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Article

Synergistic Effects of Hydrostatic Pressure and Dissolved Oxygen on the SCC Behavior of Hydrogenated Ti6Al4V Alloy in Deep-Sea Environment

by
Feifei Huang
1,
Yuxiang Zhu
2,
Meng Yu
1,
Lei Wen
1,3,* and
Ying Jin
1,3,*
1
National Center for Materials Service Safety, University of Science and Technology Beijing, Beijing 100083, China
2
Central China Company of National Petroleum and Natural Gas Pipeline Network Group, Wuhan 430000, China
3
Innovation Group of Marine Engineering Materials and Corrosion Control, Southern Marine Science and Engineering Guangdong Laboratory, Zhuhai 519080, China
*
Authors to whom correspondence should be addressed.
Metals 2023, 13(3), 449; https://doi.org/10.3390/met13030449
Submission received: 1 February 2023 / Revised: 16 February 2023 / Accepted: 17 February 2023 / Published: 21 February 2023
(This article belongs to the Section Corrosion and Protection)
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Abstract

:
This work focused on the synergistic effect of hydrostatic pressure (HP) within the range of 0.1~50 MPa and a dissolved oxygen (DO) concentration within the range of 0.18~11.8 ppm on the stress corrosion cracking (SCC) behavior of hydrogenated Ti6Al4V alloy in a simulated deep-sea environment by electrochemical measurements and slow strain rate tensile (SSRT) tests. The potentiodynamic polarization and electrochemical impedance spectra results confirmed the corrosion resistance degradation with the HP increasing to 50 MPa. The fracture morphologies showed a mixed characteristic of brittle fracture on the surface layer and ductile fracture in the inner part. Higher HPs increased SCC susceptibility while a larger DO concentration decrease that of Ti6Al4V alloy.

1. Introduction

With the growing global requirements for energy and mineral resources, the exploration and development of the deep sea have attracted broad attention, since there are more abundant mineral and energy resources in deep-sea regions [1]. Compared with the shallow sea, deep-sea regions have still been less explored due to the extreme environmental conditions. The deep-sea environment exhibits the characteristics of high hydrostatic pressure (HP), low temperatures, less sunlight, high salinity and appreciable dissolved oxygen (DO), causing the unique corrosion behavior of metals serving in these extreme environments [2]. Among these deep-sea environmental factors, HP is currently the one most considered to influence the corrosion behavior of metals being used in ocean environments [3]. Previous studies have shown that the corrosion behavior of some metal materials under deep-sea HP is obviously different from that under atmospheric pressure. Baccaria et al. [4,5,6,7,8] investigated the corrosion behavior of passive metals (Al, Ni, Cu and stainless steels) under different HPs; the results showed that the corrosion resistance decreased with the increase in HP due to the enhanced effect on localized corrosion. Zhang et al. [9,10] concluded that HP promoted pitting initiation and growth in 316L and Fe–20Cr steels. Studies on active metals have also been conducted. HP treatment also deteriorated the stress corrosion cracking susceptibility [11,12] and general corrosion performance [13] of X70 steel. Similarly, HP promoted the corrosion rate of X65 steel by accelerating the cathodic reaction rate, as reported by Li et al. [14]. However, seldom have studies focused on the influence of DO or the synergistic effects of HP and DO on the corrosion behavior of metals in the deep-sea environment.
Since the traditional engineering material of steel is susceptible in marine environments, especially in deep-sea conditions, titanium and its alloys are good alternatives to steel in marine engineering due to their low density, specific strength, excellent corrosion resistance and high-temperature resistance, and have great potential to be promising structural materials in deep-sea exploration. The excellent corrosion resistance is attributed to the formation of dense passive films on the alloy surface [15]. The passive films act as a barrier to prevent the alloys from dissolution and block out aggressive ions in the solution [16,17]. However, the corrosion resistance of passive films formed on titanium alloys can be deteriorated in the deep-sea environment. Dong [18] found that the corrosion resistance of a passive film formed in simulated deep-sea environments is lower than that formed in a simulated shallow-sea environment, since the reduction reaction was restrained under the effect of high HP and intermediate oxides with a non-stoichiometric ratio, which was involved in the film’s formation due to the low DO concentration. Liu [19] studied the influence of HP on the passive films formed on the surface of Ti6Al4V alloy and concluded that HP accelerated the dissolution of the passive film.
In addition, titanium and its alloys are susceptible to stress corrosion cracking (SCC), which is relevant to hydrogen embrittlement (HE) in hydrogen-containing environments [20,21]. Hydrogen diffuses into the matrix of titanium and its alloys, and reacts with titanium to form hydride [22,23]. At the sites of hydride formation, stress concentration occurs due to the volume expansion caused by the hydrides’ formation, triggering the initiation of HE [24,25]. Components made of titanium/ titanium alloy may be attacked by hydrogen atoms generated from galvanic corrosion between titanium-made components and carbon steel or other metals during service in the marine environment [26]. In this case, hydrogen absorption and subsequently HE will occur in titanium alloys. Previous studies mainly focused on the HE behavior in the shallow-sea environment. Only a few studies focused on the HE behavior of titanium alloy in the deep-sea environment. Liu demonstrated that high HP promoted the formation of hydrides at the α/β interface of Ti6Al4V, and thus induced HE in a simulated deep-sea environment [27]. However, the process and the mechanisms of HE for titanium alloys serving in deep-sea environments are still unclear, and the influence of HP and DO on the failure behavior of titanium alloys needs to be clarified.
In this work, electrochemical hydrogen-charged Ti6Al4V alloy was employed to study the SCC behavior in the simulated deep-sea environment. X-ray diffraction (XRD) tests were used to confirm the phase composition of the charged and uncharged samples. The effects of HP and the DO concentration on the corrosion behavior and SCC susceptibility were investigated by electrochemical measurements including potentiodynamic polarization (PDP) and electrochemical impedance spectra (EIS) tests and slow strain rate testing (SSRT), respectively. The morphologies of the tensile specimens after SSRT at different HP and DO were observed by scanning electron microscopy (SEM).

