous investigations about corrosion and/or HE testing of copper and nickel–phosphorus coatings on mild steel have been published worldwide [51,52,53,54,55,56,57,58]. Although much research effort has already been devoted to such coating–steel material systems, the findings of various research studies are often different or even contradictory. The early study by Chen and Wu [51] was focused on electrochemical hydrogen permeation experiments on copper-plated AISI 4140 steels. They revealed that the applied copper plating decreased the hydrogen absorption rate, which consequently resulted in lowering both the hydrogen permeation rate and its effective diffusivity. Khan et al. [52] studied Cu-based anti-corrosion coatings on steel implemented by surface alloying. They reported a 44% improvement in corrosion resistance for the developed Cu–Ni coating with a Ni-rich transition layer. The published works [53,54,55] presented the research efforts of the Canadian corrosion program for the long-term management of nuclear waste using copper coatings on steel or cast iron as viable corrosion barriers. The study by Samanta et al. [56] was focused on the development of amorphous Ni–P coating over API X70 steel for hydrogen barrier application. They concluded that excellent resistance to hydrogen permeation has been related to the amorphous nature of electroless Ni–P coating. Biggio et al. [57] published a comprehensive review on electroless Ni–P coatings for hydrogen barrier applications on steel. They emphasized the ease of preparation of the Ni–P coatings and suggested their suitability for applications in existing pipelines. On the contrary, the study by Michler and Naumann [58] did not indicate any significant improvements in HE resistance for their tensile test specimens with electroplated Cu and electroless Ni–P coating.

In our recent work [59], promising results have been obtained regarding the effects of electrochemical hydrogen charging on the HE resistance of copper- or nickel–phosphorus-coated X42 grade pipeline steel using smooth (i.e., notch-free) tensile test specimens. The current work represents further continuation of our preceding research about the possibility of HE suppression by using copper or nickel–phosphorus coatings on X42 steel tensile specimens with the presence of a notch. This study deals with the effects of electrochemical hydrogenation on the notch tensile properties of copper- and nickel–phosphorus-coated X42 grade pipeline steel, including the investigation of uncoated material for comparison. The adhesion of the studied electroless-plated coatings with the steel substrate was characterized by tribological scratch tests.

2. Experimental Material and Methods

Commercial API 5L X42 grade pipeline steel (further referred to as X42 steel) was used in the present study as the base material. It was supplied in the form of a hot-rolled pipe with nominal dimensions of an outer diameter of 114.3 mm and a wall thickness of 13.49 mm. The elemental chemical composition of the pipe provided by the steel manufacturer is shown in

Table 1. Elemental chemical composition of studied X42 steel [wt.%].

C	N	Mn	Si	P	S	Cu	Cr	Mo	V	Ni	Al	Sn	Fe
0.16	0.009	0.51	0.24	0.014	0.010	0.19	0.09	0.02	0.007	0.08	0.027	0.012	rest

.

Notch tensile tests were carried out at room temperature using cylindrical tensile specimens with a circumferential V-notch in their central portion (Figure 1). The specimens were produced by lathe machining, with a finished surface roughness of 0.6 µm. The notch was machined by using a turning parting knife.

The notch tensile tests were performed for three investigated material configurations, i.e., the uncoated steel (“X42_uncoated”), copper-coated steel (“X42_Cu-coated”), and nickel–phosphorus-coated steel (“X42_Ni–P-coated”). Non-electrochemical deposition of copper coating onto the prepared clean surface of notch tensile test specimens was carried out by their immersion into a solution of H₂O (1000 mL) with CuSO₄ (50 g) and concentrated H₂SO₄ (50 g) for 60 s. Afterwards, the Cu-coated notch tensile specimens were rinsed with water and dried. Electroless deposition of nickel–phosphorus coating onto the prepared clean surface of notch tensile test specimens was carried out by their immersion into a warm solution of H₂O (1000 mL) with (C₃H₅O₃)₂Ni (50 g) and NaH₂PO₂∙H₂O (15 g), maintained at 90 °C for 90 min. Subsequently, the Ni–P-coated notch tensile specimens were rinsed with water and dried. The average thickness of prepared Cu and Ni–P coatings was 1.7 µm and 8.8 µm, respectively [59]. The prepared coatings reflected limitations of the used non-electrochemical deposition processes.

