Effect of HSO 3 − and alternating current on corrosion behaviour and mechanism of CoCrFeNi HEA in a simulated marine environment

Abstract

The corrosion behavior and mechanism of CoCrFeNi HEA under AC application were systematically studied in 3.5 wt.% NaCl solution containing different concentrations of NaHSO₃. The results showed that the existence of HSO₃ ⁻ increased the defect density within the passivation film on the surface of the HEA. With the increase of HSO₃ ⁻ concentration, the passivation zone became narrow, E_p shifted negatively, and i_p increased, which indicated that the corrosion resistance of the HEA was weakened. When the HSO₃ ⁻ concentration increased to 0.1 M, a clear transition zone appeared, indicating that the increase in concentration reduced the passivity of the HEA, and its corrosion behavior changed from the typical passivation characteristic to the activation state with increasing i_AC. The imposed AC promoted the nucleation of metastable pitting, which was indicated by a remarkable increase in the average pitting number density (N_avg) from 5 to 17 mm⁻² with the enhancement of i_AC.

GRAPHICAL ABSTRACT

Keywords

CoCrFeNi HEA alternating current corrosion behavior passive film corrosion behaviour

Highlights

Corrosion behaviour of CoCrFeNi HEA with imposed AC was studied in a SO₂-polluted marine environment.

Metastable pitting behaviour was investigated in NaCl solution containing at different i_AC.

With increased concentration, the stability and protective ability of passive film were weakened.

The HEA changed from passivation characteristic to activation state with the increase of applied i_AC.

The initiation of pitting was due to the selective dissolution of Ni element, and the selective dissolution of Cr, Fe, Co and Ni elements occurred inside the oxide.

Introduction

Traditional alloys are generally designed based on one main metal component plus a small number of other metal elements to improve or obtain the desired properties, which can not satisfy the application requirements of some complex industrial environments [1,2]. In addition, the drawback of this design method is that some complex compounds or unpredictable phase compositions can be generated. Nowadays, Ye et al. [3] proposed a new concept of alloy design, that is, the alloy containing multiple primary elements has simple structures with fewer phase compositions. These alloys, known as entropy-based alloys, have attracted a lot of attention from researchers because of their superior mechanical properties [4], thermal stability [5] and wear resistance [6,7]. Among them, the equiatomic CoCrFeNi high-entropy alloy (HEA) is one of the most representative HEAs and has achieved some research results. Salishchev et al. [8] investigated the metallographic structure and mechanical properties of CoCrFeNi alloy under solidification and annealing conditions. The results showed that the HEA had a single-phase face-centered cubic (FCC) structure, and the addition of the V element caused significant strengthening accompanied by plasticity deficiency. Sathiyamoorthi et al. [9] compared the change in thermal stability of the CoCrFeNi HEA under milling and sintering conditions, demonstrating its excellent high-temperature properties. Huo et al. [10] showed the advantage of CoCrFeNi HEA in replacing conventional alloys in critical loading applications, and the results demonstrated that the HEA exhibited a good combination of high strength and plasticity under tensile deformation at high strain rates.

Based on the outstanding mechanical performances of CoCrFeNi HEA, it can be used in practical engineering structures. In engineering applications, the anti-corrosion property of materials is also an important index to evaluate the service life of equipments, and the current studies on the corrosion behaviour of CoCrFeNi HEA are very limited [11,12]. Liu et al. [13] reported that the addition of Ge element inhibited pitting nucleation and enhanced anti-corrosive performance. Hsu et al. [14] investigated that FeCoNiCr HEA with FCC structure was easy to passivate in NaCl solution, and even exhibited better anti-corrosion performance than 304L stainless steel. Zhu et al. [15] explained the effect of microstructure and chemical composition of CoCrFeNi HEA on the electrochemical properties in NaCl solution and proposed that the high content of Cr in the HEA promoted the formation of the passivation layer. Wu et al. [16] explored that a non-equiatomic FeNiCoCr possessed higher corrosion resistance than 316L SS in 0.1 M H₂SO₄ solution. The as-homogenised HEA possessed a very low passivation current density (i_p) of 4.7224 × 10⁻⁶ A/cm².

In recent years, researchers have extensively studied the corrosion behaviour of various alloys in the simulated marine environment [17,18], and the results show that Cl⁻ can adsorb in the weak position of the passivating film on the surface of the materials, destroying the stability and integrity of the film and inducing pitting nucleation. With the extensive development of marine resources, in some regions, due to industrial pollution, the marine atmosphere contains a certain amount of SO₂ and SO₂ dissolves in seawater to form , so the alloys applied in the offshore platforms and marine construction in the splashing environment of seawater are susceptible to more serious corrosion [19,20]. On the one hand, can also be adsorbed in the active sites of the passivation film and combined with Cl⁻ to accelerate the corrosion process; on the other hand, the ionisation effect of is greater than the hydrolysis effect, producing more H⁺ and increasing the local acidity, which is destructive to the metals [21]. However, there is no research report on the corrosion behaviour of CoCrFeNi HEA in SO₂-pollution seawater solution. Therefore, the study of the corrosion behaviour of the HEA in the industrial pollution environment is an important guiding value for its application in engineering. Moreover, some coastal areas and offshore platforms are disturbed by alternating current (AC) due to the surrounding high-voltage transmission lines or electric equipments [22,23], which may promote the corrosion of CoCrFeNi HEA employed in these cases. Thus, it is necessary to study the corrosion behaviour of the HEA with AC interference in NaCl solution polluted by various levels of SO₂.

In the current paper, the effect of and AC on the passivation film and corrosion behaviour of CoCrFeNi alloy in a simulated marine environment was investigated by electrochemical tests, surface analysis techniques and immersion corrosion test. In addition, the effect of AC on the initiation and development of metastable pitting was statistically analysed.