2. Materials and Methods

2.1. Experimental Materials

A commercial hot-rolled Ti6Al4V alloy with the chemical composition (wt.%) of Al 6.30%, V 4.20%, Fe 0.07%, C 0.02%, and remainder Ti was investigated in this study. The metallograph of the sample is shown in Figure 1, which shows a bimodal microstructure of equiaxed α phase and lamellar β phases.

2.2. Cathode Hydrogen Charging

The hydrogen charging samples were cut to a size of 10 mm × 10 mm × 1 mm. Epoxy resin and ethylenediamine mixed in a ratio of 10 to 1 were prepared for the sealing of the samples, leaving a 10 mm × 10 mm test area. The samples’ surfaces were ground with silicon papers step by step to 2000 grit. The hydrogen-charging experiment was performed in 3.5 wt.% NaCl solution prepared with analytical-grade NaCl and deionized water. A standard three-electrode system was adopted, with the sample as the working electrode, a saturated calomel electrode (SCE) as the reference electrode, and a platinum plate as the counter electrode. All the samples were charged for 1 h utilizing a PARSTAT 4000 electrochemical workstation (AMETEK, San Diego, CA, USA) with a constant current density of 100 mA/cm2.

2.3. Microstructural Characterization

The phase composition of the charged and uncharged samples was determined by using the X-ray diffraction (XRD) method. The XRD experiments were carried out by using a Rigaku SmartLab 9 kw X-ray diffractometer (Rigaku, Tokyo, Japan) with Cu Kα radiation at 20 kV and 45 mA at a step size of 0.02° and a scan rate of 4°/min. The surface and cross-section morphologies of the charged samples were observed by a scanning electron microscope (SEM, ZEISS MERLIN Compact, Jena, Germany).

