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Article

Comparison of Corrosion Behavior of a-C Coatings Deposited by Cathode Vacuum Arc and Filter Cathode Vacuum Arc Techniques

by
Zhiqing Feng
1,
Zhetong Zhou
2,
Junhao Zeng
2,
Ding Chen
2,
Fengying Luo
3,
Qimin Wang
2,
Wei Dai
2,* and
Ruiming Zhang
1,3,*
1
Guangdong Hydrogen Energy Institute of Wuhan University of Technology, Foshan 528200, China
2
School of Electromechanical Engineering, Guangdong University of Technology, Guangzhou 510006, China
3
National Energy Key Laboratory for New Hydrogen-Ammonia Energy Technologies, Foshan Xianhu Laboratory, Foshan 528200, China
*
Authors to whom correspondence should be addressed.
Coatings 2024, 14(8), 1053; https://doi.org/10.3390/coatings14081053
Submission received: 21 July 2024 / Revised: 9 August 2024 / Accepted: 15 August 2024 / Published: 17 August 2024
(This article belongs to the Special Issue Anti-corrosion and Anti-wear Coatings: Fundamentals, Technologies, and Applications)
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Versions Notes

Abstract

:
This study explores the utilization of cathodic vacuum arc (CVA) technology to address the limitations of magnetron sputtering technology in preparing amorphous carbon (a-C) coatings, such as having a low ionization rate, low deposition rate, and insufficiently dense structure. Specifically, a-C coatings were prepared by the cathodic vacuum arc (CVA)and the filtered cathodic vacuum arc (FCVA) technology,, one with embedded carbon particles and one without, both having closely related carbon structures. Research is currently underway on bipolar plate coatings for fuel cells. The corrosion behavior of the prepared a-C coatings was examined through Tafel polarization analysis under simulated fuel cell operating conditions as well as potentiostatic analysis at 0.6 V under normal conditions and 1.6 V under start–stop conditions for 7200 s. The coatings before and after corrosion are characterized using scanning electron microscopy, energy-dispersive X-ray spectroscopy, Raman spectroscopy, and infrared spectroscopy. The results reveal that the incorporation of conductive graphite-like particles in the coatings reduces their contact resistance. However, the gaps between these particles and the coatings act as pathways for corrosive solution, exacerbating the corrosion of the coatings. After corrosion at 0.6 V, both sets of coatings with sp2-hybridized carbon structures are contaminated by elements such as hydrogen and oxygen, leading to an increase in their contact resistance. Under high potential conditions (1.6 V), large corrosion pits and defects appear at the locations of graphite-like carbon particles. Furthermore, both sets of samples exhibit more severe oxygen contamination and a transformation of broken carbon bonds from sp3- to sp2-hybridized forms, irrespective of whether embedded graphite particles are present.

