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Article

Electrochemical Behavior and Passive Film Properties of Hastelloy C22 Alloy, Laser-Cladding C22 Coating, and Ti–6Al–4V Alloy in Sulfuric Acid Dew-Point Corrosion Environment

by
Chao Zheng
,
Zongde Liu
*,
Quanbing Liu
,
Yao Kong
,
Shengyang Guo
and
Congcong Liu
Key Laboratory of Power Station Energy Transfer Conversion and System of Ministry of Education, North China Electric Power University, Beijing 102206, China
*
Author to whom correspondence should be addressed.
Metals 2022, 12(4), 683; https://doi.org/10.3390/met12040683
Submission received: 9 March 2022 / Revised: 9 April 2022 / Accepted: 12 April 2022 / Published: 15 April 2022
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Abstract

:
The electrochemical behavior and passive film properties of Hastelloy C22 alloy, laser-cladding C22 coating, and Ti–6Al–4V alloy in sulfuric acid dew-point corrosion environment were investigated through a combination of electrochemical measurements and surface analyses. The C22 alloy and laser-cladding C22 coating exhibited similar passivation and repassivation behavior without pitting corrosion, resulting from a similar passive film with a bilayer structure consisting of a Cr2O3-dominated compact inner layer and a porous outer layer containing oxides of Mo and hydroxides of Ni and Cr. The slightly poorer corrosion resistance and higher sensitivity to localized corrosion exhibited by the C22 coating were attributed to the microscale heterogeneity of the passive film resulting from the element segregation in the microstructure introduced by the laser-cladding process. The corrosion of the TC4 alloy performed as the preference dissolution of the β phase. Compared to the C22 alloy and C22 coating, the TC4 alloy exhibited more stable passivation behavior but poorer corrosion resistance, which is attributed to a compact but less protective single-layer passive film consisting of oxides of Ti and Al. An increase in temperature degrades passive film stability and accelerates the charge transfer process.

1. Introduction

In addition to emphasizing legislation to reduce emissions of acidic gases and environmental pollution, about ninety percent of power plants in China use wet flue gas desulfurization (WFGD) technology, which uses limestone and limestone slurry to absorb SO2 [1,2]. However, the cooling and humidification of flue gas after desulfurization results in the condensation of acid droplets that mainly contain sulfuric acid, which is formed by the SO3 remaining in the flue gas, and consequently produces significant acid dew-point corrosion on the low-temperature surface of flue gas desulfurization (FGD) and subsequent facilities such as back-end ductworks, reheaters, fans, and chimneys [3,4].
The use of alloys with high resistance to dew-point acid corrosion, like titanium alloys and Ni-Cr-Mo alloys, is an efficient method to overcome the corrosion failure of these facilities and has been utilized for decades in the most corrosive environments of FGD systems [5,6] and other low-temperature equipment in the United States and Europe [7,8]. Titanium alloys are used as lining materials for chimneys for FGD systems in China. Moreover, several laboratory investigations have reported the corrosion behavior of titanium and Ni-Cr-Mo alloys in sulfuric acid dew-point corrosion environments. Shan Zhao et al. reported that a titanium alloy exhibited better corrosion resistance in stack corrosion environments under WFGD conditions compared with polyurea, vinyl ester glass flake (VEGF), fiberglass-reinforced plastic (FRP), 09CrCuSbsteel and foam glass [9]. De Souza et al. comprehensively studied the electrochemical behavior of pure titanium and titanium–tantalum alloys in 20–80 wt.% H2SO4 at 25–75 °C [10]. Shoemaker et al. verified the excellent corrosion resistance properties of Ni-Cr-Mo alloys through laboratory tests and field observation of FGD systems in service [5]. Yoshio Takizawa et al. compared the corrosion behavior of Ni-Cr-Mo alloy containing different Mo content in 60% and 80% H2SO4 at 120 °C. Yoshio and Katsuo [11], Rajendran and Rajeswari [12] and Darowicki and Krakowiak [13] studied the pitting corrosion resistance of Ni-Cr-Mo alloys in different simulated scrubbed flue gas environments. Mishra investigated the role of alloying elements in the uniform corrosion resistance of Ni-Cr-Mo (W) alloys in concentrated acids [14]. These researchers widely agree that the excellent resistance of titanium and Ni-Cr-Mo alloys is attributed to the protective oxide layer and passivation occurring on the surface. Wang et al. investigated the passivation behavior of titanium in fluoride-containing sulfuric acid (pH = 1.0 to 5.0) at 25 °C [15,16]. Cui et al. revealed pure titanium’s electrochemical behavior and surface characteristics in a mixed solution of H2SO4, HCl, HNO3, and NaF. [17]. Lloyd et al. revealed the composition and structure of the passive films of a series of Ni-Cr-Mo alloys in 1.0 mol·L−1 NaCl + 0.1 mol·L−1 H2SO4 solutions at 25 to 85 °C [18,19,20]. Many researchers applied electrochemical and surface analytical techniques (X-ray photoelectron, Auger spectroscopies and atomic force microscopy) to reveal the passive behavior of a series of Ni-Cr-Mo alloys under different conditions [21,22,23,24]. However, limited studies exist on the electrochemical behavior and correlative passive film properties of titanium alloys and Ni-Cr-Mo alloys in a sulfuric acid dew-point corrosion environment introduced by applying WFGD technology could limit the in-depth understanding and life evaluation of these materials in desulfurized flue gas environments.
It is widely agreed that the Ni-Cr-Mo alloys, just like C276 and C22, are the best candidates for construction or lining materials of facilities suffering the sulfuric acid dew-point corrosion. [5,8]. However, the extensive application of Ni-Cr-Mo alloys in FGD systems and other facilities is restricted because of the high construction cost resulting from the persistently high price of nickel and molybdenum raw materials. It is especially the shortage of nickel resources in China that leads to the application of titanium alloys with lower performance instead. [5,25] Laser cladding is a widely used technology to attain excellent surface properties by preparing a thin coating. Its confirmed advantages include high energy density, metallurgical bonding, low dilution rate, and a small heat-affected zone [26,27,28,29].
Furthermore, the high-value performance of laser-cladding technology contributes to its potential industrial application, especially in the application demanding high-performance but expensive materials like Ni-Cr-Mo alloys. Thus, preparing a Ni-Cr-Mo alloy coating on structural steel is considered cost-effective to completely replace using Ni-Cr-Mo alloys as construction or lining materials [30]. Li et al. reported that laser-cladding Ni-Cr-Mo alloy coatings exhibited equivalent or better corrosion resistance compared to bulk Ni-Cr-Mo alloy in high-temperature environments [31,32,33]. Wang et al. investigated the corrosion behavior of laser-cladding Ni-Cr-Mo alloy coating in neutral and dilute acid solutions at room temperature [34,35]. During the laser-cladding process, researchers attributed this performance improvement to grain refinement and non-equilibrium grain-boundary segregation. However, there is limited research contributing to the comparison between the corrosion behavior of laser-cladding Ni-Cr-Mo alloy coatings and bulk Ni-Cr-Mo alloy in a simulated sulfuric acid dew-point corrosion environment.
In our previous study, the laser-cladding Hastelloy C22 coating exhibited lower resistance and higher resistance in a simulated sulfuric acid dew-point corrosion environment (at 50–70 °C in 50 wt.% H2SO4) compared to the C22 alloy and TC4 alloy, respectively. [35]. Our understanding of the passive behavior of laser-cladding C22 coating in this aggressive environment is still incomplete; therefore, further investigation of the electrochemical behavior and passive film properties under these conditions will help elucidate the corrosion behavior and optimize laser-cladding technology under practical conditions. The comparison between the Hastelloy C22 alloy and Ti–6Al–4V alloy, which are considered the top-level anticorrosive materials in desulfurized flue gas environments, will introduce the application prospect of laser-cladding technology in corrosion prevention.
In the present study, the electrochemical behavior and passive film properties of Hastelloy C22 alloy, laser-cladding C22 coating on carbon structural steel, and Ti–6Al–4V alloy (TC4) in a simulated sulfuric acid dew-point corrosion environment were investigated and compared using electrochemical measurements and surface analysis. Potentiodynamic polarization was applied to evaluate and compare the corrosion behavior of these three materials. Three different potentials were chosen as different passive conditions during potentiostatic polarization; meanwhile, the corroded morphology and passive film properties at each potential are analyzed. The effect of temperature on the electrochemical behavior and passive film properties is simultaneously discussed.

