# Protective TiO x N y -based multilayer coatings for 316L stainless steel bipolar plates in a PEM water electrolyzer

Khaoula Chergui, Benoit Lescop, Michel Prestat, Geoffrey Ringot, Flavien Vucko, Valérie Demange, Ludivine Rault, Francis Gouttefangeas, Loïc Joanny, Sophie Ollivier, et al.

► To cite this version:

Khaoula Chergui, Benoit Lescop, Michel Prestat, Geoffrey Ringot, Flavien Vucko, et al. Protective TiO x N y -based multilayer coatings for 316L stainless steel bipolar plates in a PEM water electrolyzer. International Journal of Hydrogen Energy, 2025, 178, pp.151650. 10.1016/j.ijhydene.2025.151650. hal-05279958

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# Protective TiOₓNᵧ-based multilayer coatings for 316L stainless steel bipolar plates in a PEM water electrolyzer

Khaoula Chergui<sup>a</sup>, Benoit Lescop<sup>a</sup>, Michel Prestat<sup>b</sup>, Geoffrey Ringot<sup>b</sup>, Flavien Vucko<sup>b,c</sup>, Valérie Demange<sup>c,d</sup>, Ludivine Rault<sup>c</sup>, Francis Gouttefangeas<sup>c</sup>, Loïc Joanny<sup>c</sup>, Sophie Ollivier<sup>c</sup>, Michael Walls<sup>d</sup>, Gaël Maranzana<sup>e</sup>, Divino Salvador Ramirez-Rico<sup>f</sup>, Loïc Assaud<sup>g</sup>, Stéphane Rioual<sup>a,w ID</sup>

## ARTICLE INFO

Handling Editor: M Mahdi Najafpour

### Keywords:

PEM water electrolysis

Bipolar plate

Interfacial contact resistance

Coating

Corrosion

Stainless steel

Titanium oxynitride

## ABSTRACT

The present work reports the optimization of magnetron-sputtered titanium/titanium oxynitride thin-film coatings for application on 316L stainless steel bipolar plates (BPP) in proton exchange membrane water electrolyzers (PEMWE). As it will be shown, these coatings display tunable electrical conductivity and mechanical hardness to optimize the interfacial contact resistance (ICR). By controlling the elemental composition and thickness of the coating, and by using multilayers instead of single TiON layers, the ICR can be reduced to 1.5 mΩ cm², a value lower than the targeted ICR (10 mΩ cm²) for BPP in PEMWE. Moreover, due to its stable oxide nature, in contrast to uncoated steel, the coated steel was undamaged after 168 h at 60 °C under the experimental conditions found in PEMWE, in which the BPP surface potential is decoupled from the cell voltage. This proposed low-cost coating should therefore be considered as a promising material to enhance the competitiveness of the PEMWE technology.

## 1. Introduction

The production of green hydrogen is projected to be a central component in advancing decarbonized energy infrastructures and industries [1–4]. The proton exchange membrane water electrolyzer (PEMWE) is a near-ambient-temperature hydrogen production technology that is based on the electrochemical splitting of water molecules into oxygen and hydrogen [5]. A PEMWE stack comprises bipolar plates (BPP) and porous transport layers (PTL) on both sides of a catalyst-coated membrane (CCM) where oxygen and hydrogen evolution reactions take place. BPPs play a fundamental role in the performance of PEMWE since they ensure electrical conduction in the stack and maintain sealing between components. They also enable the evacuation of gases produced at the electrodes and ensure a uniform mechanical pressure [6]. PEMWEs are highly sensitive to the presence of

cationic impurities that can be produced by the dissolution of the stack elements (BPP and PTL). Indeed, metal cations such as Fe³⁺, Ni²⁺ and Cr³⁺ may accumulate in the proton exchange membrane (e.g Nafion™) and in the ionomer within the catalyst layers, due to their higher affinity for sulfonic acid groups than H⁺ [7]. This type of contamination reduces the membrane's proton conductivity, leading to the degradation of the cell performance over time. Metal-ion pollution may also reduce the effectiveness of the catalyst. Li et al. [8] reported that a contamination of only 1 ppm of Fe³⁺ under test conditions of 0.5 A cm⁻² and 60 °C is sufficient to cause a significant drop in electrolysis durability, and this was confirmed by other studies [9,10]. In the anodic compartment of the PEMWE, where dissolution processes and metal leaching are more likely to occur, titanium is currently the state-of-the-art BPP (and PTL) material. Titanium is usually coated with platinum-group metals (PGM), such as platinum or iridium, due to their good corrosion resistance and low

<footnote><sup>a</sup> Univ Brest, CNRS, Lab-STICC-UMR 6285, Brest, France</footnote>

<footnote><sup>b</sup> French Corrosion Institute, Brest, France</footnote>

<footnote><sup>c</sup> Univ Rennes, CNRS, ISCR – UMR 6226, ScanMAT-UAR 2025, Rennes, France</footnote>

<footnote><sup>d</sup> Univ of Paris-Saclay, Laboratory of Solid-State Physics, Saclay, France</footnote>

<footnote><sup>e</sup> Univ de Lorraine, CNRS, LEMTA, Vandoeuvre-les-Nancy, France</footnote>

