Introduction

Materials

Table 1. Chemical composition of AO-C1 and N-C1 pipeline steels (wt %).

Figure 1. Microstructure of a) AO-C1 steel b) N-C1 steel. Etchant: nital (5%).

Figure 1. Microstructure of a) AO-C1 steel b) N-C1 steel. Etchant: nital (5%).

Experimental

Test Specimens

Both tensile and fatigue specimens were machined from the longitudinal direction of the pipe sections, as illustrated in Error! Reference source not found..

Figure 2. Specimen orientation.

Figure 2. Specimen orientation.

The geometries of the tensile and fatigue specimens are illustrated in Error! Reference source not found.a,b respectively. These designs were influenced by limitations related to pipe thickness, as well as the capabilities of the machining and testing equipment.

Figure 3. Specimen geometries of a) Tensile specimen b) Fatigue specimen; All dimensions are in mm.

Figure 3. Specimen geometries of a) Tensile specimen b) Fatigue specimen; All dimensions are in mm.

Electrochemical Hydrogen Charging

To induce hydrogen absorption, an electrochemical charging process was employed using the test set up shown in Error! Reference source not found.. The test cell was filled with a 0.1 M NaOH electrolyte containing 3.3 g/L Ammonium Thiocyanate (NH₄SCN) as a poison to prevent the recombination of hydrogen atoms into H₂ [25]. The test cell was equipped with two one-sided valves for purging, as well as openings for pH and temperature meters, and it was fully sealed. Specimens were cleaned with acetone before insertion into the cell, and all areas except the gauge area were covered with insulating tape to ensure that only the calculated surface area was exposed to the electrolyte and electrochemically charged. Before the start of the charging process, the test cell was purged with argon gas to remove all oxygen. The charging process was carried out by a potentiostat applying a constant current to promote the hydrogen evolution reaction at the steel surface. The specimen served as the working electrode, while platinum was used as the counter electrode. The selection of electrolyte, charging current density, and charging time was based on permeation method outlined in the authors previous work [25] on the same materials using a double electrolyte cell based on the Devanathan-Stachursky method [54], in accordance with ISO 17081 [55]. The permeation tests provided effective diffusivity (D_eff), hydrogen concentration in interstitial lattice and reversible trap sites (C_0R), and time to steady state (t_ss) for both steels, as shown in Error! Reference source not found.. The charging current density for both steels was 10 mA/cm².

Table 2. Output of permeation tests.

Table 2. Output of permeation tests.

	Charging current density (mA/cm2)	Deff (m2/s)	C0R (mol/m3)	tss (min)
AO-C1	10	5.67 e-11	0.49	73
N-C1	10	3.38 e-11	3.14	123

Several researchers [56,57,58] have noted a correlation between hydrogen concentration in electrochemical charging and the square root of the charging current density, described by the equation:

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

Using the values of charging current density and hydrogen concentration from the permeation tests, the constant factor K was calculated for each steel. This constant was then used to correlate different hydrogen concentrations with charging current densities. Error! Reference source not found. summarizes the relationships between hydrogen content, charging current densities, and charging times for the electrochemical hydrogen charging experiments. To ensure maximum hydrogen saturation, the charging time was set to 1.5 times the steady state time (t_ss) determined from the permeation tests.

Table 3. Electrochemical hydrogen charging parameters.

Table 3. Electrochemical hydrogen charging parameters.

CH (wppm)	AO-C1		N-C1
	Charging time (min)	Charging current density (mA/cm2)	Charging time (min)	Charging current density (mA/cm2)
0.07	110	12.3	185	0.3
0.2		100.4		2.5
0.5		627.8		15.6
1.0		2511.2		62.5
1.4		-		122.5
2.0		-		249.9

Figure 4. Experimental setup for electrochemical hydrogen charging.

Figure 4. Experimental setup for electrochemical hydrogen charging.