1. Introduction

Proton exchange membrane fuel cells (PEMFCs) exhibit immense promise as a next-generation energy conversion device with high efficiency and minimal environmental impact [1,2]. Bipolar plates, as the core components of PEMFCs, support the membrane electrodes, manage reaction-generated water and heat, and segregate the reaction gasses within the stack [3]. Compared to graphite and its composite bipolar plates, metal bipolar plates are economical, have high mechanical strength, and can be easily processed into ultra-thin sheets. However, in the acidic working environment of PEMFCs, they develop passivation films on the surface, which increase the contact resistance and release metal ions that contaminate the catalyst, thereby reducing the output performance of the battery [4,5,6].
The advancements in physical vapor deposition (PVD) technology have made it possible to address these challenges by applying protective coatings like transition metal nitrides, carbides, MAX phases, and amorphous carbon (a-C) on the surface of metal bipolar plates [7,8,9,10]. For instance, Lee et al. [11] coated 316 L stainless steel with CrN and TiN using cathodic arc evaporation technology, achieving contact resistances of 23 mΩ·cm2 and 10 mΩ·cm2, respectively, under a pressure of 1.5 MPa. The self-corrosion current densities for CrN and TiN coatings measured in a simulated fuel cell environment were 0.3 μA/cm2 and 40 μA/cm2, respectively. Although the CrN coating exhibited good corrosion resistance with a metastable structure, the TiN coating showed delamination. Forouzanmehr et al. [10] also applied CrN and TiN coatings on the surface of 410 stainless steel using cathodic arc evaporation. The corrosion current densities of CrN and TiN were 0.07 μA/cm2 and 0.11 μA/cm2, respectively, under standard corrosion solution conditions. The TiN coating did not exhibit extensive delamination, and the interfacial contact resistances (ICRs) were 9.26 mΩ·cm2 and 11.23 mΩ·cm2 at a pressure of 1.4 MPa. The instability of CrN and TiN coatings with variations in the fabrication process parameters was attributed to defects such as droplets, pores, and cracks in the single-layer coating structure.
Metal carbides are recognized for their excellent wear and corrosion resistance [12,13]. Su et al. [14] applied CrCx coatings with varying carbon content on the surface of 316 L stainless steel substrate using unbalanced magnetron sputtering. Optimal electrical conductivity and corrosion resistance were achieved at a 52% atomic ratio of carbon, with a corrosion current density of 0.07 μA/cm2 and a contact resistance of 7 mΩ·cm2 at 1.4 MPa under simulated conditions. These findings were attributed to the formation of conductive sp2-hybridized carbon structures within the coating. Lu et al. [9] utilized DC magnetron sputtering to deposit Ti3AlC2 (MAX phase) coatings via heat treatment at 800 °C and 900 °C. The self-corrosion current densities for unheated, 800 °C heat-treated, and 900 °C heat-treated coatings were measured to be 2.1, 0.47, and 4.4 μA/cm2, respectively. The performance of the samples before heat treatment did not meet the requirements for practical application. However, the corrosion resistance and conductivity of Ti3AlC2 were enhanced after heat treatment.
Owing to their excellent electrical conductivity, chemical stability, and cost-effectiveness, a-C coatings are widely used in the modification of bipolar plates [15,16]. Currently, the a-C coatings are mainly prepared via chemical vapor deposition (CVD) or PVD. The CVD technology, which uses carbon-containing gasses as the carbon source, inevitably introduces impurity elements. By contrast, the PVD technology uses high-purity graphite as the carbon source and can control the proportion of sp3/sp2-hybridized carbon structures in the coating by adjusting the deposition bias, pressure, and temperature [17,18,19,20]. Several researchers have successfully prepared a-C coatings using magnetron sputtering technology with self-corrosion current densities below 1 μA/cm2 and an ICR below 10 mΩ·cm2 at 1.4 MPa. Furthermore, it has been reported that there is a balance between the conductivity and corrosion resistance of the coating, and the conductivity of the coating has a great relationship with the proportion of sp3 hybrid carbon structure in the coating, and with the increase in the proportion of sp3 hybrid carbon structure, the corrosion resistance of the coating is enhanced, but its conductivity is reduced [10,16,21,22]. This is mainly because sp3-hybridized carbon structures have higher bond energy and stronger chemical stability [19,20]. A higher proportion of sp3-hybridized carbon structures contributes to the corrosion resistance of the coating. However, the electrical conductivity of the a-C coating mainly benefits from the sp2-hybridized graphite clusters inside the coating. As the proportion of sp2-hybridized carbon structures decreases, the contact resistance of the coating increases [23,24]. Furthermore, the preparation of a-C coatings through magnetron sputtering technology faces challenges such as cross-sectional columnar structures, crack defects, a low ionization rate, and a low deposition rate. In addition, most of the constant potential tests on protective coatings for bipolar plates are typically based on the normal fuel cell operating potential of 0.6 V. However, there is a lack of testing at high corrosion potentials in the existing studies.
Therefore, in this study, the cathodic vacuum arc (CVA) process, renowned for its high deposition and ionization rates, is utilized to prepare a dense glass-like structure on the cross-section of the a-C coating. By adjusting the bias during deposition, the CVA process allows precise control over the proportion of sp3/sp2-hybridized carbon structures inside the a-C coating. However, neutral graphite-like carbon particles are inevitably embedded in the coating during the CVA process. This necessitates exploring the impact of these particles on the electrical conductivity and corrosion resistance of the coating. To address this issue, the filtered cathodic vacuum arc (FCVA) technology is employed. The FCVA technology relies on magnetic filter bending tubes to remove neutral carbon particles, thereby achieving a-C coatings with a smoother surface. The prepared coatings are subjected to corrosion resistance tests under extreme conditions, simulating the high operating potential of PEMFCs [18,20]. The contact resistance, carbon structure, and surface morphology before and after corrosion are characterized to explore the impact of graphite-like particles on the electrical conductivity and corrosion resistance of the a-C coating.

2. Materials and Methods

2.1. Material Preparation

The sputtering source materials used were a rectangular titanium target (99.95% purity) and a round stonewashed target (99.5% purity). The substrates used in this study primarily included commercial 304 L stainless steel squares (20.0 × 20.0 × 1 mm3) for contact resistance and corrosion resistance tests and P-type single-crystal silicon wafers (oriented in the 100 direction, with a thickness of 0.625 μm) for testing the coating cross-section morphology and residual stress. Prior to deposition, the substrates underwent sequential treatment: ultrasonic cleaning with a metal cleaner and rinsing with anhydrous ethanol, followed by immediate drying in a drying oven.