<footnote><sup>f</sup> ICMMO-ERIEE, Université Paris-Saclay, UMR CNRS 8182, 91405, Orsay, France</footnote>

<footnote><sup>g</sup> University Paris Sciences & Letters, ESPCI, ENS, CNRS UMR 8004, Institute of Porous Materials of Paris, 10 Rue Vauquelin, 75005, Paris, France</footnote>

<footnote>* Corresponding author.
E-mail address: rioual@univ-brest.fr (S. Rioual).</footnote>
interfacial contact resistance (ICR) with the adjoining PTL. However, this directly affects the capital expenditures (CAPEX) of the PEMWE technology and undermines its competitiveness. It is reported that BPPs account for ca. 50 % of stack costs [11] and are thus central to the reduction of overall costs. The choice of such costly BPP materials was first dictated by the view that BPP, PTL, and anode experience similar electrochemical conditions, i.e., the BPP would exhibit the same highly anodic potential as that of the anode, higher than 1.5V vs reversible hydrogen electrode (RHE). Ex situ corrosion investigations of BPP candidate materials were therefore mostly performed at around 2V vs RHE [12–15]. This view has been contested in recent years by the findings of Becker et al. [16] indicating that the potential within a sintered PTL remains unaffected by the current density in a PEMWE. In a subsequent work, Becker et al. described the spatial distribution of potentials in the PEMWE cells [17]. The authors reported that the potential on the anodic side decreases with distance from the CCM, gradually reaching the open circuit potential (OCP) at the vicinity of the BPP (or slightly higher if galvanically coupled with the PTL). To further prove this statement, these authors produced a BPP coated with sputtered carbon that did not exhibit any degradation after 30 days of electrolysis at 2 A cm⁻² at 60 °C. This finding clearly supports the view that the BPPs are not exposed to anodic potentials high enough to degrade carbon. This phenomenon is explained by the high resistivity of high-purity feed water, which leads to the so-called “ionic decoupling” of the potential at the BPP from the cell voltage. In other words, ionic decoupling means that in ultrapure water, the BPP acts as a simple current collector at around open-circuit potential. Hence, it seems more appropriate to use a potential around the OCP value and near-neutral pH for conducting ex situ corrosion testing of BPP materials [18]. Note that this “nominal” behavior of the BPP might be questioned in case of contamination of the water with leaching of metallic ions coming from the water feed or fluoride-anion release from an ageing membrane, which would increase the electrical conductivity of the solution. A thin PTL might also place the BPP too close to the catalyst layer, and outside the conditions for ionic decoupling [17].

The findings on ionic decoupling further suggest that stainless steel (SS) could be a viable candidate for BPPs, as it offers greater cost effectiveness, higher availability and machinability compared to titanium for industrial-scale production. However, due to the high sensitivity of PEMWEs to cationic pollution, it is desirable to coat stainless steel BPPs to prevent metal leaching in pure water at temperatures higher than 60 °C. Another corrosion phenomenon, called rouging, that is well known in the pharmaceutical industry, can indeed occur on unprotected stainless steel surfaces in contact with heated pure water (pipes, valves, tanks ...) forming a thin, reddish iron oxide layer (though other discolorations can also be observed) [19–22]. For instance, Bernardi et al. [20] reported the rouging corrosion of SS 316L exposed at 70 °C for three weeks in the so-called “water for injection” (WFI, a type of pure water with conductivity of ca. 1 µS cm⁻¹). During rouging of steel, ionic species from the SS surface are likely to be released into the water, which is not desirable for PEMWE. The mechanism and kinetics of rouging are still poorly understood. However, a possible explanation lies in the instability of the passive layer, which consists mainly of chromium oxide (Cr₂O₃), in pure water [23,24]. This leads to the dissolution of metal ions such as Fe²⁺ or Fe³⁺ into the solution and the formation of insoluble iron oxide on the surface. In addition to protection from surface oxidation and metal leaching, the second criterion for the coated BPP concerns the interfacial contact resistance (ICR), which must be lower than 10 mΩ cm² [25] to minimize ohmic losses in the water electrolysis stack. According to several authors, developing coatings with high electrical conductivity at the surface is a key point for obtaining low ICR values [26,27]. The challenge is then to produce a highly electrically conductive coating with a low value of the ICR, and with good stability, preferably without the formation of any passivation layer which may modify the electrochemical impedance in solution.

Among the different methods for producing coatings, PVD (Physical

Vapor Deposition) is very attractive since it is scalable and can be applied industrially. Some research groups have investigated PVD-coated SS for BPP applications [15,28,29]. However, only a few studies have investigated the stability of SS and PVD-coated SS in near-neutral solutions [29,30]. Rojas et al. [29] consider CrN/TiN, Ti/TiN, Ti and TiN coatings deposited on SS 316L, SS 904L and SS 321. The results indicate that the most effective substrate/coating combination is a Ti/TiN multi-layered coating on SS 321, which exhibits minimal dissolution under 2V vs SHE after 2 h. Laedre et al. [30] investigated the protective effect of a bi-layer coating composed of a tantalum base layer with an indium tin oxide (ITO) top layer applied to a titanium BPP surface. The results reveal that this coating offers good corrosion resistance with an ICR of 30 mΩ cm² after polarization at +2 V vs. RHE in sulfuric acid at pH 5.5 for more than one day.