Results and Discussion

Tensile Behavior

The stress-strain curves for the uncharged and charged tensile specimens from AO-C1 and N-C1 pipeline steels are illustrated in Error! Reference source not found.a-b, with corresponding tensile data presented in Error! Reference source not found.a-d. The results from tensile tests on uncharged samples show that AO-C1 exhibits higher ductility, whereas N-C1 demonstrates greater yield strength (YS) and ultimate tensile strength (UTS). Additionally, it is notable that the difference between YS and UTS is relatively small in N-C1, suggesting that this steel will reach final fracture more quickly when subjected to tensile loads beyond the elastic region, compared to AO-C1. This implies a more abrupt transition to failure under plastic deformation for N-C1.

Figure 5. CSRT results for Uncharged and charged with varying hydrogen contents for a) AO-C1 steel; b) N-C1 steel.

Figure 5. CSRT results for Uncharged and charged with varying hydrogen contents for a) AO-C1 steel; b) N-C1 steel.

Error! Reference source not found.a,c and illustrate the yield strength (YS) and ultimate tensile strength (UTS) for both uncharged and hydrogen-charged samples with varying hydrogen content. It is evident from the figures that the YS of both materials increases after hydrogen charging compared to their uncharged state. This observation is consistent with findings from previous studies [59,60]. The increase in YS is typically attributed to the solid solution strengthening effect of interstitial hydrogen. When hydrogen atoms occupy interstitial sites within the metal lattice, they hinder dislocation movement, thereby enhancing the material's resistance to deformation and increasing its yield strength [61]. Within the range of hydrogen contents tested in this study, the maximum difference between the tensile strength of precharged and uncharged AO-C1 samples is 36%, and for N-C1 samples, it is 16%. The more substantial increase for AO-C1 steel may be due to degradation over its extended service period, as the yield strength of the uncharged samples (290 MPa) is below the expected grade standard of 359 MPa. The maximum difference in UTS between uncharged and charged samples is 9% for AO-C1 and 6% for N-C1. Thus, while hydrogen charging significantly impacts the yield strength, its effect on the ultimate tensile strength is relatively minor, aligning with findings from aforementioned studies.

Error! Reference source not found.a,c and illustrate the yield strength (YS) and ultimate tensile strength (UTS) for both uncharged and hydrogen-charged samples with varying hydrogen content. It is evident from the figures that the YS of both materials increases after hydrogen charging compared to their uncharged state. This observation is consistent with findings from previous studies [59,60]. The increase in YS is typically attributed to the solid solution strengthening effect of interstitial hydrogen. When hydrogen atoms occupy interstitial sites within the metal lattice, they hinder dislocation movement, thereby enhancing the material's resistance to deformation and increasing its yield strength [61]. Within the range of hydrogen contents tested in this study, the maximum difference between the tensile strength of precharged and uncharged AO-C1 samples is 36%, and for N-C1 samples, it is 16%. The more substantial increase for AO-C1 steel may be due to degradation over its extended service period, as the yield strength of the uncharged samples (290 MPa) is below the expected grade standard of 359 MPa. The maximum difference in UTS between uncharged and charged samples is 9% for AO-C1 and 6% for N-C1. Thus, while hydrogen charging significantly impacts the yield strength, its effect on the ultimate tensile strength is relatively minor, aligning with findings from aforementioned studies.

Figure 6. Effect of hydrogen content on tensile properies of specimens: a,b) AO-C1 steel ; c,d) N-C1 steel.

Figure 6. Effect of hydrogen content on tensile properies of specimens: a,b) AO-C1 steel ; c,d) N-C1 steel.