For testing the adhesiveness and frictional behavior of prepared coatings on the steel substrate, tribological scratch tests were carried out employing a universal tribometer and a scratch tester Bruker Mod. UMT 2M (Bruker-Nano Surfaces TMT Unit, Campbell, CA, USA) with Vickers diamond indenter and acoustic emission (AE) detector model AE-5. The scratch tests were performed on plain surfaces of coated prismatic specimens (approx. 5 × 10 × 25 mm in size) within the load range from 1 to 30 N according to the ASTM C1624 −22 standard [60]. Three tribological scratching tests were performed for each studied coating.

The HE susceptibility of all studied material configurations was evaluated by performing conventional quasi-static tensile tests of prepared notch tensile specimens in conditions without and with electrochemical hydrogen charging at room temperature. Cathodic hydrogen charging was conducted in a solution of 1M HCl and 0.1N N₂H₆SO₄ at a current density of 200 A/m² for 24 h, employing a potentiostat/galvanostat model 173 (Princeton Applied Research, Oak Ridge, TN, USA). Prior to tensile testing, the hydrogen-charged tensile specimens were put into a thermal insulating flask with liquid nitrogen in order to avoid quick hydrogen desorption. The room-temperature tensile tests were performed at a crosshead travel speed of 0.2 mm/min, employing the universal testing machine TIRATEST 2300 (TIRA GmbH, Schalkau, Germany) in accordance with standard ISO 6892-1:2019 [61]. For all material configurations (“X42_uncoated”, “X42_Cu-coated”, and “X42_Ni–P-coated”) in each material condition (non-hydrogenated and hydrogen-charged), three individual notch tensile tests were carried out, and average notch tensile properties, i.e., notch tensile strength (NTS) and reduction of area (RA), were evaluated. The values of the hydrogen embrittlement index (HEI) were evaluated as relative changes in individual notch tensile properties according to the following equation:

$$\mathrm{HEI} = {\displaystyle \frac{{\mathrm{X}}_0-{\mathrm{X}}_{\mathrm{H}}}{{\mathrm{X}}_0}} \times 100 \%$$

(1)

where X₀ is the average value of an individual notch tensile property, e.g., NTS or RA of broken tensile test specimens in a non-hydrogenated material state, and X_H is the average value of the same notch tensile property of broken tensile test specimens in a hydrogen-charged material state. The electrochemical hydrogen charging method provides a cost-effective and simpler alternative to gaseous hydrogen charging, which requires high temperatures and pressures (e.g., 150 °C and 40 MPa, as reported in [62]), as well as specialized and expensive equipment subject to rigorous safety standards for implementation [13,48,63]. It was revealed that any change in a parameter increasing the hydrogen concentration in the steel microstructure also enhances HE sensitivity [48].

The microstructural and phase analyses were performed using the scanning electron microscope (SEM) JEOL JSM-7000F (Jeol Ltd., Tokyo, Japan) with the electron back-scattered diffraction (EBSD) detector Nordlys-I (Oxford Instruments plc, Abingdon, Oxfordshire, UK). The EBSD phase analyses were performed by means of Kikuchi diffraction pattern analysis. The EBSD analyses were performed on a drawing direction plane of prepared metallographic cross-section, and the obtained EBSD data were further processed using the CHANNEL-5, HKL software package (service pack 7, HKL technology A/S, Hobro, Denmark). Crystallographic data (unit cell parameters, space group, and the chemical composition of the phase and atom Wyckoff position in a crystal lattice) for individual anticipated phases were taken from the EBSD software database. Fractographic analyses of the fracture surfaces of broken tensile test specimens and morphological characterization of tribological scratching tracks were performed using SEM Tescan Vega-3 LMU (Tescan Brno, s.r.o., Brno, Czech Republic).

3. Results and Discussion

3.1. Microstructural and Phase Characterization

Figure 2 shows the SEM microstructure and corresponding EBSD analyses of the X42 grade pipeline steel under investigation.

An overall view of the polygonal ferritic–pearlitic grain microstructure of the as-received material is shown in Figure 2a. A detailed view of the microstructure consisting of Fe-α ferrite phase (gray grains) and pearlite microstructural constituent with alternating Fe₃C carbide/Fe-α ferrite (bright/gray lamellae structure) is shown in Figure 2b. According to our former microstructural investigations performed [59], the mean grain sizes of individual microstructural constituents are 23.9 µm and 19.4 µm for ferrite and pearlite, respectively. The corresponding area fractions are 81.6% ferrite and 18.4% pearlite. Selected examples of randomly recorded Kikuchi EBSD patterns related to the BCC-structured Fe-α ferrite matrix and Fe₃C carbide with orthorhombic crystal structure are shown in Figure 2c and Figure 2d, respectively. With respect to the chemical composition of the investigated steel, some occurrence of common non-metallic inclusions such as Al₂O₃, MnS, SiO₂, and calcium aluminates is also to be expected. AlN, as another expected minor phase, is a common precipitate in micro-alloyed steels, primarily serving to pin the austenite grain boundaries. However, only AlN and MnS were confirmed in the studied steel by means of thermodynamic calculations in our previous work [59]. The inclusions and precipitates acting like hydrogen-trapping sites (irreversible or reversible) may have either positive or negative effects on the resulting HE resistance. However, the emphasis of the present investigation is placed on the effects of two electroless-plated coatings, which do not affect the metallurgical state of non-metallic inclusions and precipitates within the studied steel.