In this context, the present work reports the protection of 316L substrates by PGM-free coatings prepared by PVD. Among the various candidates, conductive titanium oxynitride (TiOₓNᵧ) emerges as a particularly promising coating. Its oxide-based nature is expected to enhance the surface stability and suppress the formation of passive interfacial layers. Additionally, its elemental composition can be tailored to optimize the ICR. The initial phase of the present work is dedicated to characterizing a range of coatings and optimizing their ICR performance. Particular attention is given to the influence of electrical conductivity and hardness, with a comparative analysis of multilayer versus single-layer architectures. Finally, the protective behavior of the coatings under anodic conditions – mimicking BPP operation with ionic decoupling – is evaluated over extended exposure times of up to 168 h.

## 2. Materials and methods

### 2.1. Thin-film deposition process

Single-layer TiOₓNᵧ and multilayer TiOₓNᵧ/Ti/TiOₓNᵧ thin films were deposited using the DC reactive magnetron sputtering method (PRO Line, Lesker), employing a high-purity 3" titanium target (99.999 %). Different substrates were used: (100) silicon wafers, glass, and as-received SS 316L (50 × 50 × 2 mm³). Silicon wafers and glass substrates were selected for X-ray diffraction (XRD) and electrical conductivity measurements, respectively. Before deposition, the substrates were cleaned via ultrasonication in acetone and ethanol and rinsed with deionized water. Before the sputtering process, the deposition chamber was evacuated to a base pressure of 1 × 10⁻⁷ mbar, and the Ti target was then pre-sputtered for 15 min to remove any surface contamination. The deposition was carried out with a mixture of high-purity (99.9996 %) Ar and N₂ reactive gases with controlled flow rates and a total power of 75 W. To adjust the chemical compositions of the TiOₓNᵧ coatings, three sets of experimental conditions were defined, as summarized in Table 1. In particular, the amount of nitrogen and the total working pressure were varied to obtain three TiOₓNᵧ variants with different nitrogen contents. Metallic titanium was also deposited using the conditions given in Table 1. In addition to the production of single TiOₓNᵧ thin films, a multilayer approach was included in the study, alternating TiOₓNᵧ and metallic Ti layers.

### 2.2. Composition and structural characterization

Coating thickness was measured using a KLA+ (KLA tencor) mechanical profilometer. The crystallographic structure and orientation of the different coatings were analyzed by XRD, using a PANalytical Empyrean diffractometer, with a CuKα1,2 radiation source (λ = 0.15418 nm). Analyses were performed in Bragg-Brentano configuration and the XRD patterns of the coatings were compared to that of TiN powder (99.5 %, CAS 25583-20-4). The composition of the different coatings deposited on SS 316L was measured on five distinct regions using a JEOL JSM 7100 F scanning electron microscope (SEM) equipped with an energy dispersive X-ray (EDX) detector (SDD X-Max 50 mm²
<table><caption>**Table 1**
Coating deposition conditions, elemental composition and electrical conductivity of the TiOₓNᵧ films.</caption><thead><tr><th rowspan="2">Sample</th><th rowspan="2">N<sub>2</sub> flow (sccm)</th><th rowspan="2">Total working pressure (mbar)</th><th colspan="3">Atomic composition (Substrate: SS 316L)</th><th rowspan="2">Electrical conductivity (S cm<sup>-1</sup>) (Substrate: glass)</th></tr><tr><th>%Ti</th><th>%O</th><th>%N</th></tr></thead><tbody><tr><td>TiO<sub>x</sub>N<sub>y</sub>-a</td><td>1</td><td>5.99 × 10<sup>-3</sup></td><td>29.8 ± 0.6</td><td>51.8 ± 0.1</td><td>18.4 ± 0.5</td><td>50</td></tr><tr><td>TiO<sub>x</sub>N<sub>y</sub>-b</td><td>1</td><td>4.66 × 10<sup>-3</sup></td><td>34.6 ± 0.3</td><td>30.6 ± 0.3</td><td>34.8 ± 0.2</td><td>173</td></tr><tr><td>TiO<sub>x</sub>N<sub>y</sub>-c</td><td>6</td><td>4.66 × 10<sup>-3</sup></td><td>39.6 ± 0.6</td><td>15.8 ± 0.3</td><td>44.6 ± 0.5</td><td>10<sup>3</sup></td></tr><tr><td>Ti</td><td>0</td><td>5.73 × 10<sup>-3</sup></td><td>—</td><td>—</td><td>—</td><td>—</td></tr></tbody></table>

Oxford Instruments). A 20–40 lamella was prepared by Focused Ion Beam (FIB) and analyzed using high angle annular dark field (HAADF) imaging and electron energy-loss spectroscopy (EELS) in a Nion UlstraSTEM 200 scanning transmission electron microscope (STEM). The electrical conductivity of the different coatings deposited on glass substrates were assessed via the four-probe method and using the HIOKI RM3545-02 mΩ-meter.