The elongation and toughness at fracture of the steels are shown in Error! Reference source not found.b&d. For the AO-C1 sample, the total elongation decreased from 15.5% for the uncharged sample to 9% for the charged sample with 0.5 wppm hydrogen content. In the case of N-C1 steel, the elongation for the uncharged sample and the samples with 0.5 and 1.4 wppm hydrogen content were 13%, 9.7%, and 8.7%, respectively. These findings demonstrate that the total elongation at fracture for both steels is significantly reduced in hydrogen-charged samples compared to those without hydrogen charging, consistent with the phenomenon of HE, where the presence of hydrogen reduces ductility and toughness in steels and is well-documented in previous studies [29,31,34,61,62,63]. A similar trend is observed in the toughness values. For samples charged with 0.5 wppm hydrogen content, the toughness values decreased by 33% for AO-C1 steel and 25% for N-C1 steel.

To summarize the tensile test results, it is evident that both steels experience a significant loss of ductility and toughness due to hydrogen HE. However, at the same hydrogen content (i.e. 0.5 wppm), AO-C1 steel exhibits a more pronounced impact. It is also important to note that similar hydrogen content values are not achieved with similar charging current densities. As shown in Error! Reference source not found., the equivalent charging current density for achieving 0.5 wppm hydrogen content was 627.8 mA/cm² for AO-C1 steel, compared to 15.6 mA/cm² for N-C1 steel. This is also why AO-C1 steel was tested within a lower range of hydrogen contents due to the potentiostat's maximum current capacity constraints. The difference in charging current densities and the corresponding hydrogen content can be attributed to the deformation states of the steels. N-C1 steel, being heavily deformed (cold rolled), has more defects, allowing for more hydrogen trapping which leads to lower diffusivity and higher hydrogen concentration [64,65,66] as reflected in Error! Reference source not found.. However, reversible traps do not contribute to the HE phenomenon, which justifies the lower loss of ductility in N-C1 steel compared to AO-C1 steel at the same hydrogen content. Further explanation on this issue is provided in the next section on the results for fatigue tests.

Fractography of Tensile Samples

The detailed examination of fracture surfaces provides essential insights into the failure mechanisms of tested materials. Fracture surfaces of tensile specimens, both uncharged and hydrogen-charged, were analyzed using scanning electron microscopy (SEM). Error! Reference source not found.a,d show the fracture surfaces of uncharged and hydrogen-charged (1.2 wppm hydrogen content) N-C1 steel, respectively. Side views of the samples are presented in Error! Reference source not found.b,e.

The ductile fracture of the uncharged sample is effortlessly characterized by a reduced fracture area and a distinctive cup-and-cone appearance, as seen in Error! Reference source not found.a,b. Ductile fracture typically occurs in two stages: initial necking and void formation, followed by void coalescence and eventual shear lip formation. The cup-and-cone morphology features a central fibrous region surrounded by a shear lip at an angle, indicative of significant plastic deformation and energy absorption prior to fracture. In the central regions, the fracture surface is characterized by dimple fracture features (Error! Reference source not found.c). These dimples result from the nucleation, growth, and coalescence of micro-voids, reflecting the material's ability to undergo substantial plastic strain before failure. The morphology of the shear lip region in the uncharged sample is dominated by shear dimples that are elongated and oriented along the direction of shear (Error! Reference source not found.a). Compared to the central fibrous region, the shear lip region tends to have a smoother and more streamlined appearance.

Figure 7. SEM fracture surfaces of N-C1 tensile samples; a) CH=0 (Uncharged) b) Side view of “a” c) Magnification of green area in “a” d) CH=1.2 wppm e) Side view of “d” f) Magnification of green area in “d” .

Figure 7. SEM fracture surfaces of N-C1 tensile samples; a) CH=0 (Uncharged) b) Side view of “a” c) Magnification of green area in “a” d) CH=1.2 wppm e) Side view of “d” f) Magnification of green area in “d” .