Experimental

Materials and test solution

The material used for the tests was an equiatomic CoCrFeNi HEA ingot prepared from elements with purity above 99.9% wt.% by vacuum arc melting. It was machined into 10 × 10 × 3 mm specimens for the study of the electrochemical behaviour, and the remaining surfaces were sealed with epoxy resin except for a 10 × 10 mm² surface. After sanding from 180 to 2000 grit using SiC sandpaper, they were cleaned sequentially with anhydrous ethanol and deionised water in turn, dried and set aside. However, for the study of metastable pitting, the specimen was filamentary with a cross-sectional area of 0.1 mm² to avoid the inability to induce metastable pitting due to the current superposition.

In this work, the test solution was 3.5 wt.% NaCl solution with different concentrations of NaHSO₃ for simulating SO₂-polluted marine environment. All measurements were carried out at a constant temperature of 30°C.

Microstructure analysis

The CoCrFeNi specimens were etched with the reagent solution of aqua regia (HCl:HNO₃ = 3:1) for 10 s. Afterward, the microstructure of the HEA was characterised using scanning electron microscopy (SEM) equipped with energy dispersive spectroscopy (EDS).

Electrochemical tests

The electrochemical workstation (CHI660E) was employed to perform electrochemical tests using the classical three-electrode system: the saturated calomel electrode (SCE) as the reference electrode (RE), the platinum sheet as the counter electrode (CE) and the HEA sample as the working electrode (WE). Before the potentiodynamic polarisation curves tests, to maintain a steady state, the samples were immersed in 3.5 wt.% NaCl solution with various concentrations of NaHSO₃ (0, 0.01, 0.05, 0.1 M) for 1 h, and then the open-circuit potential (OCP) was recorded for 30 min. The potential range of the polarisation curve was scanned from −1 V to 0.5 V (vs SCE) and the scanning rate was 1 mV/s.

The Mott–Schottky test was performed by polarising the sample at a constant potential of 0 V for 1.5 h to form a stable passivating film on the HEA surface, and then soaking it in the solution until the OCP was stabilised. The test was carried out at a frequency of 1 kHz, with a potential range of −0.4∼0.8 V (vs. SCE) and a step voltage of 50 mV. Similarly, the electrochemical impedance spectroscopy (EIS) of the surface film was measured at a sinusoidal potential amplitude of 10 mV in the frequency range of 100 kHz ∼10 mHz. The impedance data were processed by ZSimpWin software to obtain the corresponding electrochemical parameters.

To investigate the effect of applied AC on passive film covering the HEA, the potentiodynamic polarisation curves were conducted under various alternating current densities (i_AC) of 0, 30, 50 and 100 A/m² in 3.5 wt.% NaCl solution with 0.01 M NaHSO₃. The arrangement of the electrochemical experiment under AC interference was consistent with our previous studies [24,25]. The cyclic polarisation curve was scanned from −1 to the anodic direction with a scan rate of 0.5 mV/s. When the potential reached 0.4 V, it was scanned back to the initial potential. During the progress, the potential corresponding to the sudden increase of current density was the pitting potential (E_p). Then the potential at the first intersection of the reverse scan towards the cathodic direction and the forward scan was regarded as the protection potential (E_prot).

Metastable pitting analysis

To understand the dynamic change process of metastable pitting and stable pitting under AC interference from the perspective of kinetics, the potentiostatic polarisation tests were performed in 3.5 wt.% NaCl solution with 0.01 M NaHSO₃ for 1800 s. The metastable pitting potential was 0 V_SCE according to the cyclic polarisation curve. Since the metastable pitting current was small, the applied AC was also appropriately reduced to obtain the obvious experimental phenomenon. Therefore, the applied i_AC was 0, 3, 5 and 10 A/m².

Immersion test

The electrochemical specimens polished with 0.5 μm diamond paste were immersed in 3.5 wt.% NaCl solution with various concentrations of NaHSO₃ for 8 d. Moreover, the samples were weighed before and after the immersion test, and the corrosion rate was calculated by the weight loss method. The pitting morphologies were observed by SEM and the elemental changes were analysed by EDS.

Results and discussion

Microstructure analysis

Figure 1 displays the SEM morphology of CoCrFeNi HEA after etching by aqua regia. The alloy exhibits a single-phase FCC structure, which avoids the micro-galvanic corrosion effect between different phases [26,27]. Moreover, some fine gray-white particles are sporadically distributed on the alloy surface, considered inclusions. Among them, in the process of metallographic preparation, a part of the inclusions are dislodged from the matrix due to mechanical polishing, and pores are left behind [28]. The morphology and EDS results of the inclusion at high magnification are given in Figure 2. Based on EDS analysis, the inclusion contains higher contents of O, Al and lower contents of alloy elements (i.e. Cr, Fe, Co and Ni) compared to the matrix, and the corresponding EDS mapping confirms this in Figure 2(c). Also, it can be observed that the inclusion is approximately spherical and surrounded by a hole and micro-gap in Figure 2(a,c). As widely believed, the inclusion enriched in Al and O is brittle and hard, whose thermal expansion coefficient is different from the alloy matrix, which is prone to lattice distortion nearby the inclusion [29,30]. Consequently, micro-gaps and holes are easily formed at the interface between the inclusion and the HEA. Numerous studies have reported the effect of inclusions within microstructure on corrosion resistance [31,32]. Villavicencio et al [33] investigated that alumina inclusions had little effect on the overall corrosion rate but promoted the sprouting of pitting corrosion.

Figure 1

Microstructure of CoCrFeNi HEA.

Figure 2

Original morphology (a) and EDS results (b, c) of oxide in CoCrFeNi HEA.

In summary, the single-phase structure of CoCrFeNi HEA avoids the galvanic reaction between different phases and enhances the corrosion resistance, compared to the conventional alloy [27]. The presence of inclusions increases the local corrosion activity. Besides, in the early corrosion development stage, the micro-galvanic corrosion cell is also formed at the oxide/matrix interface, preferentially sprouting pitting corrosion around the oxides, so that the original micro-gaps gradually expand.