2.2. Deposition of Coating

The a-C coatings were prepared using CVA and FCVA technologies, employing a uniform flexible metal Ti layer as the transition layer. The same deposition process was utilized for the transition layer. When the vacuum in the chamber dropped below 5 × 10−3 Pa, high-purity Ar gas was introduced for glow cleaning and ion source etching. Subsequently, the Ti transition layer was prepared by controlling the magnetron power at 2 kW, bias at −100 V, and pressure at 0.5 Pa.
During the deposition of a-C coatings using CVA and FCVA processes, the same high-purity graphite target was used, with identical control over the arc source current. There were differences between the processes laid out in the deposition pressure and bias conditions. The specific deposition parameters are listed in Table 1. In the FCVA deposition system, Ar was not introduced, and due to the presence of the filtering arc bending tube, the distance between the target and substrate was larger. To ensure sufficient energy reached the substrate ions, a higher bias was applied. The samples prepared with the CVA process containing carbon particle inclusions were denoted as T-1, while the samples prepared with the FCVA process yielding smooth surfaces were labeled as T-2.

2.3. Coating Characterization

The mechanical properties of the coatings were characterized using nanoindentation (TTX-NHT2 nanoindenter, Anton Paar, Glaz, Australia) and residual stress testing (FST-1000, Supro Instruments, Shenzhen, China). The nanoindentation tests were conducted at a constant load of 5 mN, with each test repeated at nine points to ensure accuracy. The surface and cross-sectional morphologies of the a-C coating samples, both before and after corrosion, were examined using scanning electron microscopy (SEM, SU8220, Hitachi, Tokyo, Japan), and the elemental distribution was analyzed by energy-dispersive spectroscopy (EDS). The carbon structure of the coatings before and after corrosion was analyzed using Raman spectroscopy (LabRAM HR Evolution, Jobin Yvon, Pari, France), employing a 532 nm laser, with a scanning range of 500–2500 cm−1, where multiple points were probed across the sample to ensure accuracy. This non-destructive analysis method provides qualitative and quantitative insights into the structure and degree of graphitization of carbon materials, based on the location and intensity of characteristic peaks.
Unlike Raman spectroscopy, Fourier Transform Infrared (FTIR) spectroscopy is a type of absorption spectroscopy, an analytical method based on the interaction of infrared radiation with matter to characterize the bonded structure of atoms by recording the infrared of the sample after absorption or reflection. In this paper, the carbon structure and functional groups of the a-C coatings before and after corrosion were analyzed using FTIR spectroscopy (Nicolet IS50, Thermo Fisher Scientific, Waltham, MA, USA) in the mid-infrared range of wavelengths from 4000 to 400 cm−1.
Due to technical challenges in directly testing protective coatings on bipolar plates in actual fuel cells, an ex-situ measurement method was utilized for simulating the fuel cell operating conditions. Various research institutions use different protocols, typically involving an acidic mixture of sulfuric acid and a small amount of hydrofluoric acid [25]. In this study, an electrolyte solution of H2SO4 + 0.1 ppm HF at pH 3 was formulated according to the operating environment of fuel cell bipolar plates and the test temperature was set at 80 °C [26]. Electrochemical testing included open circuit potential (OCP), Tafel polarization, and potentiodynamic corrosion analyses under simulated fuel cell conditions using a CH1660E electrochemical workstation with a three-electrode system comprising a working electrode, an R0305-saturated Ag/AgCl reference electrode, and a platinum mesh auxiliary electrode.
ICR tests were conducted using standard procedures [27], where two pieces of gas diffusion layer (GDL; Toray TGP-H-60 carbon paper) were placed on gold-plated copper blocks, with the test specimen positioned between the GDLs. A universal testing machine was used to control the experimental pressure, and the resistance values were measured using a resistance meter (Merck RK2514). The resistance value was recorded every 0.2 MPa increase in pressure and the test pressure ranged from 0.2 to 1.8 MPa.