In contrast, the fracture surface of the sample charged with 1.2 wppm hydrogen content shows two distinctive regions: the central and marginal (or edge) zones, with remarkably different characteristics. The cup-and-cone appearance with a shear lip is not observed; instead, the edges of the fracture surface are sharp (Error! Reference source not found.e), highlighting the complexity of the fracture process. The central region (Error! Reference source not found.f) is characterized by dimple fracture features similar to those of the uncharged sample, but the dimples are shallower and some microcracks are present, indicating brittleness imposed by hydrogen diffusion. Additionally, the size of the central zone is considerably smaller compared to the uncharged sample. The most apparent difference in the fracture surface of the hydrogen embrittled sample can be observed in the marginal or edge zone. During charging, hydrogen forms on the surface of the material and gradually permeates from the exterior to the interior, resulting in higher hydrogen concentrations at the exterior zones. In the marginal zones, brittle fracture characteristics are distinctly observed, including cleavage facets (Error! Reference source not found.b), quasi-cleavage facets, and intergranular fracture (Error! Reference source not found.c,d).

Cleavage facets form in steel due to hydrogen embrittlement from the highly localized and rapid propagation of cracks along specific crystallographic planes. Hydrogen atoms diffuse into the steel lattice, weakening atomic bonds and causing embrittlement. When stressed, these weakened regions are more prone to fracture, and cracks propagate along crystallographic planes, creating cleavage facets. Stress concentrations or material defects can initiate cracks along these preferred planes, which then propagate rapidly through the material along the path of least resistance. Consequently, the fracture surface exhibits flat, shiny facets perpendicular to the applied stress direction. Quasi-cleavage facets resemble cleavage facets but are less well-defined and may appear slightly curved. Intergranular fracture occurs along grain boundaries, indicating a failure mode where hydrogen-induced embrittlement weakens these boundaries, making them preferential paths for fracture propagation.

Error! Reference source not found.a-e show the fracture surfaces of an uncharged sample and a sample charged with 0.5 wppm hydrogen content of AO-C1 steel. The reduced fracture area can be observed by comparing Error! Reference source not found.a and Error! Reference source not found.e. While the reduction in area for the uncharged sample is approximately 63%, it is only 35% for the charged sample. Furthermore, the fracture surface of the uncharged sample exhibits ductile dimple features, as shown in Error! Reference source not found.b. In contrast, the fracture surface of the hydrogen-charged sample displays features indicative of brittle fracture induced by HE, such as cleavage facets, microcracks (Error! Reference source not found.d), quasi-cleavage facets, and river marks (Error! Reference source not found.e).

Figure 8. SEM fracture surfaces of N-C1 tensile samples; a) Magnification of red area marked by “a” in Figure 7a b) Magnification of red area marked by “b” in Figure 7d c) Magnification of red area marked by “c” in Figure 7d d) Magnification of red area marked by “d” in Figure 7d.

Figure 8. SEM fracture surfaces of N-C1 tensile samples; a) Magnification of red area marked by “a” in Figure 7a b) Magnification of red area marked by “b” in Figure 7d c) Magnification of red area marked by “c” in Figure 7d d) Magnification of red area marked by “d” in Figure 7d.

Figure 9. SEM fracture surfaces of AO-C1 tensile samples; a) CH=0 (Uncharged) b) Magnification of red area marked by “b” c) CH=1.2 wppm d) Magnification of red area marked by “d” e) Magnification of red area marked by “e”.

Figure 9. SEM fracture surfaces of AO-C1 tensile samples; a) CH=0 (Uncharged) b) Magnification of red area marked by “b” c) CH=1.2 wppm d) Magnification of red area marked by “d” e) Magnification of red area marked by “e”.

Fatigue Behavior

Error! Reference source not found.a,b shows the fatigue results of AO-C1 samples in bar graph format. It is important to note that the specimens were machined from a section of pipe with a prolonged service history. Thus, the relatively low number of fatigue cycles reflects the remaining fatigue life of the specimens and cannot be directly compared to steels of a similar grade with no service history. Despite the inherent scatter typically seen in fatigue test results, it is evident that specimens have experienced a substantial decrease in fatigue life with increasing hydrogen content. Specifically, the specimen exhibited approximately a 70% reduction in fatigue life at 1 wppm hydrogen content. This significant decline in fatigue life due to HE aligns with the findings of Meng et al. [50] and Faucon et al. [51], however they employed different methodologies. Both studies utilized in-situ gaseous hydrogen charging, contrasting with the electrochemical hydrogen charging method employed in this study, and while Meng estimated fatigue life using fatigue crack growth rate models, Faucon conducted actual fatigue life tests.