Atomic force microscopy (AFM) measurements were performed to probe the mechanical properties of the coatings using a Bruker Dimension Icon AFM equipped with an Sb-doped silicon tip (TESPA-V2). The Derjaguin-Muller-Toporov (DMT) model was used to measure the reduced Young's modulus in PeakForce Quantitative Nanomechanics mode, with the cantilever vibrating at 2 kHz. This frequency allows accurate mapping of mechanical properties by recording topography and analyzing force curves at each pixel. According to Lin et al. [31], a direct one to one correlation between the elastic modulus and the hardness was obtained for Ti/TiN layers. Thus, for simplification purpose, the hardness values of the coatings in this study were assimilated to the DMT moduli measured by AFM.

## 2.3. Interfacial contact resistance (ICR) measurement

ICR was measured between a coated and uncoated 316L substrate and a gold-plated 10 mm diameter copper cylinder. A manual press was used to apply a pressure of between 127 and 760 N cm⁻². The applied force was monitored using a force sensor. The resistance across this configuration was measured using a Hioki 4-wire milliohm meter, connecting one clamp to the copper cylinder (whose inherent resistance is negligible) and the other directly to the sample.

## 2.4. Electrochemical experiments

In order to evaluate the stability of uncoated and coated SS 316L, electrochemical experiments were conducted using both Origalys (Origaflex OGF500) and Gamry interface 1010E potentiostats. A double-jacketed three-electrode cell was used, containing approximately 90 mL of very dilute sulfuric acid solution at pH 5.5: Two types of solutions were used: the first consisted of highly dilute sulfuric acid with low conductivity (1.8 µS cm⁻¹); the second was adjusted to the same pH using dilute sulfuric acid and NaOH, resulting in higher conductivity (2.7 mS cm⁻¹). The temperature was kept stable at 60 °C using a LAUDA ECO RE 415 S thermostatic circulation bath. During the experiment, a platinum-coated titanium electrode and a reversible hydrogen electrode (Hydroflex, Gaskatel) were used as counter electrode (CE) and reference electrode (RE), respectively. Bare SS 316L and SS 316L coated with a TiOₓNᵧ single layer or a TiOₓNᵧ/Ti/TiOₓNᵧ multilayer, were used as working electrode (WE) with an exposed geometric surface area of 10 cm². The samples were immersed in the solution for 1 h to reach equilibrium potential. This was followed by a sequence of alternating electrochemical impedance spectroscopy (EIS) and open circuit potential (OCP) measurements. Each OCP measurement in the sequence lasted 1 h. The EIS was carried out with an excitation voltage of 10 mV, over a frequency range from 100 to 0.01 Hz.

# 3. Results and discussion

## 3.1. Characterization of single TiOₓNᵧ layers

TiOₓNᵧ coatings with different compositions were produced to optimize both the ICR value and the stability. As detailed in Table 1, the TiOₓNᵧ-a and TiOₓNᵧ-b films were grown with the same nitrogen flow of 1 sccm, but at two different pressures: 6.0 10⁻³ mbar and 4.7 10⁻³ mbar, respectively. On the other hand, the TiOₓNᵧ-c film, was produced with a high nitrogen flux of 6 sccm and a pressure of 4.7 10⁻³ mbar. The thickness of these films is 200 nm. The atomic compositions of films deposited on SS 316L substrates measured by EDX are displayed in Table 1. As shown, the atomic composition of the three TiOₓNᵧ films differs. The TiOₓNᵧ-c film shows the lowest oxygen and highest nitrogen content, while TiOₓNᵧ-a presents a high oxygen content. Table 1 also reports the electrical conductivity of these films. The electrical conductivity increases with decreasing oxygen content by two orders of magnitude, allowing the investigation of the influence of the electrical conductivity on the ICR measurements. Note that the observed variation is consistent with the effect of oxygen on electron mobility [32]. The presence of oxygen in the films might be questionable since this chemical element is not introduced in the PVD chamber during deposition. This point is well documented in the literature [32–35]. Alfath et al. [33] attribute the observed oxygen in TiN films to the presence of oxygen contaminants in the residual atmosphere of the PVD chamber. It is indeed known that titanium reacts more easily with oxygen than with nitrogen, as the formation of titanium oxide is thermodynamically more favorable than titanium nitride. The Gibbs energies of formation are more negative for titanium oxide (ΔG° ≈ −1582.3 kJ/mol) than for titanium nitride (ΔG° ≈ −86.2 kJ/mol), which favors the incorporation of oxygen into the film [34].

Fig. 1(a) shows a top-view SEM image of the 200 nm-TiOₓNᵧ-a film deposited on SS 316L, revealing a nanogranular structure with an estimated averaged grain size of 23 nm. Such nanogranular structure is also representative of other TiOₓNᵧ films (not shown). HAADF STEM analysis was then used to evaluate the internal structure of this film, as shown in Fig. 1(b and c). On Fig. 1(b), some defects are visible on the top surface of the coating, however, their depth is limited to about 10 nm, preventing the electrolyte from reaching the steel substrate during stability tests. Fig. 1(c) displays the internal structure of the film with a higher resolution. A high density is observed with the presence of 1–2 nm diameter pores.