Electrochemical measurements

The evolution of the OCP as a function of time can reveal the thermodynamic corrosion trend of the HEA, reflecting the corrosion activity [34]. Figure 3 presents the OCP curves of CoCrFeNi HEA acquired from 3.5 wt.% NaCl solution with different NaHSO₃ concentrations over 1800 s. There is no significant fluctuation in the OCP values since the samples are pre-passivated for 1 h before testing. It can be seen that as the NaHSO₃ concentration increases, the corresponding OCP value shifts negatively. Therefore, the corrosion tendency of CoCrFeNi HEA enhances with the increase of NaHSO₃ concentration.

Figure 3

OCP curves of CoCrFeNi HEA in 3.5 wt.% NaCl solution with different concentrations of NaHSO₃.

Figure 4(a) exhibits the potentiodynamic polarisation curves of the HEA measured in 3.5 wt.% NaCl solution with various NaHSO₃ concentrations. The anodic branches of the curves are roughly similar, forming a typical passivation zone, which indicates that the dissolution and passivation rates of the surface passive film are in a dynamic balance state. On the other hand, the characteristics of the plots are significantly affected by the change in NaHSO₃ concentration. The passivation region gradually shifts to the right and becomes narrow with increasing NaHSO₃ concentration, so that the passivation current density (i_p) also gradually increases. In addition, the relevant electrochemical characteristic parameters are listed in Table 1. It can be visualised that the smallest i_p is observed for the specimen tested at 0 M NaHSO₃, indicating that the HEA possesses the lowest dissolution rate of the film and the best corrosion resistance. Moreover, the sample displays stable passivation without NaHSO₃, while it appears an active–passive transition phenomenon with the addition of NaHSO₃, especially in 0.1 M NaHSO₃. Thus, the critical passive current density (i_pp) increases with the increase of NaHSO₃ concentration, indicating that may hinder the passivation of the HEA [35]. As shown in Figure 4(b), the i_p value gradually increases with the increasing NaHSO₃ concentration, indicating that the corrosion resistance gradually becomes worse, while the pitting potential (E_p) has exactly the opposite change pattern. These results indicate that the sensitivity to pitting corrosion increases.

Figure 4

Potentiodynamic polarisation curves (a), i_p and E_p values (b) of the HEA in 3.5 wt.% NaCl solution with different concentrations of NaHSO₃.

Table 1

The fitted results for potentiodynamic polarisation curves of CoCrFeNi HEA in 3.5 wt.% NaCl solution with different concentrations of NaHSO₃.

NaHSO₃ concentrations	i_p (10⁻⁵ A·cm⁻²)	i_pp (10⁻⁵ A·cm⁻²)	E_P (V)
0 M	0.259	—	0.258
0.01 M	0.497	0.593	0.162
0.05 M	0.716	0.799	0.079
0.1 M	1.703	8.165	0.022

The Nyquist and Bode plots of the HEA measured with various NaHSO₃ concentrations are presented in Figure 5. The Nyquist curves with incomplete capacitive impedance arcs are in Figure 5(a). These curves progressively decrease in radii with increasing NaHSO₃ concentration, and the corresponding corrosion resistance becomes weaker [36]. The value of impedance modulus at 0.01 Hz in the Bode plot also reflects the change in the anti-corrosion property of the HEA [37]. In Figure 5(b), the |Z| value is maximum at 0.01 Hz in NaCl solution with 0 M NaHSO₃, which reveals the strongest corrosion resistance, while the |Z| value is minimum in 0.1 M NaHSO₃ solution. Also, it can be seen from Figure 5(c) that the maximum phase angles of the samples in different concentrations are between 70° and 80°, suggesting that a uniform and stable passivating film is formed on the samples [38], which is conducive to the formation of a barrier to prevent the penetration of corrosive ions.

Figure 5

EIS curves of the HEA at the film formation potential of 0 V_SCE in 3.5 wt.% NaCl solution with different concentrations of NaHSO₃: (a) Nyquist plots; (b, c) Bode plots.

To determine the electrochemical parameters reflected by the impedance spectrum, the EIS curves for different conditions are fitted using the equivalent circuit model shown in Figure 6. Table 2 displays the relevant characteristic parameters: R_s represents the solution resistance, R_f is the resistance of the passive film covering the HEA, and Q_f as the film capacitance reflects the denseness and porosity of the film, respectively. The results demonstrate that the R_f value decreases with the increase of NaHSO₃ concentration, which implies that the protective performance of the film is weakened. In addition, the Q_f value increases with increasing NaHSO₃ concentration, indicating that the film is less dense and more porous, increasing the risk of ion reaction with the HEA matrix.

Figure 6

Equivalent circuit model used for fitting EIS data.

Table 2

Fitted electrochemical parameters for EIS data of the HEA at the film formation potential of 0 V_SCE in 3.5 wt.% NaCl solution with different concentrations of NaHSO₃.

NaHSO₃ concentrations	R_s (Ω·cm²)	Q_f (10⁻⁵ Ω⁻¹cm⁻²s^–n)	R_f (10⁴ Ω·cm²)
0 M	9.334	1.827	9.388
0.01 M	9.565	2.177	8.604
0.05 M	9.842	2.290	7.419
0.1 M	14.64	2.611	3.468