2.4. Electrochemical Corrosion Tests

All electrochemical corrosion tests were carried out in deep-sea environment simulation equipment, in which the HP, DO, temperature, salinity, pH value and mechanical load could be monitored and controlled independently. The electrochemical tests were performed by using an electrochemical workstation (PARSTAT 4000) in a conventional three-electrode cell, with the specimen as the working electrode, a platinum plate as the counter electrode, and an Ag/AgCl electrode as the reference electrode. The potentiodynamic polarization curves were measured at a scan rate of 0.5 mV/s. Electrochemical impedance spectroscopy (EIS) was conducted with a sinusoidal potential perturbation of 10 mV and a frequency range from 105 to 10−2 Hz in 3.5 wt.% NaCl solution. All electrochemical experiments were performed at pressures of 0.1 MPa, 10 MPa, 30 MPa or 50 MPa, while keeping all other environmental factors unchanged (temperature 5 °C, DO 3.8 ppm).

2.5. SSRT Tests

The SSRT samples were prepared with a gauge length of 13 mm using Wire-Cut Electrical Discharge Machining (WEDM). After the WEDM process, the samples were hydrogen-charged. SSRT tests were conducted by the loading device implemented in the deep-sea environment simulation equipment with a strain rate of 2.5 × 10−5 s−1. To determine the effect of HP on the SCC behavior of charged samples, SSRT experiments were performed under HP of 0.1 MPa, 10 MPa, 30 MPa and 50 MPa, keeping all other environmental factors unchanged (temperature 5 °C, DO concentration 3.8 ppm). Meanwhile, to investigate the effect of DO, SSRT tests were carried out at a DO concentration of 0.18, 3.8, 7.8 and 11.8 ppm, with all other environmental factors unchanged (temperature 5 °C, HP 30 MPa).

3. Results

3.1. Microstructure Characterization

Figure 2 shows the XRD patterns of Ti6Al4V samples before and after hydrogen charging for 1 h. In addition to the α-Ti and β-Ti phase peaks found in the uncharged sample, there are other peaks, which can be attributed to the formation of the hydride phase, in the XRD pattern of the hydrogen-charged sample. The XRD profiles in Figure 1 imply that the TiH and TiH2 hydrides are the dominant phase formed by 1 h hydrogenation at room temperature. As is well known, TiH with an FCT structure and TiH2 with an FCC structure are stable hydride phases at normal pressure and room temperature [23,28].
The microscopic surface morphologies of Ti6Al4V samples before and after hydrogen charging at a current density of 100 mA/cm2 for 1 h are shown in Figure 3. It can be observed that the uncharged surface was smooth and flat, with an equiaxed α phase and lamellar β phase, as shown in Figure 3a. After hydrogen charging for 1 h, some microcracks relevant to hydrogen-induced cracking (HIC), which are marked with white arrows, originated from the interface between the α phase and β phase, as shown in Figure 3b. According to previous research results [22,29], HIC cracks usually occur at the α phase or/and along the α/β interface during cathodic charging, caused by the formation and fracture of the brittle hydride phases.
The cross-section morphologies of charged and uncharged Ti6Al4V samples are shown in Figure 4. In contrast to the uncharged sample illustrated in Figure 4a, a surface hydride layer appeared and the distribution of hydrides along the depth direction of the charged sample can also be observed in Figure 4b. More needle-like hydrides were generated near the surface area of the charged sample compared with other parts of the sample. The hydride content gradually decreased with the depth increasing. In addition, it can be observed that hydrides formed more easily in the α phase and along the α/β interface. This result is consistent with other research work focusing on the electrochemically hydrogen-charged Ti6Al4V alloy [30]. The α phase and the boundaries of the α/β phase were effective trapping sites for hydrogen [31], while the β phase acted as a short path for hydrogen transportation [32]. The formation of hydrides in α + β titanium alloys is related to the difference in the diffusion rate and solubility of hydrogen in the α and β phases. The diffusion rate of hydrogen in the β phase is higher than that in the α phase [33]. Meanwhile, the solubility of hydrogen in the α phase is lower than that in the β phase. When hydrogen diffuses to the surface of α + β titanium alloy, it first enters the β phase. With the progression of diffusion, the β phase will become a hydrogen-rich area and hydrogen source for hydride formation. Hydrogen diffuses from the hydrogen-rich β phase to the hydrogen-poor α phase until it exceeds the critical hydrogen concentration of the α phase. At this stage, hydrides are formed along the interface between the α and β phases [34]. As hydrogen diffusion proceeds, hydrides form along α grain boundaries, which contain more hydrogen traps. Eventually, the hydrides grow into the interior of the α grain. When hydrides are formed within α grains, flake hydrides will form in directions related to the orientation of α grains.