Fig. 2(a–e) displays the cross-section mapping of chemical elements, and (f) shows a typical EELS spectrum measured in the coating with the N, Ti, and O peaks. The spectra at each pixel were used to obtain the elemental maps depicted in Fig. 2(a–e). However, as seen in Fig. 2(f), a quantification from the EELS spectrum is difficult due to signal overlapping, background subtraction, and cross-section uncertainties. Film compositions were therefore not determined by EELS but by EDX. Chromium and iron maps are presented in Fig. 2(d and e). These elements are observed only in the SS 316L, showing a very well-defined interface between the coating and SS 316L. Inside the coating, as shown in Fig. 2(a–c), the titanium and nitrogen contents are almost constant across the film thickness. However, the defects observed at the surface in Fig. 1(b) are still present on these two maps. In Fig. 2(b), the oxygen content is slightly decreasing from the substrate to the external
<figure><img src="image_3.png" /><figcaption>Fig. 1. (a) Top-view SEM micrograph and (b, c) HAADF STEM cross-section micrographs of 200 nm thick TiOₓNᵧ-a coated SS 316L.</figcaption></figure>

<figure><img src="image_4.png" /><figcaption>Fig. 2. (a–e) EELS maps of 200 nm TiOₓNᵧ-a coated SS 316L, (f) a typical spectrum from the deposited layer over the range including the N, Ti and O signals.</figcaption></figure>

surface. An excess of oxygen is also observed at the steel/coating interface, probably due to the native passive layer of 316L stainless steel.

Fig. 3 shows XRD patterns of thin films with 200 nm thickness on a silicon substrate. Reflections of the face-centered cubic phase are observed for the three films. The 200 peak dominates, indicating a preferential orientation along this direction. The diffraction patterns of the thin films are compared with those of TiN powder. It can be easily observed that the 200 reflection of the TiOₓNᵧ films is shifted relative to that of TiN. This shift increases towards a higher angle as the oxygen content increases. This behavior is explained by the substitution of nitrogen by oxygen in the material. A similar shift was observed by Qi et al. [35] on TiN thin films produced by DC magnetron sputtering. The authors also report the presence of oxygen in their film. Ali et al. [36] produced TiOₓNᵧ films by using both oxygen and nitrogen reactive gases during deposition and attribute the observed shift to anionic substitution [37–39], which leads to a variation of the cell parameter. As observed in Fig. 3, the 200 peak of the films becomes narrower with the decrease in oxygen content. The grain size was calculated from the width at half-maximum (FWHM) of the 200peak using the Debye-Scherrer equation. The results show a slight increase in grain size from 20 nm for TiOₓNᵧ-a to 32 nm and 26 nm for TiOₓNᵧ-b and TiOₓNᵧ-c, respectively.

<figure><img src="image_5.png" /><figcaption>Fig. 3. XRD patterns of the three deposited TiOₓNᵧ single-layers compared to TiN powder (as reference).</figcaption></figure>
## 3.2. ICR characterization of coated SS 316L by TiOₓNᵧ single layers

As stated above, the ICR value of coated steel is a crucial criterion for the development of new materials dedicated to PEMWE. The ICR of bare SS 316L and two TiOₓNᵧ-(a,c) coated SS 316L was measured as a function of the applied pressure. These two TiOₓNᵧ coatings were selected due to their high and low oxygen content. The ICR of bare SS 316L is depicted in Fig. 4(b) in black color. An almost constant low value is observed over the entire pressure range. It is below the targeted 10 mΩ cm² value represented by the horizontal dashed line in the figure. For the high-oxygen-content TiOₓNᵧ-a coating (sample A), in red color, a value of approximately 250 mΩ cm² is observed at pressures above 400 N cm⁻². The low-oxygen TiOₓNᵧ-c, which displayed a higher electrical conductivity, has a high ICR, exceeding 300 mΩ cm², the measurement range of the multimeter. The data are therefore not included in the graph. This shows that decreasing the ICR cannot be achieved by simply increasing the film conductivity. This suggests that other factors strongly influence the ICR value (such as hardness, see below). To give some further insights into the impact of the electrical conductivity of the coating (σcoating) on the measured ICR (ICRmeas), it is important to consider the theoretical relations between them. The electrical conductivity of the coating is given by the following equation:

$$ \sigma_{coating} = \frac{t_{coating}}{R \times S} \qquad (3) $$

Where tcoating is the thickness of the coating, R is the electrical resistance, and S is the surface area through which the current flows. The product of R and S is the area specific electrical resistance of the coating (ASR-coating), and the measured ICRmeas is the sum of the actual ICR and the ASRcoating:

$$ ICR_{meas} = ICR + ASR_{coating} \qquad (4) $$

The value of the ASRcoating calculated from the data of Table 1 is 0.4 10⁻³ mΩ cm² for the TiOₓNᵧ-a coating. It is well below the measured ICR (300 mΩ cm²) and targeted value (10 mΩ cm²), indicating that this term can be neglected provided that the coating material is carefully selected to have a sufficiently high conductivity. This confirms that the increasing the electrical conductivity of the coating has no influence on ICRmeas.