2. Materials and Methods

2.1. Materials and Specimens

The materials tested in this study were commercial standard Hastelloy C22 alloy, Ti–6Al–4V alloy (TC4), and self-fabricated laser-cladding C22 coatings. The chemical composition of the C22 powder used for laser cladding, C22 alloy, and TC4 alloy is listed in Table 1.
The C22 powder with a diameter in the range of 46–150 μm was provided by the Beijing General Research Institute of Mining and Metallurgy (BGRIMM), Beijing, China. The substrate for the laser-cladding process was a Q235 carbon structural steel plate with 10 mm thickness. The coaxial laser-cladding process employed a high-power fiber laser (ZKZM-3000, zKzM Laser Technology Co., Ltd., Xi’an, China) and a self-designed coaxial powder feed nozzle. The high-power fiber laser produces a collimated laser beam with a maximum 3000 W power and a beam spot of 1.4 mm in diameter. The coaxial powder feed nozzle was designed to eject three powder streams from uniformly distributed directions around the laser beam. The powder streams were formed by argon carrier gas with a flow rate of 5 L/min. A multitrack and multilayer-cladding process with an overlap ratio of 40%, and a laser scan speed of 8 mm/s was adopted to prepare 2 mm thick coating specimens. The technological process employed during the fabrication of laser-cladding C22 coatings was optimized to obtain defect-free coatings metallurgical bonding with the substrate, and the microstructure mainly contains primary dendrite and eutectics with Mo segregation forming on the grain boundary, which is specifically presented and discussed in one of our previous studies [36].
Electrochemical test specimens with a dimension of 10 mm × 10 mm × 3 mm were cut from C22 alloy, as-prepared C22 coatings, and TC4 alloy. After being abraded with 400 grit Al2O3 paper and then sequentially ultrasonically cleaned in distilled water and acetone, each specimen was fixed in epoxy resin with a brass rod attached at the back to create an electrical connection. As a working electrode, a 10 mm × 10 mm exposed surface was abraded with 2000 grit Al2O3 paper and then polished with a 1.5 μm diamond abrasive; the thicknesses of the C22 coatings were about 1600 μm after surface treatment. All prepared specimens were sequentially ultrasonically cleaned in distilled water and alcohol, then stored in a glass-drying vessel for 24 h prior to electrochemical testing.

2.2. Electrochemical Measurement

All electrochemical measurements were conducted with an electrochemical workstation (CHI660e, Chenhua, Shanghai, China) in a standard three-electrode cell, using a pure platinum sheet as the counter electrode and a mercurous sulfate electrode (MSE) with a Luggin capillary salt bridge as the reference electrode. All potential values reported in this work were versus MSE. To simulate the sulfuric acid dew-point corrosion environment in FDG systems and back-end ductworks [3,11,37], a 50 wt.% H2SO4 solution prepared using analytical grade 98 wt.% H2SO4 and deionized water was selected as the electrolyte solution, and the temperature was maintained at 50 °C, 60 °C, and 70 °C through a water bath. The open circuit potential (OCP), potentiodynamic polarization curve, and potentiostatic polarization curve of each specimen was measured, and each measurement was repeated three times. The most representative measurements were presented for discussion in this study.
Prior to the polarization tests, all three specimens were immersed in the electrolyte solution and the OCPs were measured continuously for 1 h. The OCP measurements revealed that 1 h was sufficient to achieve a stabilized state and passive films. Potentiodynamic polarization scans were performed forward and back at a scan rate of 0.5 mV·s−1 with potential ranges of −800 to 960 mVmse for the C22 alloy, −800 to 800 mVmse for the C22 coating, and −1300 to 1200 mVmse for the TC4 alloy. These parameters were set to study the passivation, transpassivation and repassivation behavior of specimens in 50 wt.% H2SO4 solution and the influence of temperature. Potentiostatic polarization tests were performed at potentials of 100 mVmse, 450 mVmse, and 600 mVmse for 10 h to form different surface states in different anodic regions. Subsequently, specimens underwent the potentiostatic polarization tests and then were rinsed in deionized water and dried to prepare for the surface analysis described below.
EIS measurements were conducted with a frequency range of 105 Hz to 10−2 Hz with an amplitude of 10 mV peak-to-peak AC signals at measured open circuit potential. Nyquist and Bode curves were obtained to evaluate corrosion resistance and analyze electrochemical behavior.

2.3. Surface Analysis

The surface morphology and composition of specimens after 10 h potentiostatic polarization were analyzed using SEM and XPS, respectively. A scanning electron microscope (FEI Quanta 200F, Eindhoven, The Netherlands) equipped with EDS (EDAX, Pleasanton, CA, USA) was applied to observe the surface morphology and elemental composition. Compositions of passive films and chemical states of alloying elements were analyzed through XPS measurements on an X-ray photoelectron spectrometer (K-Alpha, ThermoFisher, Waltham, MA, USA) with a monochromatic Al Kα (1486.6 eV). The X-ray gun was operated at 150 W (15 kV, 10 mA) and the photoelectron take-off angle was 90°. During the experiments, a base pressure was maintained at approximately 5 × 10−9 mbar. High-resolution spectra with an analyzed area of 400 μm2 were recorded at a pass energy of 50.0 eV with an energy step of 0.1 eV, whereas survey spectra were recorded at a pass energy of 150.0 eV with an energy step of 1 eV. The values of measured binding energies were corrected by referencing to C 1s (284.8 eV). A Gaussian–Lorentzian product function and a Shirley background subtraction were employed to fit the spectra by applying Avantage software (version 5.984, Waltham, MA, USA).