3. Results and Discussion

3.1. Coating Structure

The SEM images revealing the cross-sectional morphology of the T-1 and T-2 samples before corrosion are presented in Figure 1. It can be seen in Figure 1a that the surface of the T-1 sample prepared using the CVA process contains numerous protruding carbon particles. By contrast, in Figure 1c, the T-2 sample prepared using the FCVA process shows a smooth surface without apparent carbon particle protrusions, indicating a filtering effect of the bent tube on carbon particles. Figure 1b reveals large carbon particles embedded within the T-1 coating. In regions without these particles, the coating exhibits a dense, glass-like structure. On the other hand, Figure 1d displays a smooth and dense glass-like structure of the T-2 sample, without significant defect traces. We suspect that these large particles may have a negative impact on the corrosion resistance of the coating; this will be confirmed later.
Raman spectroscopy serves as a complementary technique to X-ray diffraction (XRD) and EDS techniques, expanding the characterization methods of materials. It is a low-cost, non-destructive method that relies on the rotational and vibrational modes to identify material structures and components and is particularly effective for carbon materials [28,29]. Notably, a-C exhibits Raman peak shifts within the range of 800–2000 cm−1. The vibrational breathing modes of sp2 carbon atoms in aromatic rings generate a D peak with A1 g symmetry at 1350 cm−1, indicating defective graphite structures. The E2g symmetric G peak at 1580 cm−1 arises from sp2-hybridized carbon atoms in aromatic rings and carbon chains, representing ordered graphite structures [30].
Figure 1e shows the Raman spectra of the regions with and without carbon particles in the T-1 sample before corrosion testing, alongside the T-2 sample spectrum. The curve at the bottom of the T-1 group corresponds to the non-particle region. The Raman spectra at the two positions in the T-1 sample clearly reveal different intensities of the G peak and D peak, with an ID/IG ratio of 1.65 and 1.36 for regions with and without carbon particles, respectively. This suggests that the proportion of sp3-hybridized carbon structure is higher in non-particle regions compared to carbon particle regions, indicating better mechanical and corrosion resistance properties. The Raman spectrum of the T-2 sample reveals an ID/IG of 1.10. Thus, the ID/IG value for a-C coatings prepared by the FCVA process is generally lower than those prepared by the CVA process.

3.2. Characterization of Mechanical Properties

In PEMFCs, the protective coatings on metal bipolar plates typically do not have stringent mechanical property requirements. However, changes in the mechanical properties such as the hardness of a-C coatings can be used as an indicator of the proportion of sp3/sp2-hybridized carbon structures within the coating. It has been widely established that the hardness of a-C coatings increases with the increase in the proportion of sp3-hybridized carbon [31,32,33]. The measured hardness, elastic modulus, and residual stress values of samples T-1 and T-2 are shown in Figure 2. The hardness and elastic modulus values of the two sets of samples are close, indicating a similar proportion of sp3/sp2-hybridized carbon structures in the coatings, which meets the study’s requirements for the carbon structure of the coatings. The residual stress of sample T-1 is lower than that of T-2, which can be attributed to the graphite-like carbon particles embedded in the coating. These particles reduce the coating density, thereby releasing the residual stress in the coating.