The measured ICR value of the TiOₓNᵧ-a coated SS 316L can be compared with data reported in the literature for TiN films deposited on SS 316L under a compaction force of 140 N cm⁻². Sun et al. [40] obtained an ICR of around 60 mΩ cm² for a 1 µm TiN film deposited by magnetron sputtering. Rojas et al. [29] report an ICR of around 5 mΩ cm² for a 1.8 µm TiN film. Using cathodic arc deposition for the production of a 2 µm thick TiN film, Lee et al. [41] measured an ICR of around 10 mΩ cm². All these values are well below our result, and suggest the influence of the film thickness on ICR. However, it is important to note that there is no absolute ICR value, as the measurements strongly depends on setup used for the characterization, in particular, the plates and compression pads. In our study, as in those cited above, gold-plated copper plates were used, in contrast to some authors [42–45]. We also cannot exclude some differences in the ICR values ascribed to the presence of a gas diffusion layer not used in the present study. Due to the technical differences in the setup used in the literature, to highlight the role of film thickness on the ICR, coated steel with a 400 nm thick layer of TiOₓNᵧ-a was produced and characterized. As shown in Fig. 4(a), this sample corresponds to the sample B. The ICR of this sample is shown in Fig. 4(b), in blue color. A clear decrease in the ICR is observed, confirming the influence of the thickness on the ICR. Moreover, the measured value approaches the targeted value represented by dashed line in the figure. This finding also confirms that the ASRcoating (that depends on the coating thickness, see Eq. (1) and Eq. (2)) does not govern the ICRmeas value.

<figure><img src="image_6.png" /><figcaption>Fig. 4. (a) Schematic representation of the SS 316L coated with single layers and multilayers. (b) ICR as a function of the clamping pressure of uncoated (black) and the coated SS 316L. The horizontal dashed line represents the target value of 10 mΩ cm². (c) ICR measured at 140 N cm⁻² for the coated SS 316L, the electrical conductivity of each coating is displayed on the top of the bars.</figcaption></figure>
### 3.3. ICR characterization of coated SS 316L by multilayers

An alternative approach for improving the ICR is to consider multilayer coatings. For instance, when using C/CrN and C/TiN multilayers instead of single layers, Li et al. [45] demonstrated a reduction of the ICR. The authors explained this behavior by a better homogenization of the current flowing through the multilayers. Better ICR values were also observed for CrN/CrAlN multilayer coatings compared to CrN and CrAlN coatings [27]. The authors attribute this behavior to the presence of a highly conductive layer (CrAlN) at the surface with respect to the inner ceramic layer (CrN). Similarly, in Ref. [26], the authors explained the low ICR value of Cr/CrN coatings by the high electrical conductivity of the multilayer coating. It is noticeable that, most of the time, the strategy of using multilayers in the literature is not applied to decrease the ICR but rather to prevent the formation of continuous, vertically aligned pores during columnar growth (often observed in PVD) and eliminate pinholes, thereby making the film structure more compact. This improves barrier properties, hindering the access of the electrolyte to the substrate, and enhances the corrosion resistance [45].

To investigate in detail the advantage of multilayers with respect to a single layer on the reduction of the ICR, we consider the presence of a layer of metallic titanium between two layers of TiOxNy-a. As shown in Fig. 4(a), several multilayer configurations were developed by varying the thickness of the layers. The A and B samples refer to the single layer TiOxNy-a coatings with 200 and 400 nm thickness, respectively. The C, D, E samples correspond to several multilayers with different thicknesses. Fig. 4(b) presents the measured ICR values as a function of the applied force for steel coated by multilayers. The C sample is a TiOxNy (100 nm)/Ti(50 nm)/TiOxNy (50 nm) multilayer with the same thickness as the A sample. The ICR of the C sample is lower than that of the A sample, indicating that a replacement of the inner TiOxNy part of the A coating by a metallic Ti layer improves the ICR. For the D sample, the thickness of the first deposited TiOxNy coating is increased with respect to the C sample. This change improves the ICR value, as already shown for single layers. In this case, the ICR is below the target value of 10 mΩ cm². A further improvement is obtained with the E sample with the structure TiOxNy (200 nm)/Ti (200 nm)/TiOxNy (200 nm). Fig. 4(c) shows the ICR values measured at 140 N cm⁻² for all these coated steel samples. As shown, the C sample exhibits an ICR value of 197 mΩ cm², which is lower than the 300 mΩ cm² measured for the single layer with the same thickness (200 nm). This underlines the interest of using multilayers instead of single layers. For the D and E multilayers, the measured ICR varies from 13.5 mΩ cm² to 1.53 mΩ cm². The thicker 600 nm TiOxNy/Ti/TiOxNy multilayer, with the lowest ICR, stands out consequently as the most promising coating of the series, and the most suitable for BPP applications in PEMWE. These ICR values are comparable with those reported in the literature for SS 316L with titanium based coatings: 11 mΩ cm² for a 2.75 µm Ti/TiN thickness coating [46], and 15 mΩ cm² for 1 µm TiN/CrN thickness coating [42], demonstrating the quality of the multilayers produced here.