3. Results

3.1. Open Circuit Potential (OCP) Measurements

Figure 1 presents the OCP vs. time curves of the C22 alloy, C22 coating, and TC4 alloy in a 50 wt.% H2SO4 solution. Note that the OCP of the C22 coating stays more negative than that of the C22 alloy for the duration of the test. The OCPs for the C22 alloy and the C22 coating exhibit similar variation trends with respect to time and temperature. The values measured at 50 and 60 °C slightly increase with testing time, and the increase rates gradually decline to a steady-state later in the experiment (Figure 1a,b). Similar shifts and trends of the OCP with respect to time were reported for Ni-base alloys in acid solutions and stainless in simulated desulfurized flue gas condensates [38,39,40]. In particular, the rate of increase in OCP for the C22 coating is higher earlier in the testing period, then undergoes a sharp drop midway through. There is no remarkable difference in temperature influence, 20 to 25 mV more positive at 60 °C, between the relatively stable OCPs at the end of testing at 50 and 60 °C, whereas significant negative shifts of OCPs occur as the temperature elevates to 70 °C, which are 200 and 180 mV more negative than values at 50 °C for the C22 alloy and the C22 coating, respectively. Dou et al. [38] and García-Antón et al. [39] reported a contrary influence of increasing temperature on OCP; meanwhile, Lloyd [20] and Cardoso [40] reported a similar negative shift of OCP for Ni-Cr-Mo alloys in 1.0 mol·L−1 NaCl + 1.0 mol·L−1 H2SO4 and Ni-Cr-Fe alloy in NaCl solution, respectively. The reaction between the noble air-formed oxide film and electrolyte in early testing at 50 and 60 °C leads to the rapid thickening of the passive film, and a stable thickness is acquired as the test progresses. The C22 coating presents a higher corrosion tendency and a more notable evolution of a passive film compared to the C22 alloy. As the temperature increases to 70 °C, the equilibrium between the dissolution and growth of the passive film is immediately achieved early in testing as a result of the more considerable reaction, and presents a higher corrosion tendency. The corrosion reaction thermodynamic stability and passive film growth/dissolution rate of the C22 alloy and C22 coating are clearly affected until 70 °C.
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Figure 1. Open circuit potential of the C22 alloy (a), C22 coating (b), and TC4 alloy (c) in 50 wt.% H2SO4 at 50–70 °C.
Figure 1. Open circuit potential of the C22 alloy (a), C22 coating (b), and TC4 alloy (c) in 50 wt.% H2SO4 at 50–70 °C.
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The OCP of the TC4 alloy exhibits different variation trends from those of the C22 alloy and the C22 coating under the same testing conditions. It is clear that the OCP for TC4 maintains relatively stable and more negative values at each temperature after a short testing time. A slight decrease and then increase occurs during the initial period at 60 and 70 °C. As testing concludes, the same value of OCP is obtained at 50 and 70 °C, accompanied by the 7 mV more negative value at 60 °C. The OCP gap between the different testing temperatures was too small to verify the influence of temperature on corrosion tendencies. For titanium alloys in H2SO4 solution, researchers have reported different variations with negative shifts that eventually achieve steady-state values in the active region [17,41,42,43]. Wang et al. reported a similar variation in OCP with a larger increase at initial immersion for titanium and its alloy in H2SO4 solution sulfuric acid [16,44]. The air-formed oxide film on the TC4 was only slightly transformed during testing resulting from the strong oxidizing electrolyte, while the increase in temperature only slightly affected its corrosion reaction thermodynamic stability.

3.2. Potentiodynamic Polarization Measurement

The polarization curves of the C22 alloy, C22 coating, and TC4 alloy in 50 wt.% H2SO4 solution at 50 to 70 °C are presented in Figure 2 and Figure 3, respectively. The scanning direction during polarization is marked with arrows. The electrochemical parameters obtained from the polarization curves are listed in Table 2, including corrosion potential (Ecorr), passivation current density (ip), passivation potential (Ep), pitting potential (Epit). Ep represents the beginning of the passive potential region; Epit is the potential at which the current density reached a value of 100 μA·cm−2. In Figure 3, the potentiodynamic curves show the TC4 alloy in a state between stable and unstable passivity and parameters about TC4 alloy have not been listed.
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Figure 2. Potentiodynamic polarization curves of the C22 alloy (a), C22 coating (b) in 50 wt.% H2SO4 at 50–70 °C.
Figure 2. Potentiodynamic polarization curves of the C22 alloy (a), C22 coating (b) in 50 wt.% H2SO4 at 50–70 °C.
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As Figure 2 shows, the C22 alloy and C22 coating present typical passivation behavior along with the active region, active–passive transition region, passive region and transpassive region on the forward scan of polarization curves, which is similar to that reported by other researchers for Ni-Cr-Mo alloys in sulfuric acid solutions [11,25,35,45]. Meanwhile, the passive regions for both specimens obtained similar broad potential ranges. At 70 °C, the passivation current density and passivation potential of the C22 coating have not been considered because of the shortage of the passive potential region; a transition from passive to transpassive behavior is observed at more negative pitting potential (−179 mVmse). At 5060 °C, C22 coating presents lower passivation current density (7.55 μA·cm−2 at 50 °C, 1.01 μA·cm−2 at 60 °C) and higher pitting potential (597 mVmse at 50 °C, 593 mVmse at 60 °C), which indicates that C22 coating performs a high anti-corrosion resistance. The OCP values for the C22 alloy and the C22 coating at 50 and 60 °C were in the passive region, while the OCP values at 70 °C were in the active region, meaning that the spontaneous passivation of the C22 alloy and C22 coating was inhibited as the temperature reached 70 °C. During polarization of the C22 alloy at 70 °C and the C22 coating at 60−70 °C, several small peaks arose with higher potential under forward scanning. It could be speculated that the synergistic effect of localized disruption and reformation of the passive film led to this isolated current density fluctuation on the curves. These results indicate that the C22 alloy and coating provide similar corrosion resistance and passive ability in this environment, which could be attributed to the formation of protective oxide passive film [11,12,13,14,18,19,20,21,22,23]; however, these characteristics are significantly degraded at 70 °C. Moreover, the C22 coating was more susceptible to an increase in temperature. On the backward scan, both C22 alloy and coating presented clear repassivation without hysteresis on the polarization curves, accompanied by the significant suppression of the current density. This phenomenon indicates the improved passivity of both specimens after repassivation with no notable pitting initiates, similar to results reported by other researchers [13,22,46]. However, the more positive Ecorr recorded on the backward scan for both specimens suggests that the cathodic kinetics was accelerated compared with the forward scan [19,44]. In contrast, the accelerating effect on cathodic kinetics of the C22 coating at 50 and 60 °C was weaker, as shown by the wider passive region on the backward scan. Meanwhile, the steep but small rises in the passive region indicate that initiating and fixing localized corrosion occurred during the backward scanning process.
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Figure 3. Potentiodynamic polarization curves of the TC4 alloy in 50 wt.% H2SO at 50–70 °C.
Figure 3. Potentiodynamic polarization curves of the TC4 alloy in 50 wt.% H2SO at 50–70 °C.
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As Figure 3 shows, the TC4 alloy presents more notable active-passivation behavior with a much broader active region and active–passive transition region compared with those of the C22 alloy and C22 coating, which was attributed to the presence of highly acidic and oxidizing electrolytes [44,47]. All the passive regions of the TC4 alloy at each temperature were sustained till the potential reached as high as 1.2 Vmse without the occurrence of transpassivation and pitting corrosion, which is in good agreement with the reference results [10,17,41]. Meanwhile, the onset potentials of passive regions remained almost unchanged, and the corrosion potential of the TC4 remained at −1.18 ± 0.01 Vmse without clear temperature effects. In contrast, a visible increase was observed for the current density as the temperature increased, while the distinction between the current densities at 60 and 70 °C was a little more complicated. The current density at 60 °C was slightly lower before passivation and higher after it. This phenomenon has been attributed to the more pronounced effect of temperature on hydrogen evolution reaction and dissolution behavior [10,15,16,17,43]. Meanwhile, isolated fluctuation only occurred for the current density at 70 °C, which indicated that the initiation of localized corrosion demands a higher temperature. It is concluded that the increase in temperature accelerates the dissolution of passive film and materials as well [42]. On the backward scan, TC4 presented significant suppression of the current density, and the degree of suppression decreased with an increase in temperature. Two corrosion potential peaks appeared before the potential scanned back to the initial corrosion potential at each temperature, and the potential interval between the two peaks decreased with an increase in temperature. This behavior is an indication of a high-rate hydrogen evolution reaction, and the increase in temperature facilitated this effect [10,16,17,39,44]. Compared with the C22 alloy and the C22 coating dose, TC4 exhibited a more negative corrosion potential and a higher passive current density at each temperature, which indicates poorer corrosion resistance in this environment. Nevertheless, the much broader passive region and invariable corrosion potential during the backward scan ensured more stable passivation and resistance to localized corrosion for the TC4.