Figure 3 shows photographic images of the notch tensile specimens corresponding to individual material surface configurations.

The EDX chemical micro-analyses of free surfaces of both coated material configurations revealed that the coating on Cu-coated specimens was exclusively formed of pure copper, whereas the coating on Ni–P-coated specimens was formed of the mixture of about 81 at.% nickel and 19 at.% phosphorus [59]. The X-ray diffraction phase analysis revealed that the copper coating possesses a fully crystalline face-centered cubic (FCC) crystal structure, whereas the nickel–phosphorus coating is of an amorphous nature [59].

3.2. Tribological Scratching Behavior

The adhesiveness and frictional behavior of both as-deposited coatings were tested by means of tribological scratch tests at room temperature, and the obtained results are shown in Figure 4, Figure 5, Figure 6 and Figure 7. A morphological characterization of the tribological scratching track on Cu-coated steel is depicted in Figure 4.

Morphological analysis of the wear track (Figure 4) revealed a gradual sequence of mechanisms, ranging from plastic grooving through cohesive failure to adhesive delamination. In accordance with the tribological interpretations proposed by Burnett and Rickerby [64], individual stages are related to the achievement of several critical loads. During the initial stage corresponding to the first critical load Lc1, the coating exhibited ductile behavior with material displacement but without delamination (Figure 4a). At higher loads reaching the second critical load Lc2, cracks were perpendicular to the sliding direction and localized delamination appeared (Figure 4b), confirming the onset of cohesive failure. In the final stage corresponding to the critical load Lc3, complete separation of the copper from the substrate occurred (Figure 4c), accompanied by the spallation of larger coating fragments, while the track symmetry indicated stable contact of the indenter with the substrate.

Figure 5 shows representative behaviors of the coefficient of friction (COF) and acoustic emission (AE) signal during the tribological scratch tests performed on Cu-coated steel.

The evolution of the coefficient of friction (COF) and acoustic emission (AE) correlated well with the observed morphology. Within the load range of 1–2 N, the COF fluctuated between 0.33 and 0.42, due to the initial ploughing of the indenter into the soft coating overcoming the surface topography, while the AE signal remained at a low level. In the range of 2–17 N, stable sliding occurred, with a COF of ~0.38–0.40 and a single significant AE impulse at Lc2 (~17 N), which was associated with cohesive delamination. In the load interval of 17–24 N, the plastic flow of copper dominated, the COF stabilized around 0.39, and the AE signal remained close to the noise level, confirming the ductile characteristics of the coating. The critical load Lc3 (25–30 N) corresponded to complete delamination and final coating failure. Overall, the results confirm that the copper coating on steel exhibits distinctly ductile behavior, which ensures silent deformation and stable sliding. The presence of a single dominant AE impulse indicates good, although not exceptional, adhesion to the substrate.

A morphological characterization of the tribological scratching track on Ni–P-coated steel is depicted in Figure 6.

The morphology of the scratch track of the Ni–P coating on steel demonstrates the typical behavior of a hard and brittle coating. In the initial stage corresponding to the first critical load Lc1, a brittle response with minimal plastic deformation was observed, with minimal plastic deformation and cutting of the material without flow (Figure 6a). At higher loads reaching the second critical load Lc2, wedge-shaped spallation and transgranular micro-cracks appeared, indicating the onset of cohesive failure (Figure 6b). The final stage corresponding to the critical load Lc3 was characterized by combined adhesive–cohesive failure, with the spallation of larger coating segments and local exposure of the substrate (Figure 6c). The lower pile-up accumulation compared to copper indicates the higher hardness and lower ductility of the Ni–P coating. The critical load Lc3 for the Ni-P coating was lower than that of Cu, suggesting that, while it provides a hard and durable surface, its adhesion to the substrate under progressive loading is less robust than that of the ductile copper coating.