Figure 10. Effect of hydrogen content on fatigue life of AO-C1 specimens tested at 80% YS.

Figure 10. Effect of hydrogen content on fatigue life of AO-C1 specimens tested at 80% YS.

Fatigue test results for N-C1 specimens are shown in Error! Reference source not found.a,b. Error! Reference source not found.a presents the results for specimens tested at 80% of YS of the material, while Error! Reference source not found.b shows the results for specimens tested at 50% of YS. At 50% YS, the maximum stress at the specimen's neck center is equivalent to the imposed stress on AO-C1 specimens (278 MPa). This means that when tested at the same stress levels, N-C1 specimens demonstrated a higher fatigue life compared to AO-C1 samples, with uncharged N-C1 samples showing approximately 1.5 times more fatigue cycles. However, it should be noted that AO-C1 specimens have a considerably prolonged service history.

Figure 11. Effect of hydrogen content on fatigue life of N-C1 specimens tested at; a) 50% YS b) 80% YS.

Figure 11. Effect of hydrogen content on fatigue life of N-C1 specimens tested at; a) 50% YS b) 80% YS.

When tested at 80% YS, N-C1 specimens exhibited significantly fewer fatigue cycles compared to AO-C1 samples tested at 80% YS. Based on tensile tests, the ductility of AO-C1 is higher than N-C1, and it is generally believed that ductile materials exhibit better fatigue resistance, though fracture toughness is a more accurate indicator of this property [67]. Moreover, the poor fatigue resistance of N-C1 steels at higher values of stresses can be attributed to their microstructure, which includes numerous pores as shown in Error! Reference source not found.. In uncharged samples, these pores can act as stress concentration points, reducing the fatigue life of the specimens.

Similar to AO-C1 samples, the fatigue life of N-C1 samples decreased after hydrogen charging. However, the impact of hydrogen content on fatigue life reduction is not as severe in N-C1 samples. For instance, at 1 wppm hydrogen content, the fatigue life decreased by approximately 23%, significantly lower than the 70% reduction observed in AO-C1 samples at the same hydrogen content when both were tested at 80% of their respective yield strengths. This difference can be attributed to the presence of pores in the microstructure of N-C1 steel. Generally, pores can act as either reversible or irreversible trapping sites for hydrogen, depending on their specific characteristics. Yaktiti et al. [68] investigated the role of porosities in hydrogen diffusion in steel, demonstrating that pores act as reversible trap sites. These pores delay hydrogen diffusion, and the trapped hydrogen can be desorbed from these sites at room temperature. Therefore, the higher hydrogen content in N-C1 steel can be attributed to reversible hydrogen trapped in porosities that desorb during testing, not contributing to the hydrogen embrittlement of the specimens.

Figure 12. Observed porosities in N-C1 microstructure; a) As polished (not etched) specimen to avoid impact from etchant b) Magnification of the yellow square in “a” c) 3d profile of the pores in “b”.

Figure 12. Observed porosities in N-C1 microstructure; a) As polished (not etched) specimen to avoid impact from etchant b) Magnification of the yellow square in “a” c) 3d profile of the pores in “b”.