3.3. Potentiostatic Polarization Measurement

As the potentiodynamic polarization curves in Figure 2 and Figure 3 exhibit, the potentials of 100 mVmse, 450 mVmse, and 600 mVmse for potentiostatic polarization are almost all located in the passive region of the three specimens. In contrast, the potential of 600 mVmse for the C22 alloy and C22 coating are in the initial stage of the transpassive region. The potentiostatic polarization curves of the three specimens at different potentials and temperatures are presented in Figure 4 and Figure 5. The average values of current density from potentastic polarization curves for the last 1800 s are listed in Table 3
As Figure 4 shows, the variation trends of current density were similar for the C22 alloy and C22 coating at each potential and temperature. The current density dropped significantly in the initial period, which can be attributed to the formation and evolution of a passive film on the surface [22,39,48]. Moreover, the significant reduction of the duration of the drop with an increase in potential and temperature implies an acceleration effect on the formation of the passive film. As the polarization proceeded, the current density exhibited different variations at different potentials and temperatures. It is clear that a considerable fluctuation of current density arises at potentials of 100 mVmse and 450 mVmse for both the C22 alloy and coating and that the amplitude of fluctuation at 100 mVmse is much larger than that at 450 mVmse. Furthermore, the measured current density at the end of polarization at 100 mVmse is higher than that at 450 mVmse for 60 and 70 °C. With the measured corrosion potential near 100 mVmse that was indicated on the backward scan of the potentiodynamic polarization for two specimens, it could be deduced that a high-rate hydrogen evolution reaction combines with an unstable passive system at 100 mVmse to bring about this phenomenon. The initiation of localized corrosion may arise at 100 Vmse, as shown by the significant fluctuation for the C22 coating at this potential, which will be verified by the results of surface analysis later in this study. The formation and dissolution of the passive film dominated the electrochemical reaction at the 600 mVmse potential, leading to stability at a higher current density. Similarly, the fluctuation observed in the late period for the C22 coating at 600 mVmse could be attributed to the initiation and further passivation of metastable localized corrosion [43]. It can be concluded that the proportion of current density contributed by cathodic reaction and anodic reaction changed with the potential in the passive region. The increase in temperature simply accelerated the electrochemical reaction that manifested as an increase in current density at each potential, which was also reported by other researchers studying the Ni-Cr-Mo alloy [18,19]. The measured current density of the C22 coating was slightly larger and more sensitive to localized corrosion than that of the C22 alloy, which is consistent with the results of the potentiodynamic polarization.
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Figure 4. Current density vs. time plots recorded on C22 alloy (a) and C22 coating (b) during potentiostatic polarization.
Figure 4. Current density vs. time plots recorded on C22 alloy (a) and C22 coating (b) during potentiostatic polarization.
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As Figure 5 shows, the current density of the TC4 alloy was relatively stable at each condition during the potentiostatic polarization measurement. Similar to the C22 alloy and C22 coating, the current density of the TC4 alloy initially underwent a notable drop that increased in amplitude with higher potentials or temperatures. This can be attributed to the formation of a passive film and the high-rate hydrogen evolution reaction that occurred during initial measurements [22,39,43,48,49,50]. The current density remained relatively stable without notable fluctuation as the polarization proceeded, implying stability of the passive system and immunity to localized corrosion. Moreover, the influence of temperature on current density was much more significant than that of potential. It can be concluded that the TC4 alloy yields a stable passive system at each applied potential, and only the increase in temperature accelerates the anode reaction. In comparison with the C22 alloy and C22 coating, the TC4 alloy exhibited higher current density (Table 3) at the same temperature and potential, resulting in a higher corrosion rate that is in good agreement with the results of potentiodynamic polarization and our previous study [36].
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Figure 5. Current density vs. time plots recorded on TC4 alloy during potentiostatic polarization.
Figure 5. Current density vs. time plots recorded on TC4 alloy during potentiostatic polarization.
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Table
Table 3. The average values of the current density of the C22 alloy, C22 coating and TC4 alloy from potentastic polarization curves for the last 1800 s at different temperatures and potentials.
Table 3. The average values of the current density of the C22 alloy, C22 coating and TC4 alloy from potentastic polarization curves for the last 1800 s at different temperatures and potentials.
SampleTemperatureCurrent Density/μA·cm−2
100 mVmse450 mVmse600 mVmse
C22 alloy50 °C0.180.3357.48
60 °C23.820.9686.01
70 °C18.643.32217.44
C22 coating50 °C0.701.8679.50
60 °C4.575.02109.11
70 °C61.3128.3889.14
TC4 alloy50 °C14.4317.4514.59
60 °C74.8054.5550.81
70 °C176.84168.38233.69