Numerous studies have shown that the semiconductor properties of the passivating film formed on the alloy surface are highly correlated with corrosion behaviour [39,40]. The Mott–Schottky theory is widely applied to study the semiconductor behaviour of the film [41]. Based on the theory, the semiconductor behaviour and carrier density of the film formed on CoCrFeNi HEA in 3.5 wt.% NaCl solution with different NaHSO₃ concentrations are investigated. The space charge layer capacitance followed by external electrode potential is calculated by the following Equation (1). (1) where the ‘+’ indicates n-type semiconductor, and the ‘−’ denotes p-type semiconductor. ϵ symbolises the dielectric constant of the passivation film, which is considered as ϵ = 12 [42]. ϵ₀ presents the vacuum dielectric constant (8.854 × 10⁻¹⁴ F·cm⁻¹). k is the Boltzmann constant (1.38 × 10⁻²³ J·K⁻¹), e represents the electronic charge (1.602 × 10⁻¹⁹ C). E and E_FB mean the applied electrode potential and the flat-band potential (V). T stands for the absolute temperature (K). Besides, N denotes the carrier concentration, including acceptor (N_A) and donor densities (N_D), which can be obtained from the slope (m) of C⁻² versus E by the following Equation (2). Among them, m is the slope of straight lines of the curves. (2) To further explore the effect of concentrations on the passivation film, the following Equation (3) is adopted to calculate the film thickness [43]. (3) Figure 7(a) shows the Mott–Schottky curves of the passivation films formed on the HEA in 3.5 wt.% NaCl solution with different NaHSO₃ concentrations. It can be seen that these plots contain two linear regions, showing the different electronic structures of the films. The region with a positive slope exhibits the n-type semiconductor behaviour, which is because the defects within the film are mainly oxygen vacancies and metal cation interstices. While the region with a negative slope displays p-type semiconductor behaviour, corresponding to the cation vacancies [44]. The calculated N_D and N_A in Figure 7(b) can reflect the degree of defects within the film: the lower the density, the more stable the internal structure of the film and the better the protection performance. Therefore, with the increase of concentration, the protection performance gradually drops. And when the HEA specimen is in the concentration of 0.1 M NaHSO₃, the defects within the film are the most abundant and the corrosion resistance of the HEA is the worst.

Figure 7

Mott–Schottky plots (a), carrier densities and thickness (b) of the passive films formed on the HEA at the film formation potential of 0 V_SCE in 3.5 wt.% NaCl solution with different concentrations of NaHSO₃.

Furthermore, the blue curve in Figure 7(b) depicts the thickness of the passivation films on the HEA surface in 3.5 wt.% NaCl solution with different concentrations. The thicker the film is, the better its protection is. Therefore, in the 0 M NaHSO₃ solution, the film is the thickest and its protective performance is the best. The results of the Mott–Schottky curve are consistent with those of the EIS tests, and the pattern of carrier density is in accord with that of Q_f value above in Table 2.

AC interference analysis

Figure 8 displays potentiodynamic polarisation curves and fitted corrosion current density (i_corr) values of CoCrFeNi HEA in 3.5 wt.% NaCl solution with 0.01 M NaHSO₃ under different i_AC. It can be observed that the shapes of polarisation curves measured at 0 and 30 A/m² are similar and both show apparent passivation region, which indicates the formation of a passivating film on the HEA surface. However, the curves tested at 50 and 100 A/m² present the obvious activation behaviour. Comparing the fitted i_corr values for polarisation curves exhibited in Figure 8(b) and Table 3 demonstrates that the increase in NaHSO₃ concentration significantly enhances the occurrence of corrosion, especially at high i_AC of 50 and 100 A/m². This may be due to the fact that more H⁺ ions are generated with an increased hydrogen evolution reaction under AC application [25], which promotes the corrosion reaction rate. Moreover, when the i_AC increases from 0 A/m² to 100 A/m², the corrosion potential (E_corr) generally tends negative shift, indicating a gradual increase in corrosion tendency. As listed in Table 3, the HEA without AC has a lower passive current density (i_p) and nobler pitting potential (E_p), compared with those tested at 30 A/m², revealing that the passivity of the HEA is dropped even at a low i_AC of 30 A/m², that is, the formation of pits is facilitated.

Figure 8

Potentiodynamic polarisation curves (a), i_corr values (b) of the HEA interfered at different i_AC in 3.5 wt.% NaCl solution with 0.01 M NaHSO₃.

Table 3

The fitted results for potentiodynamic polarisation curves of the HEA under various i_AC in 3.5 wt.% NaCl solution with 0.01 M NaHSO₃.

i_AC (A/m²)	E_corr (V)	i_corr (10⁻⁵ A·cm⁻²)	E_p (V)	i_p (10⁻⁵ A·cm⁻²)
0	−0.441	0.469	0.162	0.489
30	−0.485	3.190	0.035	1.104
50	−0.442	3.666	—	—
100	−0.525	12.351	—	—

The development process of pitting is generally as follows: pitting initiation, metastable pitting growth and stable pitting formation. Generally, E_p reflects the generation of stable pitting, and the potentials of pitting initiation and metastable pitting development are usually between E_p and E_prot [45,46]. Figure 9(a) compares the cyclic polarisation curves of CoCrFeNi HEA with or without AC in 3.5 wt.% NaCl solution containing 0.01 M NaHSO₃. Several current peaks between E_p and E_prot can also be observed, which implies the production of metastable pitting [47]. Moreover, as shown in Figure 9(b), the value of |E_p−E_prot| at 0 A/m² is close to that at 30 A/m², while the hysteresis loop area at 30 A/m² is much larger than that at 0 A/m². Table 4 lists the specific values of the above two index data. As known to all, the hysteresis loop area can be used to evaluate the resistance to pitting corrosion of metals, i.e. the smaller the hysteresis loop area, the less susceptible to pitting corrosion [48]. Thus, the imposed AC weakens the pitting resistance of the HEA.

Figure 9

Cyclic potentiodynamic polarisation curves (a), the values of |E_pit − E_prot| and hysteresis ring area (b) of the HEA at different i_AC.

Table 4

The fitted results for cyclic potentiodynamic polarisation curves of the HEA under various i_AC in 3.5 wt.% NaCl solution with 0.01 M NaHSO₃.

i_AC (A/m²)	\|E_p − E_prot\| (V)	Hysteresis ring area (10⁻⁴ A·cm⁻²·V)
0	0.482	0.854
30	0.447	1.721

To further investigate the effect of different applied i_AC on the formation and development process of metastable pitting, various low i_AC (0, 3, 5 and 10 A/m²) are selected and applied to the HEA electrode. The influence of AC interference on the pitting process of the HEA is statistically analysed by the potentiostatic polarisation test.