To provide some further insights into the role of the electrical conductivity on the ICR, we also measured the electrical conductivity of the studied coatings, as shown in Fig. 4(c). When comparing C and D multilayers, the electrical conductivity increases by a factor of 1.5, while the ICR is reduced by a factor of 14.5. The same behavior is observed when changing from the D to the E multilayer. Similarly, samples A and C display close ICR values but their electrical conductivity values differ by two orders of magnitude. These findings indicate that the intrinsic electrical conductivity of the coating is not the main factor which governs the ICR and differ from the explanation of S. Pugal Mani et al. [27] on the low ICR values of the CrN/CrAlN multilayers.

AFM measurements were also performed to probe the mechanical properties of the coatings. DMT moduli are presented in Fig. 5 for the C, D and E multilayers. As shown, the hardness of the multilayer changes, with measured values of 8.34, 2.99, and 2.47 GPa respectively. This reduction in hardness may explain the decrease in the ICR. According to Persson et al. [47], electrical contact resistance is closely related to the mechanical properties of the interface. The hardness influences the nature of the asperity contacts, which largely determine the resistance, due to current restrictions across these contact regions. The presence of metallic titanium might modify the ICR by reducing the hardness of the multilayer structure. Indeed, the hardness of Ti is expected to be lower than that of a TiOxNy thin film. The fact that titanium in the multilayer affects its overall hardness is reinforced by the work of Lin et al. [31], who compared the hardness of a TiN monolayer and a Ti/TiN multilayer. They demonstrated that the Ti/TiN multilayer has a lower hardness (7.5 GPa) than the TiN monolayer (14 GPa). However, more research is needed to confirm our observations and to correlate the measured ICR with hardness.

### 3.4. Exposure tests on uncoated and TiOxNy (200 nm)/Ti (200 nm)/TiOxNy (200 nm) coated steel

Following the previous section, TiOxNy (200 nm)/Ti (200 nm)/TiOxNy (200 nm) coated steel was selected as the most promising material for PEMWE. Fig. 6(a) displays the STEM HAADF cross-sectional image of TiOxNy (200 nm)/Ti (200 nm)/TiOxNy (200 nm) multilayer on 316L. Despite the presence of some external pollutants (light-contrast round spots) well defined interfaces are observed with the presence of some pores in the inner part of the coating. However, these latter do not propagate from the surface to the steel substrate and are unlikely to impact the protective properties of the coating in the solution. Fig. 6 (b–d) displays the EELS mapping of chemical elements. Metallic titanium is present in the intermediate layer in between the two inner and outer oxynitride layers that are oxygen- and nitrogen-rich, as expected.

In order to assess the protective properties of this material, the electrochemical behaviour of bare and coated 316L samples were compared. First, bare SS 316L was exposed to a dilute solution of H₂SO₄ at pH 5.5, maintained at OCP and 60 °C. Under these conditions, the initial conductivity of the solution was 1.8 µS cm⁻¹. EIS and OCP

<figure><img src="image_7.png" /><figcaption>Fig. 5. DMT modulus maps obtained for TiOxNy/TiOxNy multilayer-coated SS 316L with different architectures (samples C, D, E of Fig. 4(a)).</figcaption></figure>
<figure><img src="image_8.png" /><figcaption>**Fig. 6.** (a) HAADF STEM cross-section micrograph and (b–e) EELS maps of TiOxNy (200 nm)/Ti (200 nm)/TiOxNy (200 nm) multilayer-coated SS 316L. The round-shaped stains are artefacts (pollution) coming from the TEM lamella preparation.</figcaption></figure>

measurements were carried out over a period of 168 h (7 days), as shown in Fig. 7. The evolution of the low-frequency phase angles, illustrated in the Bode diagram in Fig. 7(a), shows a gradual increase towards higher values, suggesting capacitive behavior over time due to surface passivation. Fig. 7(b) shows the evolution of the modulus of impedance $|Z|$ as a function of the frequency. The values at high frequency reflect the electrolyte resistance. A gradual decrease in this resistance over time is observed, which is attributed to the leaching of 316L constituents in the solution. This is confirmed by the increase in conductivity, from 1.8 µS cm⁻¹ to 30 µS cm⁻¹ at the beginning and end of the experiment, respectively. In addition, a progressive decrease in OCP was recorded during exposure (Fig. 7(c)) which also indicates the release of metallic cations into the solution. A similar phenomenon is reported in the literature for SS316L BPPs [48,49] and is attributed to the dissolution of metallic elements such as Fe, Cr, Mo, Ni and Mn [49].