3.4. Electrochemical Impedance Spectroscopy Measurement

The measured EIS profiles of the C22 alloy, C22 coating, and TC4 alloy at different potentials and temperatures are presented as the Nyquist and Bode plots in Figure 6, Figure 7, Figure 8 and Figure 9.
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Figure 6. Nyquist plots for the C22 alloy in 50 wt.% H2SO4 at 100 mVmse (a), 450 mVmse (b) and 600 mVmse (c).
Figure 6. Nyquist plots for the C22 alloy in 50 wt.% H2SO4 at 100 mVmse (a), 450 mVmse (b) and 600 mVmse (c).
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Figure 7. Nyquist plots for the C22 coating in 50 wt.% H2SO4 at 100 mVmse (a), 450 mVmse (b) and 600 mVmse (c).
Figure 7. Nyquist plots for the C22 coating in 50 wt.% H2SO4 at 100 mVmse (a), 450 mVmse (b) and 600 mVmse (c).
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Figure 8. Bode plots for the C22 alloy in 50 wt.% H2SO4 at 100 mVmse (a), 450 mVmse (b) and 600 mVmse (c).
Figure 8. Bode plots for the C22 alloy in 50 wt.% H2SO4 at 100 mVmse (a), 450 mVmse (b) and 600 mVmse (c).
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Figure 9. Bode plots for the C22 coating in 50 wt.% H2SO4 at 100 mVmse (a), 450 mVmse (b) and 600 mVmse (c).
Figure 9. Bode plots for the C22 coating in 50 wt.% H2SO4 at 100 mVmse (a), 450 mVmse (b) and 600 mVmse (c).
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The depressed capacitive arcs in the high-frequency regions of Nyquist plots (Figure 6 and Figure 7) correspond to the charge transfer on the surface [35,51]. The near −80° phase angles for both specimens with slight fluctuations over a wide range of frequencies in Bode plots (Figure 8 and Figure 9) indicate the pure capacitive response of stable passive films and the existence of more than one time constant, which is related to the dielectric properties and bilayer of the passive film. [40] It was observed that the diameter of the capacitive arc and peak of phase angle for both specimens roughly decreased with an increase in temperature at each potential, which suggests a localized disruption or thinning of the passive film [50,51]. The exceptional variation for the C22 alloy at 600 mVmse and the C22 coating at 100 mVmse, both at 60 °C, coincides with the measured instantaneous current density during potentiodynamic polarization. It could be speculated that the passive film of both specimens was exposed to a destabilizing impact at 50 and 60 °C, and significantly degraded at 70 °C. The straight line in the low-frequency region of Nyquist plots for the C22 alloy at 100 mVmse, 70 °C and for the C22 coating at 100 mVmse is related to the diffusion of the oxidation products [24,35,51,52,53]. The transformation from a straight line to an incomplete second capacitive arc with an increase in applied potential is related to the charge transfer at the disruption sites of the passive film. [45,51,54] The presence of inductance loops along with the significant reduction of capacitive arcs, phase angle, and impendence modulus in the transpassive region (600 mVmse) is associated with the relaxation process of adsorption species which come from the dissolution and disruption of passive film. [45,54] It was concluded that the controlling behavior of passive film dissolution of both specimens was affected by the applied potential and temperature, which primarily resulted from the dissolution and disruption of passive film. Most of the capacitive arc diameters for the C22 alloy were larger than that of the C22 coating at the same temperature and potential, except for those at 600 mVmse and 50 and 70 °C. This suggests that the passive film of the C22 alloy exhibited higher corrosion resistance within the passive region, which is consistent with the difference between the measured current density of the two specimens during potentiostatic polarization. The influence of the composition and structure of the passive film for the two specimens will be revealed through subsequent surface analysis.
As Figure 10 and Figure 11 show, the Nyquist and Bode plots of the TC4 maintain a similar shape as the applied potential and temperature vary, indicating unaltered electrochemical behavior and interface characteristics. It is comprehensible that each applied condition was covered by the passive region measured during potentiodynamic polarization. The Nyquist plots of the TC4 alloy consist of a depressed capacitive arc followed by an incomplete second capacitive arc, indicating that the electrochemical behavior of the TC4 alloy was mainly controlled by charge transfer. The depressed capacitive arc and the broad peak of the phase angle imply that the predominantly capacitive behavior of the compact passive film combines with the reaction products absorbing on the surface. This result is different from the time constant measured by Mogoda in sulfuric acid with a lower concentration [42], but is similar to the results reported by other studies of titanium under more aggressive conditions [16,17,55,56]. It could be attributed to the larger amount of absorption reaction products resulting from a more concentrated sulfuric acid. The sharp reductions in the capacitive arc, impedance modulus, and the phase angles as the temperature increases to 70 °C indicate the significant thinning or localized disruption of the passive film [16,50,51,54,56]. As the applied potential increased, the diameter of the capacitive arc and impedance modulus at 50 and 70 °C basically remained unchanged as the applied potential increased, while those at 60 °C slightly increased. It suggests the passive film performs higher resistance against disruption and dissolution at 60 °C, which is consistent with the decrease in current density as the applied potential increased during potentiostatic polarization at 60 °C. This stability within the passive region is inconsistent with the results reported by Wang on TC4 in 0.5 mol/L sulfuric acid [54], which could be attributed to the more rapid formation of a compact passive film in concentrated sulfuric acid. The incomplete second capacitive arc is related to the charge transfer on the surface of metal under the passive film. Meanwhile, the phase angles at a low frequency (approximately 10−1 Hz) at 450 mVmse and 600 mVmse, both at 70 °C, are clearly higher than those at lower temperatures, implying that very little reaction product absorption occurred at 70 °C. It could be concluded that the corrosion resistance of the TC4 alloy remains stable in the passive region but significantly degrades with an increase in temperature. Similarly, the influence of the composition and structure of passive films will be revealed in subsequent surface analysis. Judging from the EIS results, the corrosion resistance of the TC4 alloy is clearly lower than that of the C22 alloy and C22 coating, which is consistent with the results of the electrochemical measurements.
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Figure 10. Nyquist plots for the TC4 alloy in 50 wt.% H2SO4 at 100 mVmse (a), 450 mVmse (b) and 600 mVmse (c).
Figure 10. Nyquist plots for the TC4 alloy in 50 wt.% H2SO4 at 100 mVmse (a), 450 mVmse (b) and 600 mVmse (c).
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Figure 11. Bode plots for the TC4 alloy at 100 mVmse (a), 450 mVmse (b) and 600 mVmse (c).
Figure 11. Bode plots for the TC4 alloy at 100 mVmse (a), 450 mVmse (b) and 600 mVmse (c).
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3.5. Surface Morphology