Metastable pitting analysis

Figure 10 depicts the current–time transient curves obtained for CoCrFeNi HEA under different i_AC in 3.5 wt.% NaCl solution with 0.01 M NaHSO₃. Clearly, the number and amplitude of the current density peaks increase with increasing i_AC, which demonstrates that AC promotes the development degree of metastable pitting corrosion. Since the applied AC generates more active sites on the HEA surface, resulting in high pitting sensitivity. Zhu et al. [24] reported that a large number of H atoms produced by imposed AC accumulate at the weak position and defects within the passivating film, promoting pitting nucleation. To quantitatively analyse the effect of AC interference on metastable pitting behaviour, the statistical method was adopted. The background current density (i_bg) in Figure 10 is determined to be 1.2 nA/cm² by potentiostatic polarisation test of CoCrFeNi HEA tested in Na₂SO₄ solution. When the peak value of the current transient (i_{peak value}) is greater than the i_bg, the peak is counted. Besides, for overlapping current density peaks, the current densities of adjacent two peaks differ by less than 50%, which can be viewed as the same metastable pitting peak, as shown in Figure 12 (b).

Figure 10

Current-time transient curves of CoCrFeNi HEA at different i_AC in 3.5 wt.% NaCl solution with 0.01 M NaHSO₃.

Figure 11(a) exhibits the variation trend of the average pitting number density (N_avg), and average i_{peak value} of metastable pitting on the HEA at different i_AC in 3.5 wt.% NaCl solution with 0.01 M NaHSO₃. The values of N_avg and average i_{peak value} increase significantly with the increase of applied i_AC. This indicates that the metastable pitting is easily excited to nucleation as i_AC rises. Figure 11(b) displays the average and maximum lifetime of the metastable pit at different i_AC, reflecting the growth and repassivation rates of metastable pitting. It can be seen that the average and maximum lifetime become longer with increasing i_AC. The application of AC extends the life of metastable pitting, stimulates its full nucleation and growth, and inhibits the repassivation process. Therefore, metastable pitting is easier to evolve into stable pitting, which is difficult to repair from repassivation. The related literatures [49,50] indicate that AC promotes the migration and ionisation of ions. As a result, metastable pitting gradually expands to stable pitting as metal elements dissolve.

Figure 11

Average pit number density, average peak value of current transient (a), lifetime of metastable pit (b) on CoCrFeNi HEA at different i_AC.

Figure 12 presents the single peak transient and overlapped peaks transient of the metastable pitting. The current peak in Figure 12(a) goes through three stages of development: increases slowly, rises rapidly and decreases sharply, which corresponds to the initiation, growth propagation and repassivation of the metastable pitting [51]. Where i_peak represents the peak current density; t_grow and t_rep are the growth time and repassivation time of metastable pitting, and t_Lifetime denotes the whole development process of metastable pitting. The overlapping peak transient is shown in Figure 12(b), which can be explained by the fact that metastable pitting occurs when structural defects exist in the vicinity of the pitting, leading to multiple pitting nucleations [52]. On the other hand, metastable pitting is an unstable process, causing overlapping peaks and current fluctuations.

Figure 12

Single peak transient (a) and overlapped peaks transient (b) of metastable pitting of the HEA.

Immersion test

Figure 13 shows the average corrosion rate of CoCrFeNi HEA in 3.5 wt.% NaCl solution with different NaHSO₃ concentrations. The corrosion rate of the HEA rapidly increases with increasing NaHSO₃ concentration. Combined with the above analysis results, it can be concluded that the corrosion resistance of the HEA is weakened with the increase of NaHSO₃ concentration. The corrosion morphologies of the HEA immersed in NaCl solution with different concentrations of NaHSO₃ for 8 d are exhibited in Figure 14. It can be observed that the number of pits produced on the HEA surface increases with an improved NaHSO₃ concentration. For 0.1 M NaHSO₃, some pits fuse and expand in the local region of the matrix. Therefore, the corrosion characteristics also confirm that the presence of reduces the anti-corrosion performance of the HEA.

Figure 13

The average corrosion rates of CoCrFeNi HEA in 3.5 wt.% NaCl solution with different concentrations of NaHSO₃.

Figure 14

Corrosion morphologies of the HEA immersed in 3.5 wt.% NaCl solution with different concentrations of NaHSO₃ for 8 d: (a,b) 0 M, (c,d) 0.05 M, (e,f) 0.1 M.

Figure 15 depicts the morphologies and EDS results of pitting on the HEA immersed in 3.5 wt.% NaCl solution with 0.1 M NaHSO₃ for 8 d. Through the semiquantitative analysis result (Figure 15b), it can be found that the content of Ni element inside the pitting is significantly reduced compared with the matrix, and the percentage of Fe element is almost unchanged, whereas the contents of Cr and Co increase, which reveals that the occurrence of pitting may be induced by the selective dissolution of Ni element. The mapping result in Figure 15(c) also further proves this. The dissolution behaviour can also reduce the compactness and stability of the passive film covering the HEA, and its protection will be weakened.

Figure 15

Morphologies and EDS results of pits on the HEA immersed in 3.5 wt.% NaCl solution with 0.1 M NaHSO₃ for 8 d.

Moreover, the composition of dissolved oxide inclusion after the immersion test is detected in Figure 16. Compared with the original oxide inclusion in Figure 2, the contents of Cr, Fe, Co and Ni elements markedly decrease. Obviously, pitting occurs at its surrounding oxide, which may be caused by the effect of micro-galvanic corrosion cell formed between oxide and matrix. This phenomenon also exists in our previous study [53], which is detrimental to the corrosion resistance of the alloy.

Figure 16

EDS results of oxide inclusion within the HEA immersed in 3.5 wt.% NaCl solution with 0.1 M NaHSO₃ for 8 d.