3.2. Electrochemical Behavior of Hydrogenated Ti6Al4V Alloy

3.2.1. Influence of Hydrogen Charging on the Corrosion Behavior of Ti6Al4V Alloy

The potentiodynamic polarization results of the charged and uncharged sample in 3.5 wt.% NaCl solution are illustrated in Figure 5. Compared with the uncharged sample, the curve of the charged one shifts to the side with higher current density, indicating that the electrochemical process, especially the anodic process, is accelerated by the formation of hydrides. A much narrower passivation region can be observed in the anodic curve of the charged sample compared to that of the uncharged one, which indicates an imperfect passive state in the charged sample. In addition, the passivity current density of the charged sample is much larger than that of the uncharged one, suggesting that the protection effect of the passive film is undermined by the hydrogen-charging process.
Figure 6 shows the Nyquist plots for the charged and uncharged samples. Both plots exhibit an uncompleted single capacitive arc, implying the highly capacitive behavior of the samples. The passive film formed on the titanium alloys determines the capacitive behavior. The radius of the capacitive loop decreases with hydrogen charging, indicating the worse corrosion resistance of the charged sample. By using the equivalent circuit (EC) embedded in Figure 6, the experimental results can be well fitted.

3.2.2. Influence of HP on the Corrosion Behavior of Charged Sample

The potentiodynamic polarization curves of the charged sample under different hydrostatic pressures are shown in Figure 7. With the increase in hydrostatic pressure, the curve gradually moves to the high-current-density side, i.e., the right side, indicating that the electrochemical reaction is accelerated by the high hydrostatic pressure. Moreover, the passive region can be recognized from all the curves, but the potential range becomes narrower with the increasing hydrostatic pressure, which implies that the hydrostatic pressure plays a role in the protection effect of the passive film.
Figure 8 shows the Nyquist plots for the sample charged for 1 h under different HPs (0.1 10, 30 and 50 MPa). It can be observed that all the curves show an uncompleted single capacitive arc, which indicates the highly capacitive behavior of the sample. This capacitive behavior of the charged titanium alloys corresponds to the electrochemical process of the passive film formed on the sample surface. The radius of the capacitive loop decreases with increasing pressure, which indicates that pressure promotes the process of metal dissolution. The experimental results can be well fitted by the EC in Figure 6.

3.3. SCC Behavior of Hydrogenated Ti6Al4V Alloy

3.3.1. Influence of HP on the SCC Behaivor of Hydrogenated Ti6Al4V Alloy

The stress–strain curves of Ti6Al4V alloy under the hydrostatic pressures of 0.1, 10, 30 and 50 MPa are shown in Figure 9. As seen in the figure, the elongation and the tensile strength of the specimens under high HPs are less than those under 0.1 MPa. The specimen at 50 MPa exhibits the smallest average values for a fracture strain of 18.2% due to the more aggressive environment.
The fracture morphologies of specimens after SSRT were examined by SEM to reveal the effect of hydrostatic pressure on the fracture behavior, and the fracture edge morphologies of the 1 h charged samples under different HPs are shown in Figure 10. Secondary cracks with the length of several hundreds of microns can be observed. It can be observed from Figure 10a,c that the fracture side surface under 50 MPa exhibits a larger area of secondary cracks than that of the sample under 0.1 MPa, indicating that the ductility of the sample decreases with the increased HP.
Figure 11 shows the highly magnified fracture surfaces of the samples under different HPs. A mixed dimple morphology and quasi-cleavage characteristics are recognized, which proves that the brittle and ductile fracture occur simultaneously. According to the different characteristics, the fracture surface of the charged samples could be divided into two parts. The region near the hydrogen-charging surface shows the typical brittle feature, while the region away from the charging surface presents ductile characteristics. The region near the hydrogen-charging surface shows tearing dimples (0.1 MPa) and a quasi-cleavage feature (10, 30 and 50 MPa). This region is caused by the brittle cracking of the hydride layer. The regions away from the hydrogen-charging surface under different HPs mainly consist of dimples, indicating that this region is derived from the ductile fracture of the unaffected matrix. In addition, some obvious cracks were located on the fracture surface for the specimen that failed at 50 MPa, as shown in Figure 11d. It can be noticed that the area of brittle fracture increased with the increasing HPs.