SEM-equipped EDS was used to observe the corroded surface morphologies of the C22 alloy, C22 coating, and TC4 alloy after 10 h of potentiostatic polarization at each applied potential and temperature. There was no occurrence of clearly corroded morphology on the surface of the C22 alloy in the passivation region (100 mVmse to 450 mVmse) and only the surface morphologies in the transpassive region (600 mVmse) are presented in Figure 12. The surface morphologies and elemental content analysis of the C22 coating and the TC4 alloy at each condition are presented in Figure 13, Figure 14 and Figure 15.
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Figure 12. Corroded morphologies of the C22 alloy after potentiostatic polarization at 600 mVmse 50 °C (a), 600 mVmse 60 °C (b) and 600 mVmse 70 °C (c).
Figure 12. Corroded morphologies of the C22 alloy after potentiostatic polarization at 600 mVmse 50 °C (a), 600 mVmse 60 °C (b) and 600 mVmse 70 °C (c).
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Figure 13. Corroded morphologies of the C22 coating after potentiostatic polarization at 100 mVmse 50 °C (a), 100 mVmse 60 °C (b), 100 mVmse 70 °C (c), 450 mVmse 50 °C (d), 450 mVmse 60 °C (e), 450 mVmse 70 °C (f), 600 mVmse 50 °C (g), 600 mVmse 60 °C (h) and 600 mVmse 70 °C (i).
Figure 13. Corroded morphologies of the C22 coating after potentiostatic polarization at 100 mVmse 50 °C (a), 100 mVmse 60 °C (b), 100 mVmse 70 °C (c), 450 mVmse 50 °C (d), 450 mVmse 60 °C (e), 450 mVmse 70 °C (f), 600 mVmse 50 °C (g), 600 mVmse 60 °C (h) and 600 mVmse 70 °C (i).
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As Figure 12 and Figure 13 show, both C22 alloy and C22 coating exhibit intergranular corrosion behavior. The corroded section of the C22 alloy is the grain boundary and contrasts with that of the C22 coating, which is eutectics in a primary grain boundary. Similar intergranular corrosion has been reported by other published studies pertaining to Ni-based alloys and coatings in environments containing sulfuric acid [14,50]. However, the material loss shown by the corroded morphology at each condition indicates that the C22 alloy exhibits better corrosion resistance than the C22 coating. In addition, the influence of temperature and the applied potential on the corroded morphology of two specimens is consistent with the results of electrochemical measurements. In observing the variation in the corroded morphologies of the C22 coatings with applied potential and temperature increases, it is speculated that the corrosion was initiated at the locations of micropores or defects in the eutectics and then attacked the surrounding eutectics, which is in agreement with the speculations of our previous study [36]. Within the passivation region (100 mVmse−450 mVmse) shown in Figure 13a–f, there was no significant variation in the corroded morphology of the C22 coating, but only a slight increase of the corroded micropores and depth with increases in the applied potential and temperature. The increase in temperature within the transpassive region aggravates the dissolution of the C22 alloy (Figure 12), which is revealed by deeper and more dissolution sites in the grain interiors; however, it has no significant impact on the corrosion of the C22 coating (Figure 13g–i). Please note the nanoscale pitting exhibited in the corroded morphologies of the C22 coating, which contradicts the measured results of potentiodynamic and potentiostatic polarization.
Due to the homogeneous austenite of the C22 alloy, the elemental content of the corroded surface presents no clear fluctuation and maintains a composition that is almost the same as the original. Figure 14 presents the elemental content analysis for different sites of the C22 coating that possess the representative corroded morphology and experienced potentiostatic polarization at 600 mVmse and 70 °C; the elemental contents are listed in Table 4. It could be seen that the content of the Mo element in the corroded eutectics (Point 1) is similar to that in the primary grain (Point 2), which contradicts the segregation of the Mo element in the eutectics of the prepared C22 coating [40]. It is speculated that the eutectics containing a higher Mo element content were dissolved to expose the internal primary grain. The black pitting (Point 3) that contains a much lower content of the Mo element may have been caused by the overlapping position of external and internal eutectics. In contrast, the content of other elements remained nearly unchanged at the three measured points. It could be deduced that the eutectics of the C22 coating with a higher Mo element content were preferentially corroded. Moreover, a similar content of O element was detected at different sites of both the corroded C22 alloy and coating, indicating that the passive film in the form of oxide covered the whole surface of the two specimens. The nanoscale and existence of passive film effectively suppressed the active dissolution within the formed pits of the C22 coating, explaining why no notable pitting behavior was detected during electrochemical measurements.
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Figure 14. SEM morphology of the C22 coating after potentiostatic polarization at 600 mVmse 70 °C.
Figure 14. SEM morphology of the C22 coating after potentiostatic polarization at 600 mVmse 70 °C.
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The corroded morphologies of the TC4 alloy at different applied potentials and temperatures are shown in Figure 15. Figure 16 shows the elemental content analyses for different sites of uncorroded and corroded specimens at 450 mVmse and 70 °C. It can be seen that clear, localized dissolution occurred and was gradually aggravated by the increase in applied potential and temperature. The influence of the applied potential and temperature on the corroded morphology is consistent with the measured current density during potentiostatic polarization (Figure 5). By comparing the element content (Table 5) of the uncorroded and corroded specimens shown in Figure 16, it is apparent that the localized corrosion was caused by the preferential dissolution of a bright β phase that contains more V element (Point 2 in Figure 16a) [57,58]. The decrease in the V element content at the localized corrosion site (Point 4 in Figure 16b) indicates the subsequent exposure of an internal α phase. It is notable that the clear dissolution of an α phase adjacent to a β phase occurred as the temperature rose to 60 °C. Moreover, the passive film in the form of oxide covered the surface, which is inferred from the existence of the O element at different sites (Point 3 and Point 4 in Figure 16b) of the corroded morphology. Judging by the corroded morphology, the corrosion resistance of the TC4 alloy was clearly worse than that of the C22 alloy and coating.
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Figure 15. Corroded morphologies of the TC4 alloy after potentiostatic polarization at 100 mVmse 50 °C (a), 100 mVmse 60 °C (b), 100 mVmse 70 °C (c), 450 mVmse 50 °C (d), 450 mVmse 60 °C (e), 450 mVmse 70 °C (f), 600 mVmse 50 °C (g), 600 mVmse 60 °C (h) and 600 mVmse 70 °C (i).
Figure 15. Corroded morphologies of the TC4 alloy after potentiostatic polarization at 100 mVmse 50 °C (a), 100 mVmse 60 °C (b), 100 mVmse 70 °C (c), 450 mVmse 50 °C (d), 450 mVmse 60 °C (e), 450 mVmse 70 °C (f), 600 mVmse 50 °C (g), 600 mVmse 60 °C (h) and 600 mVmse 70 °C (i).
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Figure 16. SEM morphologies of the TC4 alloy before (a) and after (b) potentiostatic polarization at 450 mVmse 70 °C.
Figure 16. SEM morphologies of the TC4 alloy before (a) and after (b) potentiostatic polarization at 450 mVmse 70 °C.
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Table
Table 5. EDS analysis results of the TC4 alloy.
Table 5. EDS analysis results of the TC4 alloy.
OAlTiV
Point 1wt%007.7389.9902.28
at%012.9785.0102.02
Point 2wt%006.2786.2107.52
at%010.6682.5706.78
Point 3wt%07.3707.4482.4102.78
at%18.3510.9868.5002.17
Point 4wt%08.0106.4482.9602.58
at%19.8609.4668.6702.01