3.3. Electrochemical Corrosion Behavior Test

Figure 3 shows the surface morphology and EDS elemental mapping images of the T-1 and T-2 samples after 7200 s of corrosion at a constant potential of 0.6 V. After this corrosion period, no significant corrosion pits or defects were observed on the surface of both the T-1 and T-2 sample groups. Elemental analysis also confirms the absence of elements other than carbon, indicating that both samples maintain good corrosion resistance under a constant corrosion potential of 0.6 V. As shown in Figure 4, when the corrosion potential is increased to 1.6 V, the T-1 sample exhibits large corrosion pits and defects, with significant loss of the Ti transition layer at the bottom of the pits, exposing the underlying iron matrix elements. By contrast, the T-2 sample shows no significant corrosion pits, retaining a relatively intact surface morphology. Notably, after corrosion at the high potential of 1.6 V, traces of oxygen contamination are observed on the surfaces of both the samples, which is independent of whether neutral graphite particles are embedded in the coating.
Unlike Raman spectroscopy, FTIR spectroscopy is an absorption spectroscopy technique used to characterize atomic bonding structures based on the interactions between infrared radiation and materials [34]. By converting the time-domain infrared spectrum (revealing the radiation absorbed or reflected by the sample) into a frequency-domain signal, FTIR spectroscopy proves invaluable in characterizing carbon materials. In this study, mid-infrared spectra ranging from 4000 to 400 cm−1 were used to analyze the carbon structures and functional groups of a-C coatings before and after corrosion.
The corrosive environment of PEMFCs, which includes dissolved oxygen as well as corrosive ions like H+ and F−, reacts with carbon materials to form functional groups comprising C-H, C-O, and C-OH bonds. These characteristic bond structures are challenging to characterize using Raman spectroscopy. Table 2 [35,36,37,38] presents the characteristic peak positions in the FTIR spectra of a-C coatings, which are in the range of 600 to 3500 cm−1. The sp2-hybridized C=C peak appears around 1580 cm−1, while a strong peak at 1250 cm-1 indicates both sp2- and sp3-hybridized C-C bonds. The peaks in the range of 2800–3000 cm−1 are attributed to C-H structures, while the peak around 3400 cm−1 corresponds to the -O-H bond vibration. In the case of oxygen contamination, the characteristic peak at 1720 cm−1 indicates the presence of C-O and C=O structures.
The FTIR spectra of samples T-1 and T-2 before and after corrosion are shown in Figure 5. Both sets of samples exhibit characteristic absorption peaks in the wavenumber range of 2800–3400 cm−1 after corrosion, indicating significant hydrogen and oxygen contamination. Furthermore, as the corrosion potential increases from 0.6 V to 1.6 V, the intensity of these absorption peaks slightly increases. After corrosion at 1.6 V, the characteristic absorption peak representing graphite becomes more pronounced, indicating an increase in graphite-like structures within the coating. The FTIR spectroscopic results are consistent with the Raman characterization, illustrating that as the corrosion potential increases, hydrogen and oxygen contamination intensifies, leading to the formation of characteristic functional groups such as C-OH and C-H, along with a transition from sp3-hybridized carbon structures to sp2-hybridized ones.
Prior to the actual electrochemical tests, the samples were immersed in an corrosion test solution for 12 h to stabilize their surfaces. Subsequently, an OCP test lasting 1800 s was conducted. Dynamic potential and constant potential tests were performed once the OCP became stable, with fluctuations not exceeding 5%. To minimize testing errors, 2–3 repeated tests were conducted on each a-C coating sample.
Figure 6a shows the OCPs of bare 304 L stainless steel substrate and T-1 and T-2 samples. The addition of a-C protective coatings results in a significant increase in the OCP of the samples, stabilizing around 0.2 V, indicating improved corrosion resistance; this is due to the fact that corrosion of the coating occurs in the area of carbon particles on the surface, and there are fewer carbon particles in the T-2, which was confirmed earlier. By analyzing the Tafel region using the Tafel extrapolation method, the self-corrosion current density and self-corrosion potential of the coatings can be determined. A higher self-corrosion potential and lower self-corrosion current density indicate a lower rate of electrochemical corrosion, suggesting higher corrosion resistance of the coating [28,39,40].
Figure 6b displays the dynamic corrosion potentials of bare 304 L stainless steel substrate and T-1 and T-2 samples. The Tafel extrapolation fitting results for the dynamic potential curves are presented in Table 3. Both sets of samples with a-C protective coatings exhibit a noticeable decrease in the self-corrosion current density compared to the substrate, indicating enhanced corrosion resistance. The corrosion current densities of both the samples meet the 2025 technical specifications for bipolar plates set by the U.S. DOE. Notably, the corrosion current density of the T-2 sample is significantly lower than that of the T-1 sample. This is primarily ascribed to the presence of penetrating graphite-like carbon particles, and the gaps between the coating and these particles within the T-1 sample prepared using the CVA process act as channels for the corrosive solution to erode the substrate. Additionally, the difference in corrosion potentials between the graphite-like carbon particles and the coating exacerbates corrosion by initiating galvanic effects.
The potentiostatic polarization curve is used to analyze the corrosion current density of a-C coatings under constant potential. Considering that the corrosion rate at the cathode side is typically higher than that at the anode side in simulated fuel cell conditions, the potentiostatic analyses were all conducted under conditions simulating the cathode [29,41,42,43]. The surface potential of the metal bipolar plates significantly affects the corrosion rate of the coatings. Under normal operating conditions, the corrosion potential is usually around 0.6 V, while during high-power operation, the corrosion potential typically remains between 0.85 and 0.9 V, and can even approach 1.1 V [44]. Additionally, during fuel cell start–stop or significant power changes resulting in the insufficient supply of hydrogen or oxygen, a hydrogen void interface is formed at the cathode side, causing reverse electron migration and momentarily raising the surface potential of the metal bipolar plates to approximately 1.6 V [45].
Potentiostatic tests were conducted at a constant potential of 0.6 V for 7200 s under normal conditions and at 1.6 V for start–stop conditions. At 0.6 V, the corrosion current density for both sets of samples initially decreases and then stabilizes over time without significant fluctuations, indicating good corrosion stability of the coatings over extended corrosion processes. However, at 0.6 V, the T-1 sample exhibits a higher corrosion current density than the T-2 sample, indicating a higher corrosion rate of the T-1 sample. Under a high potential of 1.6 V during simulated start–stop scenarios, the corrosion current density of the samples increases significantly, accelerating the corrosion rate. The potentiostatic polarization curve of the T-1 sample exhibits large fluctuations, while that of the T-2 sample remains stable. The results under two different constant potential conditions indicate that the incorporation of graphite-like carbon particles in the T-1 coating reduces its corrosion resistance.
Figure 7 shows the Raman spectra and relevant fitting results of the T-1 and T-2 sample groups before corrosion and after 7200 s of corrosion at constant potentials of 0.6 V and 1.6 V. The parameters such as the integrated intensity ratio of the fitted D and G peaks ID/IG, G peak position, and full width at half maximum (FWHM) of G peak are provided in Table 4. The fitting results indicate that the T-1 and T-2 sample groups have similar carbon structures before corrosion. However, after corrosion at a constant potential of 0.6 V, the ID/IG shows a decreasing trend, suggesting the presence of oxygen-contaminated sp2-hybridized carbon structures within the coatings, consistent with previous studies. Notably, the ID/IG ratios of both groups of samples were increased compared to corrosion at 0.6 V after 7200 s of corrosion at 1.6 V, indicating a decrease in the proportion of sp3-hybridized carbon structures and an increase in sp2-hybridized carbon structures. Under high-potential corrosion conditions, sp3-hybridized carbon bonds in a-C coatings tend to break and recombine, leading to a transformation from sp3-hybridized to sp2-hybridized carbon forms.