3.3.2. Influence of DO on the SCC Behavior of Hydrogenated Ti6Al4V Alloy

The influence of dissolved oxygen content on the SCC behavior of hydrogenated Ti6Al4V is shown in Figure 12. It can be observed that the tensile strength and elongation of samples under low dissolved oxygen content (0.18 ppm and 3.8 ppm) are less than those under higher dissolved oxygen content (7.8 ppm and 11.8 ppm). The elongation of charged samples increases with the increasing DO concentration, indicating that the increase in DO concentration can effectively reduce the SCC susceptibility, as shown in Figure 13b.
The fracture surfaces under different DO concentrations were observed to analyze the deformation behavior of hydrogenated Ti6Al4V during SSRT, and the results are shown in Figure 13. As shown in Figure 13a, the fracture surface of specimens under lower DO concentrations (0.18 ppm) shows a brittle facture characteristic, while the fracture surface under higher DO concentrations (11.8 ppm) shows mixed fracture characteristics, mainly consisting of dimples and flat zones composed of quasi-cleavage fractures. The results indicate that the lower DO concentration enhanced the brittle fracture.

4. Discussion

During the hydrogen-charging process, hydrogen atoms first adsorbed on the surface of the passive film and reacted with titanium to form hydrides after diffusing into the substrate [30]. The formation of hydrides can be evidenced by Figure 2 and Figure 4. The deterioration of the corrosion resistance of the charged alloy can be attributed to two factors. On the one hand, the conversion of the passive film from TiO2 to TiOOH led to the production of the absorbed hydrogen in the passive TiO2 film [35], which is regarded as an electron donor to reduce the resistivity of the film [36]. On the other hand, hydride formation led to the occurrence of HIC on the surface of the alloy, as shown in Figure 3. Therefore, the corrosion resistance of Ti6Al4V decreased due to the synergistic effect of the degradation of the passive film and the occurrence of surface defects, resulting in the larger corrosion current density (Figure 5) and smaller capacitance arc (Figure 6).
In simulated deep-sea environments, the charged alloy was subjected to the synergistic influence of HP, DO, temperature and a corrosive medium (in this research, DO 3.8 ppm, 4 °C, 3.5% NaCl solution and HP varied from 0.1 MPa to 50 MPa). According to previous research, HP facilitated the anodic dissolution process by thinning the electric double layer in deep-sea environments [37]. HP also accelerated the dissolution rate of the alloy by promoting the adsorption of Cl− on the metal surface [38], and the activity of Cl− was enhanced as well [39]. The adsorption of Cl− in the HIC on the surface of the charged Ti6Al4V alloy could enhance the dissolution within the cracks. In addition, HP reduced the stability of the passive film formed on the Ti6Al4V alloy by decreasing the TiO2 content and facilitated the dissolution of the passive film and metal substrate [19]. Therefore, the charge transfer resistance of Ti6Al4V alloy decreased dramatically under the conditions of high HPs (shown in Figure 8), and the anodic current density increased by approximately one order under 50 MPa HP compared with that of 0.1 MPa (shown in Figure 7), indicating that the corrosion resistance of the charged sample was reduced. As a result, the localized corrosion process in the defect region and the global dissolution of the passive film and metal substrate were promoted in the deep-sea environment.
When applying SSRT on the charged Ti6Al4V alloy under the simulated deep-sea environment, the stress caused microstructure evolution, such as dislocations and microcrack propagation. Tensile stress changed the microstructure of the metal surface under the passive film, the DO concentration influenced the composition and stability of the passive film at the solid side of the metal/solution interface, and HP facilitated the dissolution of the passive film and compressed the double layer at the liquid side of the interface [3]. In addition, the presence of HIC would accelerate the deterioration process of the metal. Therefore, the stress, DO, HP and HIC defect played a synergistic role in the corrosion and SCC behavior of the charged Ti6Al4V alloy under the simulated deep-sea environment.
According to the fracture morphologies under different HPs (Figure 10 and Figure 11), brittle areas and ductile areas could be found for all the tested samples. With the increasing HP, the brittle area with a quasi-cleavage characteristic increased, and the ductile area featured by dimples decreased, indicating that HP increased the SCC susceptibility of the charged Ti6Al4V alloy, which is consistent with the tensile results in Figure 9. According to a previous study, the cleavage and quasi-cleavage features resulted from the crack propagating through the hydride [22]. HP accelerated the anodic dissolution process, resulting in the breakdown of the passive film [27]. Faster anodic reactions acidified the solution at the crack tips to a lower pH value, resulting in the production of more hydrogen under higher HP. More hydrogen atoms would be absorbed at the crack tip to form more hydrides under high HPs due to the lower protective ability of the passive film under high HPs. Therefore, a larger brittle area appeared on the facture surface under higher HPs (Figure 11). However, the SCC susceptibility of the charged Ti6Al4V alloy decreased slightly with the increasing DO concentration, as shown in Figure 12. The higher DO concentration facilitated the formation of a passive oxide film not only on the alloy surface but also at the crack tip area, resulting in the formation of less hydrides. Then, the formation of hydrides was reduced to some extent under a higher DO concentration. As a result, the brittle fracture was promoted by the low DO concentration (Figure 13).