Corrosion mechanism

According to all the above-mentioned results, the corrosion mechanism of CoCrFeNi HEA under different conditions is shown in Figure 17. A large number of studies have reported that Cl⁻ has a significant negative effect on the corrosion sensitivity of alloys in NaCl environment [54 56]. As displayed in Figure 17(a), the corrosion of HEA is mainly affected by Cl⁻ in 3.5 wt.% NaCl solution and the main cathode reaction is the oxygen reduction reaction (Reaction 4). During the corrosion process, Cl⁻ ions are adsorbed on the active sites and defects of the passivating film, promoting the anodic dissolution (Reaction 5), which is the main reason for the pitting initiation on the HEA surface [11]. (4) (5) In the NaCl solution containing (Figure 17(b)), the thickness of the passive film thins and its integrity becomes worse, compared with that without the addition of . Moreover, the cathodic reaction is an oxygen-consumption reaction in the initial corrosion stage. The presence of has a certain effect on the cathode reaction process [57]. With the ionisation of (Reaction 6), a large amount of H⁺ enters the solution system, increasing H⁺ concentration in the local region of the bulk solution. That is, hydrogen evolution may occur as part of cathode reactions (Reaction 7). Thus, many H atoms penetrate into the defect sites within the passive film, which increases film instability [58]. Therefore, more pits are produced as NaHSO₃ concentration increases. It is well known that SO₃ ^2- is unstable in solution and tends to oxidise (Reactions 8), leading to the formation of FeSO₄ (Reaction 9). Where Fe²⁺ comes from the chemical dissolution of oxides presented in Figure 16. The FeSO₄ continuously oxidises and forms H₂SO₄ through the cyclic acidification mechanism (Reactions 10,11) [59], resulting in further acidification of the solution. Therefore, SO₄ ^2- plays a catalytic role in the reaction process, accelerating the corrosion of the HEA. In addition, Cl⁻ ion increases the conductivity of the matrix/solution interface and promotes the electrochemical reaction. Hence, the combined effect of Cl⁻ and promotes the anodic dissolution of the HEA (Figure 13). (6) (7) (8) (9) (10) (11)

Figure 17

Corrosion mechanism diagram of the HEA under different conditions (a) in 3.5 wt.% NaCl solution, (b) in 3.5 wt.% NaCl solution with NaHSO₃, (c) interfered by AC in 3.5 wt.% NaCl solution with NaHSO₃.

As expressed in Figure 17(c), the hydrogen evolution reaction (Reaction. 7) is the dominant cathode reaction due to the influence of an external AC. The role of the applied AC is generally reflected in two aspects [22]. On the one hand, in the positive half cycle of AC, the electrolysis of water is motivated (Reaction. 12), and the continuous production of oxygen accelerates the oxidation process of SO₃ ^2- and the cyclic acidification, thus generating a large amount of H⁺. Additionally, under AC application, with the increase in the migration rate [60] and adsorption [25,61] of ions, the reaction between harmful ions and oxides/hydroxides in the passivating film is intensified, which makes the corrosion of the HEA more serious and even difficult to form a passive film. On the other hand, in the negative half cycle of AC, the hydrogen evolution reaction is accelerated (Reaction. 13) [60]. Hydrogen ions produced by electrolysis of water and ionisation of participate in the cathodic reaction process, thus stably generating hydrogen atoms. Some parts of H atoms combine together to form H₂, as confirmed by bubbles on the counter electrode surface. The produced H₂ diminishes the bonding force between the passivated film and the matrix, inducing a certain internal pressure. As a result, the defects within the film increase, and its thickness thins, which facilitates the rupture of the film [25,62,63]. (12) (13) The influence of imposed AC on the corrosion behaviour of the HEA can be reflected by Figure 8(a). With the increase of i_AC, the surface state of the HEA is changed from passivation to activation. When low i_AC is applied, the formation and dissolution rates of the passivating film are easy to reach a dynamic equilibrium state, i.e. the HEA surface is passivated. This phenomenon can be due to the fact that the number of generated hydrogen ions and atoms are relatively small, which difficulty affects the formation of the equilibrium. However, as the i_AC increases to 50 A/m², the concentration of produced hydrogen ions and atoms reaches a high level, which greatly damages the stability of the film, accelerates its dissolution, and reduces the passivity of the HEA.

Conclusions

The corrosion behaviour and mechanism of CoCrFeNi HEA with imposed AC in a simulated marine environment containing different concentrations of NaHSO₃ were systematically investigated, and the main conclusions were as below:

With the increase of NaHSO₃ concentration, the passivation region became narrow, and the i_p value increased, suggesting that the corrosion of the HEA was facilitated. In 0.1 M NaHSO₃ solution, the polarisation curve exhibited a visible transition zone, indicating that the increase in NaHSO₃ concentration decreased the passivity of the HEA.

As the concentration of NaHSO₃ increased, the defect densities within the passivation film increased and its protection performance became weaker, thus diminishing the anti-corrosion property of the HEA. Moreover, the corrosion rate of the HEA in NaCl solution with 0.1 M NaHSO₃ was about 2.4 times that without the addition of NaHSO₃.

Due to the influence of an external AC, the corrosion behaviour of the HEA was evolved from passivation characteristic at low i_AC to active state at high i_AC. In addition, the application of AC prolonged the life of metastable pitting, stimulated its full nucleation and growth, and inhibited the repassivation process. With the increase of i_AC, the N_avg value significantly increased from 5 to 17 mm⁻².

The formation of pitting on the HEA surface was related to the selective dissolution of Ni element. And the selective dissolution of Cr, Fe, Co and Ni elements occurred inside the oxide inclusion.

Footnotes

Disclosure statement

No potential conflict of interest was reported by the author(s).

References

Sharp

Billingham

Robinson

MJ.

The risk management of high-strength steels in jack-ups in seawater. Mar Struct. 2001; 14: 537–551.

Dong

Liu

Zhang

. Corrosion failure analysis of low alloy steel and carbon steel rebar in tropical marine atmospheric environment: outdoor exposure and indoor test. Eng Fail Anal. 2021; 129: 105720.

Yeh

Chen

Lin

. Nanostructured high-entropy alloys with multiple principal elements: novel alloy design concepts and outcomes. Adv Eng Mater. 2004; 6: 299–303.

Zhang

Zhao

. Effects of stress states and strain rates on mechanical behavior and texture evolution of the CoCrFeNi high-entropy alloy: experiment and simulation. J Alloy Compd. 2021; 851: 156779.

Dolique

Thomann

Brault

. Thermal stability of AlCoCrCuFeNi high entropy alloy thin films studied by in-situ XRD analysis. Surf Coat Tech. 2010; 204: 1989–1992.