3.4. ICR Test

The primary source of internal resistance in PEMFCs is the contact resistance between the metal bipolar plates and the GDL. Excessive internal resistance can significantly impair the cell’s output performance. Thus, ICR is a crucial measure of the performance of bipolar plate protective coatings. Numerous studies have established that the content of sp2-hybridized carbon structures governs the electrical conductivity of a-C coatings.
In this study, the contact resistance of the coatings was evaluated under pressures ranging from 0.2 MPa to 1.8 MPa using the method described by Wang et al. [5]. As shown in Figure 8, under a standard stack pressure of 1.4 MPa, the contact resistances of samples T-1 and T-2 before corrosion are 4.1 and 17.9 mΩ·cm2, respectively. Following corrosion, the contact resistance of both the samples increases. Specifically, after corrosion at a potential of 0.6 V, the contact resistances of samples T-1 and T-2 increase to 8.3 and 28.9 mΩ·cm2, respectively. After corrosion at 1.6 V, these values increase to 5.9 and 23.2 mΩ·cm2, respectively.
The increase in contact resistance can be primarily attributed to the oxygen contamination of the conductive sp2-hybridized carbon structures within the coatings. Interestingly, after undergoing corrosion at a constant potential of 1.6 V, the contact resistance of both sets of samples is lower than that after corrosion at 0.6 V but higher than that before corrosion. This suggests that although the coatings undergo a transformation from sp3-hybridized carbon structures to sp2-hybridized structures after corrosion at 1.6 V, the rate of this transformation is lower than the rate of contamination of the sp2-hybridized carbon structures by elements such as hydrogen and oxygen.
Notably, the contact resistance of sample T-1 is significantly lower than that of sample T-2, both before and after corrosion. The contact resistance of sample T-1 consistently meets the technical requirement set by the U.S. DOE, which specifies that the contact resistance should be less than 10 mΩ·cm2. This superior performance can be attributed to the embedded graphite-like carbon particles that reduce the longitudinal resistance of the coating. Additionally, the T-1 sample represented in Figure 9a has more and coarser graphite-like particles on the coating surface, while the T-2 sample represented in Figure 9b has a smoother surface; the protruding graphite-like carbon particles on the coating surface increase the contact area between the coating and the GDL, and the embedded graphite-like particles act as conductors to increase the longitudinal electrical conductivity of the coating, thereby reducing the contact resistance. By contrast, the sample T-2, which lacks embedded graphite-like conductive particles and has a smooth surface, exhibits excellent corrosion resistance but higher contact resistance.
Based on the above experimental results, a schematic diagram of the corrosion behavior of a-C coatings with and without large carbon particles is presented in Figure 10. As shown in Figure 10a, due to the difference in carbon structure between the carbon particles and coating, there are fissure defects between the coating and carbon particles, which will become the channel for erosion by corrosive solution, and the difference in carbon structure leads to a corrosion potential difference between carbon particles and coating, which will also form the primary cell corrosion effect between them to accelerate the corrosion and the structure of the carbon particles is not as dense as that of the coating. The lower corrosion potential of the carbon particles causes them to lose electrons first and corrosion occurs. As corrosion proceeds, corrosive ions continue to enter the gap and the carbon particles continue to corrode, and ultimately the gap between the carbon particles and the coating will expand; the corrosive solution can penetrate into the substrate through the gap, which ultimately leads to the failure of the protective coating. As these gaps widen, loosely attached particles may detach from the coating surface. By contrast, no obvious corrosion pits are observed in Figure 10b, and the coating surface morphology remains relatively intact. This is mainly ascribed to the relatively smooth surface of the coating without large carbon particles.