5. Conclusions

This work sought to clarify the synergistic effect of HP and the dissolved oxygen concentration on the SCC behavior of the hydrogenated Ti6Al4V alloy in a simulated deep-sea environment, and led to the following conclusions:
  • Needle-like hydrides, mainly TiH and TiH2, and hydrogen-induced cracks formed on the surface of the Ti6Al4V alloy after electrochemical charging for 1 h. The anodic current density and capacitance arc increased and shrunk, respectively, after hydrogen charging.
  • The SCC of the hydrogen-charged Ti6Al4V alloy in the simulated deep-sea environment was hydrogen embrittlement. The fractography indicated a mixed fracture that was both brittle and ductile. The higher the HPs, the larger the brittle area.
  • The SCC susceptibility of the charged Ti6Al4V alloy was enhanced by the higher HPs and/or the lower DO concentrations.

Author Contributions

Conceptualization, F.H. and Y.Z.; methodology, M.Y.; software, F.H.; validation, F.H., Y.Z. and M.Y.; formal analysis, F.H., L.W. and M.Y.; investigation, Y.Z.; resources, L.W.; data curation, Y.Z. and M.Y.; writing—original draft preparation, F.H., Y.Z. and M.Y.; writing—review and editing, F.H., L.W. and Y.J.; visualization, F.H. and L.W.; supervision, L.W. and Y.J.; project administration, L.W. and Y.J.; funding acquisition, L.W. and Y.J. All authors have read and agreed to the published version of the manuscript.

Funding

This work was supported by the National Key Research and Development Program of China (grant no. 2022YFA1603803).

Data Availability Statement

Not applicable.

Conflicts of Interest

The authors declare no conflict of interest.