Hsu

Sheu

Yeh

. Effect of iron content on wear behavior of AlCoCrFe_xMo_0.5Ni high-entropy alloys. Wear. 2010; 268: 653–659.

Chuang

Tsai

Wang

. Microstructure and wear behavior of Al_xCo_1.5CrFeNi_1.5Ti_y high-entropy alloys. Acta Mater. 2011; 59: 6308–6317.

Salishchev

Tikhonovsky

Shaysultanov

. Effect of Mn and V on structure and mechanical properties of high-entropy alloys based on CoCrFeNi system. J Alloy Compd. 2014; 591: 11–21.

Sathiyamoorthi

Basu

Kashyap

. Thermal stability and grain boundary strengthening in ultrafine-grained CoCrFeNi high entropy alloy composite. Mater Design. 2017; 134: 426–433.

10.

Huo

Zhou

Fang

. Strain-rate effect upon the tensile behavior of CoCrFeNi high-entropy alloys. Mater Sci Eng A. 2017; 689: 366–369.

11.

Nascimento

Donatus

Ríos

. Passive film composition and stability of CoCrFeNi and CoCrFeNiAl high entropy alloys in chloride solution. Mater Chem Phys. 2021; 267: 124582.

12.

Yang

Shang

Wang

. Effect of constituent elements on the corrosion resistance of single-phase CoCrFeNi high-entropy alloys in NaCl solution. Acta Metall Sin. 2018; 54: 905–910.

13.

Liu

Tsai

CW.

Effect of Ge addition on the microstructure, mechanical properties, and corrosion behavior of CoCrFeNi high-entropy alloys. Intermetallics. 2021; 132: 107067.

14.

Hsu

Chiang

JK.

Corrosion behavior of FeCoNiCrCux high-entropy alloys in 3.5% sodium chloride solution. Mater Chem Phys. 2005; 92: 112–117.

15.

Zhu

Liu

Zhang

CY.

Laser cladding of CoCrFeNi high-entropy alloy coatings: compositional homogeneity towards improved corrosion resistance. Mater Lett. 2022; 318: 132133.

16.

Gan

Yan

. A non-equiatomic FeNiCoCr high-entropy alloy with excellent anti-corrosion performance and strength-ductility synergy. Corros Sci. 2021; 183: 109341.

17.

Matei

Pencea

Branzei

. Corrosion resistance appraisal of TiN, TiCN and TiAlN coatings deposited by CAE-PVD method on WC-Co cutting tools exposed to artificial sea water. Appl Surf Sci. 2015; 358: 572–578.

18.

Baltazar-Zamora

Bastidas

Santiago-Hurtado

. Effect of silica fume and fly ash admixtures on the corrosion behavior of AISI 304 embedded in concrete exposed in 3.5% NaCl solution. Materials (Basel). 2019; 12: 4007.

19.

Liu

. Effect of SO₂ content on SCC behavior of E690 high-strength steel in SO₂-polluted marine atmosphere. Ocean Eng. 2018; 164: 256–262.

20.

Milito

Murano

Castellano

. Antioxidant and immune response of the sea urchin Paracentrotus lividus to different re-suspension patterns of highly polluted marine sediments. Mar Environ Res. 2020; 160: 104978.

21.

Gong

Liu

GX.

Stress corrosion cracking behavior of rusted X100 steel under the combined action of Cl⁻ and HSO₃ ⁻ in a wet-dry cycle environment. Corros Sci. 2020; 165: 108382.

22.

Pan

Liu

. Electrochemical and stress corrosion mechanism of submarine pipeline in simulated seawater in presence of different alternating current densities. Materials (Basel). 2018; 11: 1074.

23.

Kim

Muralidharan

. Electrochemical studies on the alternating current corrosion of mild steel under cathodic protection condition in marine environments. Electrochim Acta. 2006; 51: 5259.

24.

Zhu

. Effect of AC on stress corrosion cracking behavior and mechanism of X80 pipeline steel in carbonate/bicarbonate solution. Corros Sci. 2014; 87: 224–232.

25.

Zhu

Zhao

Yuan

. Study on corrosion behavior and mechanism of CoCrFeMnNi HEA interfered by AC current in simulated alkaline soil environment. J Electroanal Chem. 2021; 882: 115026.

26.

Bei

Otto

. Recovery, recrystallization, grain growth and phase stability of a family of FCC-structured multi-component equiatomic solid solution alloys. Intermetallics. 2014; 46: 131–140.

27.

Song

Zhu

. The corrosion of Al-supersaturated Mg matrix and the galvanic effect of secondary phase nanoparticles. Corros Sci. 2021; 184: 109410.

28.

Cooper

Dhers

. Effect of oxygen content upon the microstructural and mechanical properties of type 316L austenitic stainless steel manufactured by hot isostatic pressing. Metall Mater Trans A. 2016; 47: 4467–4475.

29.

Moravcik

Peighambardoust

Motallebzadeh

. Interstitial nitrogen enhances corrosion resistance of an equiatomic CoCrNi medium-entropy alloy in sulfuric acid solution. Mater Charact. 2020; 172: 110869.

30.

Garet

Brass

Haut

. Hydrogen trapping on non metallic inclusions in cr-mo low alloy steels. Corros Sci. 1998; 40: 1073–1086.

31.

Pao

Muto

Sugawara

Pitting at inclusions of the equiatomic CoCrFeMnNi alloy and improving corrosion resistance by potentiodynamic polarization in H₂SO₄. Corros Sci. 2021; 191: 109748.

32.

Hou

Xiong

Liu

. Effects of LaAlO₃ and La₂O₂S inclusions on the initialization of localized corrosion of pipeline steels in NaCl solution. Scripta Mater. 2020; 177: 151–156.

33.

Villavicencio

Ulloa

Lozada

. The role of non-metallic Al₂O₃ inclusions, heat treatments and microstructure on the corrosion resistance of an API 5L X42 steel. J Mater Res Technol. 2020; 9: 5894–5911.