3.6. Composition of Passive Film

The passive film of the C22 alloy, C22 coating and TC4 alloy formed by potentiostatic polarization at different applied potentials and temperatures were characterized through XPS. The analyses were performed according to the binding energy published in the NIST X-ray Photoelectron Spectroscopy Database. High-resolution spectra of Ni 2p3/2, Cr 2p3/2 and Mo 3d for the C22 alloy and C22 coating were registered to determine the chemical state, while Ti 2p and Al 2p were used for the TC4 alloy. The curve fits for the high-resolution spectra of the three specimens in different polarization conditions are presented in Figure 17, Figure 18 and Figure 19. The influence of applied potential on the composition of the passive film is illustrated by the spectra at 60 °C, whereas that of the temperature is illustrated by the spectra at 450 mVmse.
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Figure 17. High-resolution XPS profiles for Ni 2p3/2 (a), Cr 2p3/2 (b) and Mo 3d (c) of the surface of the C22 alloy after potentiostatic polarization.
Figure 17. High-resolution XPS profiles for Ni 2p3/2 (a), Cr 2p3/2 (b) and Mo 3d (c) of the surface of the C22 alloy after potentiostatic polarization.
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Figure 18. High-resolution XPS profiles for Ni 2p3/2 (a), Cr 2p3/2 (b) and Mo 3d (c) of the surface of the C22 coating after potentiostatic polarization.
Figure 18. High-resolution XPS profiles for Ni 2p3/2 (a), Cr 2p3/2 (b) and Mo 3d (c) of the surface of the C22 coating after potentiostatic polarization.
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Similar to the previous sections, the surface compositions of the C22 alloy and C22 coating are described in comparison with one another. The surface composition content determined by the survey spectra of both specimens (not shown here) exhibits a similar descending order of Cr > Mo > Ni, except for C and O, with a depletion of Ni compared to the concentration of the interior. There is no considerable difference in the composition content between the two specimens. As Figure 17 and Figure 18 show, the passive films of the C22 alloy and C22 coating possess similar composition consisting primarily of oxides and hydroxides of Ni, Cr, and Mo—specifically, NiO, Cr2O3, CrO3, MoO2, MoO3, Ni(OH)2 and Cr(OH)3. Specifically, Ni(OH)2, Cr(OH)3 and MoO3 were the primary oxidized species of Ni, Cr, and Mo elements, respectively. These observations are in agreement with the results reported by other researchers on alloys containing these elements in similarly aggressive environments [19,22,24,38,48,59,60,61,62].
It can be seen that there was a clear increase in the hydroxide content occurring on both two specimens as the applied potentials and temperatures increased, with an especially significant increase as the potential entered into the transpassive region. This indicates that the transformation from dense oxide to porous hydroxide depends on the applied potential and temperature. In terms of the relative content of metal, there was a marked and slight decrease for both specimens with the respective increase in applied potential and temperature. This implies that the thickness of the passive film depends more on the potential than on the temperature. The similar passive film composition yielded different corrosion resistance for the C22 alloy and coating. This contradiction will be discussed in the following section.
The surface composition of the TC4 alloy after polarization in each condition as determined by the survey spectra (not shown here) contained mostly C, O, Ti and Al elements, but no V element. As the applied potential and temperature increased, the ratio of Ti/Al slightly increased and remained nearly constant, respectively, which could be attributed to the conversion of Ti suboxide to ideal oxide facilitated by the increase in the applied potential. As the high-resolution spectra of Ti 2p and Al 2p (Figure 19) show, the composition of the passive film consisted of oxides of Ti and Al elements, specifically predominant amounts of Ti2O3 and TiO2, and a small quantity of Al2O3, which has been observed and reported in other research on passive films of titanium alloys in different environments [16,43,48,55,57]. It could be seen that the peak of these oxides remained nearly unchanged with the applied potential and temperature, whereas the peaks of the metallic states of Ti and Al maintained a low intensity and nearly vanished as the potential and temperature increased, indicating that the passive film maintained a greater thickness compared to those of the C22 alloy and coating.
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Figure 19. High-resolution XPS profiles for Ti 2p (a) and Al 2p (b) of the surface of the TC4 alloy after potentiostatic polarization.
Figure 19. High-resolution XPS profiles for Ti 2p (a) and Al 2p (b) of the surface of the TC4 alloy after potentiostatic polarization.
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4. Discussion

As the results of the OCP and potentiostatic polarization tests indicate, a remarkable transformation occurred in the air-formed oxide film of both the C22 alloy and C22 coating in 50 wt.% H2SO4. This was indicated by the positive shift of the measured OCP and the drop of the current density in the initial two tests, respectively. In addition, the more sharply increase in OCP of the C22 coating in the early stage indicates the less protective air-formed oxide film in comparison to the C22 alloy. This electrochemical phenomenon has been frequently recognized in other research on alloys possessing passivation features and resulted from the improvement or thickening of a previous oxide film [38,39,40,54]. Compared to the EDS element content analysis on the as-prepared surfaces in our previous study [36], a uniformly distributed O element was detected on the corroded surfaces of both the C22 alloy and laser-cladding C22 coating in this study, confirming the formation of an intact and thicker oxide passive film. Due to the reaction with the electrolyte, another variation in composition contributed to the formation of the passive film. The air-formed oxide film on Ni-Cr-Mo alloys was found to primarily consist of NiO, Cr2O3, CrO3, MoO2 and MoO3 in other published literature [18,24,61]. By comparison, the results of the XPS analysis in this study indicated that additional Ni(OH)2 and Cr(OH)3 existed in the passive film of both the C22 alloy and laser-cladding C22 coating after polarization. The Ni depletion occurred in the passive film, and the hydroxides along with MoO3 predominated the composition of the respective elements. According to the observations of Lloyd et al. [19,22,48,62,63], the passive film of the C22 alloy and coating in 50 wt.% H2SO4 solution exhibited a bilayer structure consisting of a Cr2O3-dominated compact inner layer and a porous outer layer containing oxides of Mo and hydroxides of Ni and Cr. The two capacitive arcs of the Nyquist plots and multiple time constants observed in the EIS measurement confirm this bilayer structure [35,40,50,62].
The similar composition and structure of a passive film implies that the C22 alloy and C22 coating exhibited similar passivation behaviors in 50 wt.% H2SO4, which was confirmed by the potentiodynamic polarization and EIS measurements. Besides the broad passive region, the phenomenon of improved passivity after repassivation without initiation through pitting was observed for both specimens [13,22,46]. With an increase in the applied potential and temperature within the passive region, the controlling behavior of the passive film dissolution was transformed from a combination of charge transfer and diffusion processes to that of two-charge transfer processes but maintained a combination of charge transfer and intermediate product relaxation processes in the transpassive region. In combination with a change in passive film composition with applied potential and temperature, it is concluded that the transformation of controlling behavior is associated with the formation of a porous outer layer and the acceleration of the transformation from a dense oxide to a porous hydroxide is the predominant behavior of degradation [48].
The higher corrosion tendency of the C22 coating at each temperature is demonstrated by the more negative OCP, indicating that the C22 coating exhibits poorer thermodynamic stability compared to the C22 alloy. In terms of corrosion kinetics, the C22 coating also exhibited a faster corrosion rate according to the higher current density and lower impedance during polarization and EIS measurements. In addition, the more notable fluctuation observed for measured current density during polarization indicates that C22 coating is more sensitive to localized corrosion under the same conditions. The relatively poor corrosion resistance of the C22 coating could be interpreted by the characteristics of the passive film and microstructure revealed in the surface analysis. In spite of the similar composition of the passive film analyzed by XPS, considerable Mo segregation was observed by SEM within the eutectics of the C22 coating, which results in the microscale heterogeneity of the passive film. This heterogeneity leads to the preferential dissolution and fracture of the passive film and presents as the degradation of corrosion resistance and the intergranular corrosion on eutectics that contain less Cr and Ni. As the SEM results suggest, the micropores or defects within the eutectics of the C22 coating facilitate localized corrosion as initiation sites; meanwhile, the Mo segregation in the eutectics results in a slower reformation of the passive film in the adjacent low-Mo sites [18,19,46]. The synergistic effect of these two behaviors makes C22 coating more susceptible to disruption from localized corrosion.
Similar passivation behavior is exhibited by the TC4 alloy in 50 wt.% H2SO4 at 50 to 70 °C. As the results of the OCP and XPS analysis show, no clear composition transformation occurs on the air-formed oxide film, but there is probable thickening after immersion into the electrolyte. The compact and thick passive film, consisting of TiO2, Ti2O3, and Al2O3, provides the stable passivation to the TC4 alloy in this aggressive environment [15,16,17,42,44,48,58] demonstrated by the very broad passive region, stable current density and suppression of current density during polarization tests. The stable OCP that localizes in the action region demonstrates the spontaneous active dissolution behavior of a passive film on the TC4 alloy under testing conditions [10,16,41,42,48]. The three corrosion potential peaks that appeared during the potential scanning back were an indication of a high-rate hydrogen evolution reaction; simultaneously, the increase in temperature facilitated this effect [10,16,17,39,44]. Meanwhile, the electrochemical behavior of the TC4 alloy was controlled by the charge transfer and the diffusion processes under each testing condition, as demonstrated by the EIS measurement [17]. The corrosion behavior revealed by the SEM observation is the preferential dissolution of a low-content β phase, which contradicts the observations of Chen et al. [57]. Moreover, the corrosion tendency and rate of the TC4 alloy was just slightly promoted by the increase in applied potential and temperature, indicating that no clear transformation occurs on the passivation behavior and passive film [51]. In spite of possessing more stable passivity, the TC4 alloy still exhibited less corrosion resistance based on the results of electrochemical testing and active dissolution behavior compared to the C22 alloy and C22 coating.