4. Conclusions

CVA and FCVA techniques were utilized to prepare a-C coatings with distinct surface morphologies while maintaining similar sp3/sp2-hybridized carbon structures. These coatings were subjected to high and low corrosion potential tests to simulate the operating conditions of PEMFCs. The coating morphologies, carbon structures, and contact resistance before and after corrosion were comprehensively examined. The main findings of the study are summarized as follows:
  • The embedded graphite-like particles effectively reduced the longitudinal resistance of the coatings. Under pressure, these particles increased the contact area between the coating and the GDL, significantly reducing the ICR. However, during corrosion, tiny gaps between the particles and the coating served as pathways for corrosive ions, reducing the coating’s corrosion resistance. Additionally, the corrosion potential difference between the particles and the coating led to galvanic effects, with the graphite-like particles being preferentially corroded due to their lower corrosion potential. The samples prepared using FCVA technology, which lacked embedded carbon particles, demonstrated excellent corrosion resistance under various conditions but exhibited higher contact resistance due to the absence of graphite-like particles.
  • The sp2-hybridized carbon structures within the a-C coatings were susceptible to contamination by elements such as hydrogen and oxygen in the solution after corrosion, leading to an increase in contact resistance. This phenomenon occurred regardless of whether graphite-like particles were embedded within the coatings. Under high-potential corrosion conditions, a noticeable transition from sp3-hybridized to sp2-hybridized carbon structures was observed.
  • The integration of techniques such as magnetron sputtering, CVA, and FCVA processes can facilitate the fabrication of composite a-C coatings, which comprise a metal Ti transition layer, a dense corrosion barrier layer without embedded graphite-like carbon particles (obtained via FCVA), and a conductive surface layer containing graphite-like carbon particles (obtained using CVA). This composite coating design aims to meet the stringent requirements for long-term protection of bipolar plates under the operational conditions of PEMFCs.
Overall, this study emphasizes the crucial role of embedded graphite-like particles in enhancing the conductivity and reducing the contact resistance of a-C coatings, albeit with potential drawbacks in terms of corrosion susceptibility.

Author Contributions

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

Funding

This research was funded by the Key Technologies R&D Program of Guangdong Province (No. 2023B0909060002).

Institutional Review Board Statement

Not applicable.

Informed Consent Statement

Not applicable.

Data Availability Statement

The data used to support the findings of this study are available from the corresponding author upon request.

Conflicts of Interest

Author Zhiqing Feng was employed by Guangdong Hydrogen Energy Institute of Wuhan University of Technology. Author Fengying Luo was employed by National energy key laboratory for new hydrogen-ammonia energy technologies. Author Ruiming Zhang was employed by the above two units. The remaining authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