34.

Zhao

Pan

Wang

. Corrosion and passive behavior of Al_xCrFeNi_3-x (x = 0.6, 0.8, 1.0) eutectic high entropy alloys in chloride environment. Corros Sci 2022; 208: 110666.

35.

Cheng

YF.

Electrochemical polarization behavior of X70 steel in thin carbonate/bicarbonate solution layers trapped under a disbonded coating and its implication on pipeline SCC. Corros Sci. 2010; 52: 2511–2518.

36.

Zhang

Dai

. Anisotropic corrosion behavior of Sc and Zr modified Al-Mg alloy produced by selective laser melting. Corros Sci. 2020; 170: 108657.

37.

Martins

Osório

Souza

MEP

. Effects of Zr content on microstructure and corrosion resistance of Ti-30Nb-Zr casting alloys for biomedical applications. Electrochim Acta. 2008; 53: 2809–2817.

38.

Yang

Cheng

YF.

Passivity degradation and photocorrosion of X52 carbon steel under visible light illumination in concentrated carbonate/bicarbonate solutions. Corros Sci. 2020; 172: 108727.

39.

Alnajjar

Christien

Barnier

. Influence of microstructure and manganese sulfides on corrosion resistance of selective laser melted 17-4 PH stainless steel in acidic chloride medium. Corros Sci. 2020; 168: 108585.

40.

Ogunsanya

Hansson

CM.

The semiconductor properties of passive films and corrosion behavior of stainless steel reinforcing bars in simulated concrete pore solution. Materialia. 2019; 6: 100321.

41.

Cui

Chen

Dou

. Passivation behavior and surface chemistry of 2507 super duplex stainless steel in artificial seawater: influence of dissolved oxygen and pH. Corros Sci. 2019; 150: 218–234.

42.

Hao

. Effect of nitrogen on corrosion behaviour of a novel high nitrogen medium-entropy alloy CrCoNiN manufactured by pressurized metallurgy. J Mater Sci Technol. 2018; 34: 1781–1790.

43.

Ningshen

Kamachi

Mittal

. Semiconducting and passive film properties of nitrogen-containing type 316LN stainless steels. Corros Sci. 2007; 49: 481–496.

44.

Taveira

Montemor

da Cunha Belo

. Influence of incorporated Mo and Nb on the Mott-Schottky behaviour of anodic films formed on AISI, 304L. Corros Sci 2010; 52: 2813–2818.

45.

Yuan

Wang

Cao

. Natural passivation behavior and its influence on chloride-induced corrosion resistance of stainless steel in simulated concrete pore solution. J Mater Res Technol. 2020; 9: 12378–12390.

46.

Zaid

Saidi

Benzaid

. Effects of pH and chloride concentration on pitting corrosion of AA6061 aluminum alloy. Corros Sci. 2008; 50: 1841–1847.

47.

Al Ameri

Cho

. Critical conditions for pit initiation and growth of austenitic stainless steels. Corros Sci. 2015; 92: 209–216.

48.

Guo

Chen

. Pitting corrosion behavior of friction-surfaced 17-4PH stainless steel coatings with and without subsequent heat treatment. Corros Sci. 2021; 193: 109887.

49.

Cheng

YF.

Effects of alternating current on corrosion of a coated pipeline steel in a chloride-containing carbonate/bicarbonate solution. Corros Sci. 2010; 52: 612–619.

50.

Kuang

Cheng

YF.

Understand the AC induced pitting corrosion on pipelines in both high pH and neutral pH carbonate/bicarbonate solutions. Corros Sci. 2014; 85: 304–310.

51.

Bai

. Effects of boron on microstructure and metastable pitting corrosion behavior of super 304H austenitic stainless steel. J Electrochem Soc. 2015; 162: 473–481.

52.

Tang

Zuo

Zhao

XH.

The metastable pitting behaviors of mild steel in bicarbonate and nitrite solutions containing Cl⁻. Corros Sci. 2008; 50: 989–994.

53.

Zhu

Yuan

. A comparative study on the corrosion behavior of CoCrNi medium-entropy alloy and 316L stainless steel in simulated marine environment. Intermetallics. 2021; 139: 107370.

54.

Yue

Ren

Huang

. The role of Cl⁻ in the formation of the corrosion products and localised corrosion of 15Cr martensite stainless steel under an CO₂-containing extreme oilfield condition. Corros Sci. 2022; 194: 109935.

55.

Wang

Corrosion of low carbon steel in atmospheric environments of different chloride content. Corros Sci. 2009; 51: 997–1006.

56.

Wang

Shi

Xiang

. Study of the corrosion behaviour of sensitized 904L austenitic stainless steel in Cl⁻ solution. Corros Sci. 2021; 181: 109234.

57.

Dong

Sun

. Effects of solution pH and Cl⁻ on electrochemical behaviour of an aermet100 ultra-high strength steel in acidic environments. Corros Sci. 2011; 53: 4159–4165.

58.

Zhong

Wang

. Synergistic effect of erosion and hydrogen on properties of passive film on 2205 duplex stainless steel. J Mater Sci Technol. 2021; 67: 1–10.

59.

Nishimura

Shiraishi

Maeda

Hydrogen permeation and corrosion behavior of high strength steel MCM 430 in cyclic wet-dry SO₂ environment. Corros Sci. 2004; 46: 225–243.

60.

Yin

. Development of a real-time AC/DC data acquisition technique for studies of AC corrosion of pipelines. Corros Sci. 2012; 61: 215–223.

61.

Wang

Zhang

Gao

. Galvanic effect and alternating current corrosion of steel in acidic red soil. Metals (Basel). 2022; 12: 12020296.

62.

Pan

Sun

Liu

. Hydrogen effects on passivation and SCC of 2205 DSS in acidified simulated seawater. Corros Sci. 2022; 208: 110640.

63.

Shubina Helbert

Nazarov

Vucko

. Hydrogen effect on the passivation and crevice corrosion initiation of AISI 304L using scanning Kelvin probe. Corros Sci. 2021; 182: 109225.