5. Conclusions

The electrochemical behavior and passive film properties of Hastelloy C22 alloy, laser-cladding C22 coating and Ti–6Al–4V alloy (TC4) in a simulated sulfuric acid dew-point corrosion environment (50 wt.% H2SO4 solution at 50 to 70 °C) were investigated by OCP, potentiodynamic polarization, potentiostatic polarization, EIS, SEM, and XPS analyses. The main conclusions obtained from this study are presented below:
(1)
Both the laser-cladding C22 coating and C22 alloy exhibit similar passivation and repassivation behavior without pitting corrosion, while TC4 alloy exhibits passivation behavior with a much broader passive region, and a more stable current density. An increase in temperature aggravates the corrosion of three materials by degrading the stability of passive film and accelerating the charge transfer process.
(2)
The passive film formed on the laser-cladding C22 coating possessed a similar bilayer structure in comparison with C22 alloy, consisting of a Cr2O3-dominated compact inner layer and a porous outer layer containing oxides of Mo and hydroxides of Ni and Cr. The degradation behavior of passive film of the C22 coating and C22 alloy performs as the formation of a porous outer layer and the transformation from the dense oxide to a porous hydroxide. The passive film of the TC4 alloy possesses a compact single-layer structure consisting of TiO2, Ti2O3 and Al2O3.
(3)
The slightly poorer corrosion resistance and higher sensitivity to localized corrosion of laser-cladding C22 coating than C22 alloy is attributed to the element segregation in microstructure introduced by the laser-cladding process and the consequent microscale heterogeneity of the formed passive film. The poorer corrosion resistance of the TC4 alloy in comparison to the C22 coating and C22 alloy is attributed to spontaneous active dissolution behavior of the β phase and less protective passive film.
(4)
The laser-cladding C22 coating on steel can be considered as a potential method to solve the sulfuric acid dew-point corrosion problem in a flue gas environment because of the comparable and better corrosion resistance in comparison with the completely using of the C22 alloy and TC4 alloy, respectively.

Author Contributions

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

Funding

This research received no external funding.

Institutional Review Board Statement

Not applicable.

Informed Consent Statement

Not applicable.

Data Availability Statement

Not applicable.

Conflicts of Interest

The authors declare no conflict of interest.

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Table
Table 1. Chemical compositions of the C22 powder, C22 alloy, and TC4 alloy in wt.%.
Table 1. Chemical compositions of the C22 powder, C22 alloy, and TC4 alloy in wt.%.
MaterialNiCrMoMnFeSiCoVWTiAlCuSbC
C22 powderBal21.313.2-2.93-2.0-3.0----0.08
C22 alloyBal.20.013.80.455.00.081.830.33.2----0.001
TC4 alloy----0.30.15-3.8-Bal.6.1---
Table
Table 2. Electrochemical parameters of the C22 alloy and C22 coating at different temperatures.
Table 2. Electrochemical parameters of the C22 alloy and C22 coating at different temperatures.
SampleTemperatureEcorr (mVmse)ip (μA·cm−2)Ep (mVmse)Epit (mVmse)
C22 alloy50 °C−64524.55−458588
60 °C−5956.31−412591
70 °C−5928.51−381613
C22 coating50 °C−5917.55−332597
60 °C−6561.01−434593
70 °C−594--−179
Table
Table 4. EDS analysis results of the C22 coating.
Table 4. EDS analysis results of the C22 coating.
OWMoCrFeNiCo
Point 1wt%03.6206.2912.5819.7404.7652.310.70
at%12.8401.9407.4521.5904.850.6750.65
Point 2wt%03.2005.3112.2421.1505.0053.090
at%11.3901.6407.2623.1505.1051.460
Point 3wt%02.0604.1108.7321.3505.3358.430
at%07.3901.2805.2223.5505.4757.090
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Zheng, C.; Liu, Z.; Liu, Q.; Kong, Y.; Guo, S.; Liu, C. Electrochemical Behavior and Passive Film Properties of Hastelloy C22 Alloy, Laser-Cladding C22 Coating, and Ti–6Al–4V Alloy in Sulfuric Acid Dew-Point Corrosion Environment. Metals 2022, 12, 683. https://doi.org/10.3390/met12040683

AMA Style

Zheng C, Liu Z, Liu Q, Kong Y, Guo S, Liu C. Electrochemical Behavior and Passive Film Properties of Hastelloy C22 Alloy, Laser-Cladding C22 Coating, and Ti–6Al–4V Alloy in Sulfuric Acid Dew-Point Corrosion Environment. Metals. 2022; 12(4):683. https://doi.org/10.3390/met12040683

Chicago/Turabian Style

Zheng, Chao, Zongde Liu, Quanbing Liu, Yao Kong, Shengyang Guo, and Congcong Liu. 2022. "Electrochemical Behavior and Passive Film Properties of Hastelloy C22 Alloy, Laser-Cladding C22 Coating, and Ti–6Al–4V Alloy in Sulfuric Acid Dew-Point Corrosion Environment" Metals 12, no. 4: 683. https://doi.org/10.3390/met12040683

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