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Coatings 14 01053 g001
Figure 1. Surface and cross-sectional morphology of T-1 and T-2 samples before corrosion: (a) surface of T-1 sample, (b) cross-section of T-1 sample, (c) surface of T-2 sample, (d) cross-section of T-2 sample, and (e) Raman spectra of T-1 and T-2 sample.
Figure 1. Surface and cross-sectional morphology of T-1 and T-2 samples before corrosion: (a) surface of T-1 sample, (b) cross-section of T-1 sample, (c) surface of T-2 sample, (d) cross-section of T-2 sample, and (e) Raman spectra of T-1 and T-2 sample.
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Figure 2. The mechanical properties of samples T-1 and T-2: (a) hardness, (b) elastic modulus, and (c) residual stress.
Figure 2. The mechanical properties of samples T-1 and T-2: (a) hardness, (b) elastic modulus, and (c) residual stress.
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Figure 3. Surface morphology and elemental distribution after 7200 s of corrosion under potentiostatic conditions at 0.6 V: (a) T-1 sample and (b) T-2 sample.
Figure 3. Surface morphology and elemental distribution after 7200 s of corrosion under potentiostatic conditions at 0.6 V: (a) T-1 sample and (b) T-2 sample.
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Figure 4. Surface morphology and elemental distribution after 7200 s of corrosion under potentiostatic conditions at 1.6 V: (a) T-1 sample, (b) T-2 sample.
Figure 4. Surface morphology and elemental distribution after 7200 s of corrosion under potentiostatic conditions at 1.6 V: (a) T-1 sample, (b) T-2 sample.
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Figure 5. Infrared spectra of T-1 and T-2 samples before and after corrosion: (a) T-1 sample and (b) T-2 sample.
Figure 5. Infrared spectra of T-1 and T-2 samples before and after corrosion: (a) T-1 sample and (b) T-2 sample.
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Figure 6. Corrosion polarization curves of T-1 and T-2 samples. (a) Open circuit potential, (b) potentiodynamic polarization curves, (c) potentiostatic polarization curves at 0.6 V, and (d) potentiostatic polarization curves at 1.6 V.
Figure 6. Corrosion polarization curves of T-1 and T-2 samples. (a) Open circuit potential, (b) potentiodynamic polarization curves, (c) potentiostatic polarization curves at 0.6 V, and (d) potentiostatic polarization curves at 1.6 V.
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Figure 7. Raman spectra before and after corrosion: (a) T-1 sample and (b) T-2 sample.
Figure 7. Raman spectra before and after corrosion: (a) T-1 sample and (b) T-2 sample.
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Figure 8. Interface contact resistance of T-1 and T-2 samples before and after corrosion; (a) T-1 sample and (b) T-2 sample.
Figure 8. Interface contact resistance of T-1 and T-2 samples before and after corrosion; (a) T-1 sample and (b) T-2 sample.
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Figure 9. Schematic diagram of the interfacial contact resistance of T-1 and T-2 samples; (a) T-1 sample and (b) T-2 sample.
Figure 9. Schematic diagram of the interfacial contact resistance of T-1 and T-2 samples; (a) T-1 sample and (b) T-2 sample.
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Figure 10. Schematic corrosion of a-C coatings with large particles (a) and without large particles (b).
Figure 10. Schematic corrosion of a-C coatings with large particles (a) and without large particles (b).
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Table 1. The experimental parameters.
Table 1. The experimental parameters.
Sample TypeProcessP/PaVbias/Vt/minIon Source Current/AArc Suoce Current/APower/kWThickness/nm
T-1Cleaning1.5−90020////
Ion etching0.8−600205///
Ti transition layer0.5−10024//2200
a-C layer0.2−20018/100/800
T-2Cleaning1.5−90020////
Ion etching0.8−600205///
Ti transition layer0.5−10024//2200
a-C layer0.005−2000120/100/800
Note: Vbias—substrate bias voltage; Power—magnetron sputtering power.
Table 2. Characteristic peak locations of infrared spectra in a-C materials.
Table 2. Characteristic peak locations of infrared spectra in a-C materials.
Wavenimber (cm−1)Probable AssignmentCharacteristic Vibration Mode
3400O-HDLC, oxy-hydrogenated film
3035~3085C-HDLC, a-C–H, (aromatic) sp2
2860~2920C-HDLC, a-C–H, DLC sp3
2180C≡CDLC sp1
1720C=ODLC contaminated with oxygen
1600~1640, 1640C=CDLC, C=C (Olefinic), sp2
1550~1600, 1580C=CDLC, C=C (Aromatic), sp2
1245C-Csp3/sp2
1100C-O-CDLC contaminated with oxygen
720graphiteDLC, attributable to the graphite form
Table 3. The fitted values of the dynamic potential polarization curves.
Table 3. The fitted values of the dynamic potential polarization curves.
Sample TypeSelf-Corrosion Potential Ecorr/VSelf-Corrosion Current Density Ecorr/μA·cm−2
T-1−0.0540.48
T-2−0.0850.2
Table 4. Raman analysis results before and after 0.6 V and 1.6 V corrosion tests.
Table 4. Raman analysis results before and after 0.6 V and 1.6 V corrosion tests.
Sample TypeID/IGFWHM (G)G-Peak
T-11.36167.91566.3
T-21.10195.31565
T-1—0.61.291781565.3
T-2—0.60.92204.91558
T-1—1.61.13192.1200
T-2—1.61.031564.71563
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MDPI and ACS Style

Feng, Z.; Zhou, Z.; Zeng, J.; Chen, D.; Luo, F.; Wang, Q.; Dai, W.; Zhang, R. Comparison of Corrosion Behavior of a-C Coatings Deposited by Cathode Vacuum Arc and Filter Cathode Vacuum Arc Techniques. Coatings 2024, 14, 1053. https://doi.org/10.3390/coatings14081053

AMA Style

Feng Z, Zhou Z, Zeng J, Chen D, Luo F, Wang Q, Dai W, Zhang R. Comparison of Corrosion Behavior of a-C Coatings Deposited by Cathode Vacuum Arc and Filter Cathode Vacuum Arc Techniques. Coatings. 2024; 14(8):1053. https://doi.org/10.3390/coatings14081053

Chicago/Turabian Style

Feng, Zhiqing, Zhetong Zhou, Junhao Zeng, Ding Chen, Fengying Luo, Qimin Wang, Wei Dai, and Ruiming Zhang. 2024. "Comparison of Corrosion Behavior of a-C Coatings Deposited by Cathode Vacuum Arc and Filter Cathode Vacuum Arc Techniques" Coatings 14, no. 8: 1053. https://doi.org/10.3390/coatings14081053

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