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Figure 1. Metallograph of Ti6Al4V alloy.
Figure 1. Metallograph of Ti6Al4V alloy.
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Figure 2. XRD patterns of Ti6Al4V alloy before and after hydrogen charging for 1 h in 3.5 wt.% NaCl solution.
Figure 2. XRD patterns of Ti6Al4V alloy before and after hydrogen charging for 1 h in 3.5 wt.% NaCl solution.
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Figure 3. Surface morphologies of Ti6Al4V alloy before (a) and after (b) 1 h hydrogen charging in 3.5 wt.% NaCl solution.
Figure 3. Surface morphologies of Ti6Al4V alloy before (a) and after (b) 1 h hydrogen charging in 3.5 wt.% NaCl solution.
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Figure 4. Cross-section morphologies of Ti6Al4V alloy before (a) and after (b) 1 h hydrogen charging in 3.5 wt.% NaCl solution.
Figure 4. Cross-section morphologies of Ti6Al4V alloy before (a) and after (b) 1 h hydrogen charging in 3.5 wt.% NaCl solution.
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Figure 5. Potentiodynamic polarization curves of TC4 alloy before and after hydrogen charging for 1 h in 3.5 wt.% NaCl solution.
Figure 5. Potentiodynamic polarization curves of TC4 alloy before and after hydrogen charging for 1 h in 3.5 wt.% NaCl solution.
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Figure 6. EIS results and the corresponding EC of TC4 alloy before and after hydrogen charging for 1 h in 3.5 wt.% NaCl solution.
Figure 6. EIS results and the corresponding EC of TC4 alloy before and after hydrogen charging for 1 h in 3.5 wt.% NaCl solution.
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Figure 7. The polarization curves of TC4 alloy after 1 h charging in 3.5% NaCl solution at different HPs.
Figure 7. The polarization curves of TC4 alloy after 1 h charging in 3.5% NaCl solution at different HPs.
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Figure 8. The EIS results for charged Ti6Al4V alloy under different HPs in 3.5 wt.% NaCl solution.
Figure 8. The EIS results for charged Ti6Al4V alloy under different HPs in 3.5 wt.% NaCl solution.
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Figure 9. The stress–strain curves (a) and elongation (b) of Ti6Al4V alloy after 1 h charging under different HPs.
Figure 9. The stress–strain curves (a) and elongation (b) of Ti6Al4V alloy after 1 h charging under different HPs.
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Figure 10. Fracture edge morphologies of Ti6Al4V alloy under different hydrostatic pressures: (a,b) 0.1 MPa; (c,d) 50 MPa.
Figure 10. Fracture edge morphologies of Ti6Al4V alloy under different hydrostatic pressures: (a,b) 0.1 MPa; (c,d) 50 MPa.
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Figure 11. Fracture surface morphologies of 1 h charged Ti6Al4V alloy after SSRT under different HPs: (a) 0.1 MPa; (b) 10 MPa; (c) 30 MPa; (d) 50 MPa.
Figure 11. Fracture surface morphologies of 1 h charged Ti6Al4V alloy after SSRT under different HPs: (a) 0.1 MPa; (b) 10 MPa; (c) 30 MPa; (d) 50 MPa.
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Figure 12. The stress–strain curves (a) and elongation (b) of Ti6Al4V alloy after 1 h charging under different DO concentrations.
Figure 12. The stress–strain curves (a) and elongation (b) of Ti6Al4V alloy after 1 h charging under different DO concentrations.
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Figure 13. Fracture surface morphologies of 1 h charged Ti6Al4V alloy after SSRT under different DO concentrations: (a) 0.18 ppm; (b) 11.8 ppm.
Figure 13. Fracture surface morphologies of 1 h charged Ti6Al4V alloy after SSRT under different DO concentrations: (a) 0.18 ppm; (b) 11.8 ppm.
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MDPI and ACS Style

Huang, F.; Zhu, Y.; Yu, M.; Wen, L.; Jin, Y. Synergistic Effects of Hydrostatic Pressure and Dissolved Oxygen on the SCC Behavior of Hydrogenated Ti6Al4V Alloy in Deep-Sea Environment. Metals 2023, 13, 449. https://doi.org/10.3390/met13030449

AMA Style

Huang F, Zhu Y, Yu M, Wen L, Jin Y. Synergistic Effects of Hydrostatic Pressure and Dissolved Oxygen on the SCC Behavior of Hydrogenated Ti6Al4V Alloy in Deep-Sea Environment. Metals. 2023; 13(3):449. https://doi.org/10.3390/met13030449

Chicago/Turabian Style

Huang, Feifei, Yuxiang Zhu, Meng Yu, Lei Wen, and Ying Jin. 2023. "Synergistic Effects of Hydrostatic Pressure and Dissolved Oxygen on the SCC Behavior of Hydrogenated Ti6Al4V Alloy in Deep-Sea Environment" Metals 13, no. 3: 449. https://doi.org/10.3390/met13030449

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