Figure 7 shows representative behaviors of the coefficient of friction (COF) and acoustic emission (AE) signal during the tribological scratch tests performed on Ni–P-coated steel. The evolution of the coefficient of friction (COF) and acoustic emission (AE) further supports the above-discussed morphological observations. In the load range of 1–6 N, the COF sharply increased to ~0.40, while the AE signal reached 10 V, corresponding to the formation of numerous micro-cracks. At loads of 6–12 N, pre-spalling occurred, accompanied by strongly fluctuating AE activity. In the interval of 12–24 N, the COF remained stable (~0.39), corresponding to the steady propagation of micro-cracks without significant material accumulation.

At ~24 N (Lc3), a complete delamination of the coating occurred, accompanied by transgranular fracture and a sharp spike in the AE signal, which indicated catastrophic failure. Beyond this stage, the indenter slid directly on the substrate, with residual Ni–P fragments. Overall, the Ni–P coating exhibits hard and brittle behavior with a stable COF and high AE activity (~10 V), distinct from the ductile behavior of Cu. The lower critical load Lc3 of Ni–P coating compared to Cu coating suggests its inferior adhesive strength, although its high hardness is confirmed.

3.3. Notch Tensile Properties and HE Sensitivity

Figure 8 shows typical engineering stress–strain curves for the studied material configurations (i.e., X42_uncoated, X42_Cu-coated, and X42_Ni–P-coated) in non-hydrogenated and hydrogen-charged material states.

All studied material configurations show pronounced yielding behavior, i.e., deformation instabilities at the onset of plastic deformation. The observed behavior can be generally related to overcoming the pinning forces in so-called Cottrell atmospheres (i.e., carbon and nitrogen atoms) by moving dislocations [65]. Moreover, in the present case, this behavior is even more complex due to the additional effect of the tensile specimen geometry, i.e., the stress-concentrating effect of the notch.

Figure 9 depicts mean values of notch tensile properties for both the non-hydrogenated and hydrogen-charged material states, which were subsequently used for the calculation of corresponding values of the hydrogen embrittlement index. The resulting HEI values calculated according to the changes in notch tensile strength (HEI_NTS) and reduction of area (HEI_RA) are shown in Figure 10.

From the results shown in Figure 9, it is obvious that the effects of the coatings and hydrogen charging procedure did not result in any significant changes in the measured values of notch tensile strength. On the contrary, the aforementioned effects induced remarkable changes in the values of the reduction of area. The application of both coatings in non-hydrogenated material conditions resulted in a clear decrease in the reduction of area compared with the uncoated material. This observation is likely related to the application of non-electrochemical coatings itself, inducing some degree of hydrogen embrittlement due to the presence of hydrogen ions in the coating solutions used for electroless plating. Moreover, the lower surface quality of the produced coatings [59] likely enhanced this behavior. On the other hand, both coated material configurations after hydrogen charging showed just the opposite tendency, i.e., the higher values of reduction of area in comparison with the uncoated material exposed to hydrogen charging. This observation indicates a certain potential for protective barrier effects of both studied coatings against hydrogen permeation. The determined values of the hydrogen embrittlement index (Figure 10) can be interpreted in a similar way, i.e., the HEI values calculated according to the changes in notch tensile strength (HEI_NTS) do not show any significant variations among individual material configurations, whereas the HEI values calculated according to the changes in reduction of area (HEI_RA) show significant changes in favor of the coated systems. In other words, the HEI_RA values of the coated material configurations were significantly (almost four times) lowered in comparison with the uncoated material, which indicated a highly improved HE resistance (Figure 10). More importantly, HEI_RA criterion for HE resistance should be considered more relevant, since the RA values characterize the material failure resistance as they are directly related to the deformation required to produce fracture [59].

As already shown in Figure 4 and Figure 6, the deposited coatings do not possess completely smooth and defect-free surface structures. Instead, various surface defects have been observed. These may serve as preferential degradation sites during hydrogen charging and afterwards as failure nucleation sites during tensile loading. In addition, Figure 11 shows that the adhesiveness of the coatings became highly impaired after the applied electrochemical hydrogen charging, especially in the case of Cu coating showing local spallation areas (see Figure 11a). In the case of Ni–P coating, a formation of superficial bubbles and/or flocks has been observed as a consequence of hydrogen charging (see Figure 11b). It is believed that the observed deterioration occurred thanks to the intensive electrochemical reaction of cathodic hydrogen evolution, leading to the weakening of the bonding strength at the coating/steel matrix interface.