Fractography of Fatigue Samples

Fracture surfaces of fatigue specimens, both uncharged and hydrogen-charged with varying hydrogen content, were analyzed using SEM. In the study of fatigue failure surfaces, three primary features are consistently observed: the origin, the fatigue zone, and the overload zone. The origin is the point where the crack initiates. A single crack origin typically suggests a failure under low overstress conditions, whereas multiple origins might indicate high stress or significant stress concentrations. The crack then propagates slowly through the fatigue zone, where progression marks (or beach marks) are usually observed. Eventually, the crack reaches a critical size, and the remaining material can no longer sustain the applied stress, leading to the overload zone. In this zone, most cracks exhibit macroscopically brittle fracture characteristics.

The fractography of N-C1 fatigue samples tested at 50% of YS is provided in Error! Reference source not found.a-h. In these figures, the overload zone is marked with a dashed yellow line, crack origins are marked with white arrows, and ratchet marks are highlighted with transparent blue lines. The green square indicates the area for which the magnified image is provided in the adjacent figure.

Figure 13. SEM fracture surfaces of N-C1 fatigue samples tested at 50% of YS; a) CH=0 (Uncharged) b) Magnification of green area in “a” c) CH=0.8 wppm d) Magnification of green area in “c” e) Magnification of green area in “d” f) CH=2.0 wppm g) Magnification of green area in “f” h) Magnification of green area in “g”.

Figure 13. SEM fracture surfaces of N-C1 fatigue samples tested at 50% of YS; a) CH=0 (Uncharged) b) Magnification of green area in “a” c) CH=0.8 wppm d) Magnification of green area in “c” e) Magnification of green area in “d” f) CH=2.0 wppm g) Magnification of green area in “f” h) Magnification of green area in “g”.

Error! Reference source not found.b shows the overload zone for the uncharged (CH=0) sample, featuring dimples characteristic of ductile fracture. The morphology of this zone is similar in the two hydrogen-charged samples. However, its location shifts to the center of the fracture surface in charged samples, indicative of higher stresses [68]. As similar bending stress is applied to all samples during the test, this increase in stress can be explained by the HELP mechanism. According to this theory, hydrogen presence around dislocations leads to a local decrease in yield stress, facilitating local dislocation movement [25]. This effectively increases the ratio of imposed stress to local yield stress.

Additionally, Error! Reference source not found.c&f show an increase in the number of crack origins and ratchet marks with higher hydrogen content. The sample with 2 wppm hydrogen content exhibits a significantly higher number of origins. This increase in crack initiation sites is attributed to hydrogen-induced defects and microcracks, indicative of HE.

The fatigue zone also shows a quasi-cleavage surface in the sample with 0.8 wppm hydrogen content, which transitions to a cleavage surface in the sample with 2 wppm hydrogen content, reflecting the reduced ductility caused by HE, as illustrated in Error! Reference source not found.d and Error! Reference source not found.g. Moreover, there is an increase in the number and density of secondary cracks, as hydrogen promotes crack branching and secondary crack formation. This is evident from Error! Reference source not found.e and Error! Reference source not found.h, where higher hydrogen content correlates with larger secondary cracks.

Figure 14. SEM fracture surfaces of AO-C1 fatigue samples (base metal) tested at 80% of YS; a) CH=0 (Uncharged) b) Magnification of green area in “a” c) CH=1.0 wppm d) Magnification of green area in “c” e) Magnification of green area in “d”.

Figure 14. SEM fracture surfaces of AO-C1 fatigue samples (base metal) tested at 80% of YS; a) CH=0 (Uncharged) b) Magnification of green area in “a” c) CH=1.0 wppm d) Magnification of green area in “c” e) Magnification of green area in “d”.

Similar features were observed in the fracture surfaces of AO-C1 fatigue samples and N-C1 fatigue samples tested at 80% YS. Error! Reference source not found.a-h shows the SEM images for uncharged and charged samples (with 1.0 wppm hydrogen content) from AO-C1 steel. The primary difference between these images with those of N-C1 steel lies in the morphology of the overload zone, which exhibits more ductile features in AO-C1 samples. This is attributed to the fact that AO-C1 steel is inherently more ductile than N-C1 steel.

Conclusion

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