EIS and OCP curves measured during the exposure of steel coated by the TiOxNy (200 nm)/Ti (200 nm)/TiOxNy (200 nm) multilayer are shown in Fig. 8. In contrast to results depicted in Fig. 7, no significant variation of the impedance and open circuit potential is observed, indicating the protective nature of the coating. During the 7 days of immersion, the electrolyte's conductivity remains constant at 1.8 µS cm⁻¹, in agreement with the absence of any coating dissolution. The protective nature of the multilayer is attributed to its dense structure (Fig. 6). This is likely to hinder the permeation of the electrolyte into the coating and the transport of potentially released metal ions (Fe, Ni, Cr) from the SS 316L substrate to the solution, as mentioned above. Fig. 4(b, c) reports the ICR of the material after exposure (sample E'). The ICR also remains low (3.01 mΩ cm²), confirming the coating's performance over the 168 h of the experiment. Similar tests were performed during three months of exposure. Fig. 1 of the supplementary data compares the SEM images before and after exposure. As seen, no sign of corrosion or delamination of the coating exist demonstrating the stability of the

coating over few months. As observed in Fig. 8(a and b), the low conductivity of the dilute H₂SO₄ solution does not allow for an accurate measurement of the process occurring at the electrode/electrolyte interface. The electrolyte resistance dominates in the Bode plots. To address this issue, the multilayer was exposed to a more conductive solution (2.7 mS cm⁻¹) while keeping the same pH (5.5). As shown in Fig. 8(a and b), the bode plots reveal no noticeable change in impedance over 7 days of exposure. This confirms the electrochemical stability of the film, indicating its role as an anti-corrosion barrier. To investigate the interest of the multilayer versus single layers in terms of stability, the same tests were also performed with the single-layer 200 nm TiOxNy-a coated steel (Sample A). A high stability was also observed in this case, indicating that the added value of the multilayer approach consists mainly in the optimization of the ICR.

## 4. Conclusion

In the context of reducing capital expenditures, titanium/titanium oxynitride thin-film coatings deposited by PVD on stainless steel are shown to be promising alternatives to the PGM-coated titanium currently used in the industry as bipolar plates in PEM water electrolysis systems. Indeed, under *ex situ* BPP-like experimental conditions (while taking into account ionic decoupling), TiOxNy layers demonstrated good chemical stability and effective prevention of metal leaching from 316L stainless steel at 60 °C for 168 h. However, despite their adequate electrical conductivity, single TiOxNy layers coatings suffer from a high interfacial contact resistance, which is detrimental to PEMWE performance. To overcome this limitation, a multilayer structure composed of TiOxNy/TiOxNy was developed, significantly reducing ICR values down to 1.5 mΩ cm², below the targeted value of 10 mΩ cm². AFM measurements indicate that mechanical properties, particularly hardness, influence the ICR, suggesting that optimized thin film design and

<figure><img src="image_9.png" /><figcaption>**Fig. 7.** EIS investigation of uncoated 316L in diluted H₂SO₄ solution (pH = 5.5) at 60 °C for 168 h. (a, b) Bode plots as function of the exposure time; (c) open circuit potential.</figcaption></figure>
<figure><img src="image_10.png" /><figcaption>**Fig. 8.** Electrochemical investigation of 600 nm-TiOxNy/Ti/TiOxNy multilayer-coated SS 316L in dilute H2SO4 solution (pH = 5.5), at the OCP and 60 °C. Two different electrical conductivities of the solution were used. (a, b) Bode plots, (c) open circuit potential.</figcaption></figure>

multilayer architectures can effectively mitigate resistance issues. This finding paves the way for new approaches aiming at the development of novel protective coatings. Future work will focus on integrating stainless steel BPPs coated with the TiOxNy/Ti/TiOxNy multilayer into actual PEMWE systems to assess their in situ protective performance under operational conditions.

### CRediT authorship contribution statement

**Khaoula Chergui:** Writing – review & editing, Writing – original draft, Visualization, Validation, Investigation, Formal analysis, Data curation.
**Benoit Lescop:** Writing – review & editing, Validation, Investigation.
**Michel Prestat:** Writing – review & editing, Writing – original draft, Validation, Supervision, Methodology, Investigation, Conceptualization.
**Geoffrey Ringot:** Validation, Investigation.
**Flavien Vucko:** Writing – review & editing, Validation, Investigation.
**Valérie Demange:** Writing – review & editing, Validation, Supervision, Methodology, Investigation.
**Ludivine Rault:** Writing – review & editing, Validation, Investigation.
**Francis Gouttefangeas:** Investigation.
**Loïc Joanny:** Investigation.
**Sophie Ollivier:** Investigation.
**Michael Walls:** Writing – review & editing, Validation, Methodology, Investigation.
**Gaël Maranzana:** Writing – review & editing, Validation, Methodology, Investigation.
**Divino Salvador Ramirez-Rico:** Validation, Investigation.
**Loïc Assaud:** Writing – review & editing, Validation, Supervision, Project administration, Methodology, Investigation, Funding acquisition.
**Stéphane Rioual:** Writing – review & editing, Writing – original draft, Validation, Supervision, Methodology, Investigation.

### Declaration of competing interest

The authors declare the following financial interests/personal relationships which may be considered as potential competing interests: Rioual Stephane reports financial support was provided by French National Research Agency. If there are other authors, they declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

### Acknowledgments

The authors gratefully acknowledge the financial support of the French National Research Agency (agreement ANR-22-PEHY-0009 under the France 2030 program and agreement ANR-22-CE93-0008-01).

### Appendix A. Supplementary data

Supplementary data to this article can be found online at https://doi.org/10.1016/j.ijhydene.2025.151650.

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