Despite this unfavorable phenomenon observed, the obtained results from the notch tensile tests (Figure 9 and Figure 10) clearly indicated that both applied coatings acted as hydrogen permeation barriers, suppressing the HE of the steel substrate within the timescale of the present investigation.

3.4. Fracture Micro-Mechanisms

Characterization of fracture micro-mechanisms for individual material configurations (uncoated steel, Cu-coated steel, and Ni–P-coated steel) was carried out by observation of fracture surfaces of ruptured notch tensile test specimens in studied material conditions with respect to hydrogen charging application (see Figure 12).

The fracture micro-mechanism related to the uncoated and non-hydrogenated X42 material after the notch tensile test consists almost completely of ductile dimple tearing produced by microvoid coalescence. Shear bands and a little portion of transgranular quasi-cleavage fracture areas are also visible. (Figure 12a). This observation is in good agreement with the findings of other research studies dealing with fracture micro-mechanisms of ferritic–pearlitic steels [66,67]. The observed microcavities within the ferritic grains are preferentially nucleated at the ferrite/pearlite grain boundaries. The size and shape of the dimples on the fracture surface strongly depend on the local microstructural features influencing the transgranular propagation of cracks leading to the final fracture of tensile specimens [68]. The fracture micro-mechanism related to the uncoated and hydrogen-charged X42 steel after the notch tensile test shows a mixed fracture micro-mechanism consisting of ductile dimple tearing (i.e., microvoid coalescence fracture), as well as an increased portion of brittle transgranular quasi-cleavage with a typical “river marks” morphological pattern (Figure 12b). This fracture behavior may indicate the complex HELP+HEDE mechanism of hydrogen embrittlement as also discussed in [37]. On the other hand, both coated material systems exhibit mostly ductile dimple tearing fracture micro-mechanisms (i.e., microvoid coalescence fracture) with both finer and larger dimples including minor portions of transgranular quasi-cleavage areas regardless of hydrogen charging application (Figure 12c–f). Thus, the fractographic characteristics of the coated material configurations are quite similar to those of the original non-hydrogenated X42 steel without coating. This observation supports the obtained findings of the present investigation, indicating that the applied Cu and Ni–P coatings on X42 pipeline steel may efficiently suppress its HE susceptibility. Future investigations should be focused on the studies under gaseous hydrogen charging conditions with respect to the mechanical behavior and improvement of surface quality, especially the adhesiveness and long-term durability of the coatings on the steel substrate.

4. Conclusions

This study was focused on the effects of electrochemical hydrogen charging on the notch tensile properties of copper- and nickel–phosphorus-coated API 5L X42 grade pipeline steel. To distinguish the effects of the coatings themselves, the study also included the investigation of reference (uncoated) material. The coatings’ adhesiveness with the steel substrate was characterized by tribological scratch tests. Here are the main conclusions:

• Neither the application of coatings nor electrochemical hydrogenation resulted in any significant changes in notch tensile strength. However, both aforementioned effects induced noticeable changes in the reduction of area. The sole application of individual coatings without hydrogen charging resulted in a clear decrease of reduction of area compared to the uncoated material, which might be related to hydrogen embrittlement due to the presence of hydrogen ions in the coating solutions as well as the observed superficial imperfections in produced coatings.
• Both coated material configurations after electrochemical hydrogenation exhibited higher reduction of area in comparison with the hydrogenated material without coating, which indicated certain protective barrier effects of both investigated coatings against hydrogen permeation. These findings have been supported by the performed fractographic observations.
• The HEI values calculated according to the changes in notch tensile strength (HEI_NTS) do not show any significant variations among individual material configurations. On the other hand, the HEI values of the coated material configurations calculated according to the changes in reduction of area (HEI_RA) were remarkably (almost four times) reduced in comparison with the uncoated material.
• The results of tribological scratch tests indicated stable adhesive bonding with the steel substrate in the initial material condition, i.e., the state without application of electrochemical hydrogen charging. The frictional and acoustic emission behaviors of the individual coatings studied corresponded fairly with their microstructural characteristics.
• After the electrolytic hydrogen charging, the adhesiveness of the coatings was lowered. Bubbles and flocks occurred within hydrogenated coatings. Moreover, a local spallation of the Cu coating also occurred after the electrochemical hydrogenation. Thus, further research efforts should be focused on improvements in the coatings’ surface quality and long-term durability, including the study of their cyclic mechanical behavior in pressurized hydrogen gas environments.