# Hydroxyl transport mechanisms upon passivation of Cr-Fe-Co-Ni-Mo multi-principal element alloy surfaces investigated by isotopic labelling

Xueying Wang, Dimitri Mercier, Sandrine Zanna, Antoine Syeux, Loïc

Perrière, Mathilde Laurent-Brocq, Ivan Guillot, Vincent Maurice, Philippe
Marcus

► To cite this version:

Xueying Wang, Dimitri Mercier, Sandrine Zanna, Antoine Syeux, Loïc Perrière, et al. Hydroxyl transport mechanisms upon passivation of Cr-Fe-Co-Ni-Mo multi-principal element alloy surfaces investigated by isotopic labelling. Applied Surface Science, 2024, 655, pp.159558. 10.1016/j.apsusc.2024.159558. hal-04439901

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Submitted on 5 Feb 2024

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# Hydroxyl transport mechanisms upon passivation of Cr-Fe-Co-Ni-Mo multi-principal element alloy surfaces investigated by isotopic labelling

Xueying WANG¹, Dimitri MERCIER¹*, Sandrine ZANNA¹, Antoine SEYEUX¹, Loïc PERRIERE², Mathilde LAURENT-BROCQ², Ivan GUILLOT², Vincent MAURICE¹, Philippe MARCUS¹*

¹ PSL Research University, CNRS - Chimie ParisTech, Institut de Recherche de Chimie Paris, Physical Chemistry of Surfaces Group, 11 rue Pierre et Marie Curie, 75005 Paris, France

2 Université Paris Est Creteil, CNRS, ICMPE, UMR7182, Thiais F-94320, France

*Corresponding authors: D. Mercier (dimitri.mercier@chimieparistech.psl.eu), P. Marcus (philippe.marcus@chimieparistech.psl.eu).

## Abstract

Time-of-flight secondary ion mass spectrometry (ToF-SIMS) combined with deuterium labelling (D₂O), and X-ray photoelectron spectroscopy (XPS) were applied to investigate the interfacial transport mechanisms of hydroxyls between aqueous electrolyte and outer part of the protective oxide film during passivation of Cr₁₅Fe₁₀Co₅Ni₆₀Mo₁₀ multi-principal element alloy surfaces. A ToF-SIMS signal treatment methodology was developed to retrieve the in-depth distribution of deuterated hydroxides in the nanometer-thick surface oxide film. After immersion at free potential or anodic passivation in acidified D₂O, the initial bilayer structure is retained. The use of heavy water electrolyte has no effect on the composition alterations induced by passivation. Deuterated species were detected mostly in the outer layer of the surface oxide film. Pure inward hydroxyl diffusion from electrolyte/oxide to oxide/metal interface is excluded from being primarily responsible for oxide film growth induced by anodic passivation. The hydroxyls from the dissolving hydroxide species are retained in the oxide film and participate in forming Cr hydroxide, thus reducing the hydroxyl surface uptake from the electrolyte. This work provides deeper insight into the mechanisms of oxide growth and corrosion protection induced by anodic passivation.

**Keywords:** ToF-SIMS; Deuterium labelling; Corrosion protection; Passive oxide film; Hydroxyl transport mechanisms; Multi-principal element alloy
# 1 Introduction

Passivation is one of the most effective ways to protect alloys against corrosion. Passivation consists in the spontaneous formation of a self-repairing surface oxide film preventing the metallic substrate from being exposed to the environment. To be protective, the oxide film should be either thermodynamically stable or dissolving kinetically too slowly in the given conditions. Cr is one of the widely used passivating elements for endowing alloys with passivity. Its high oxygen affinity[1,2] ensures the spontaneous formation and self-repair of a Cr oxide layer; and the compactness and slow dissolution rate of Cr oxide [3,4] ensures the protection effect of the passive film. Thus, the passive state of Cr-containing alloys is a steady state in which oxide formation and dissolution occur simultaneously but very slowly, as described in various oxide growth models [5–10].

In view of the important role of passive oxide films in impeding corrosion processes, their composition, structure, and morphology have been intensively studied using techniques such as X-ray photoelectron spectroscopy (XPS) [11,12], time-of-flight secondary ion mass spectrometry (ToF-SIMS) [13,14], scanning tunneling microscopy (STM) [15–18], and transmission electron microscopy (TEM) [19,20]. Many works have revealed that protective oxide films formed on Cr-containing Fe- or Ni-based alloys are extremely thin with a thickness range of about 1-3nm and present an in-depth stratified structure[13,21–24] and lateral heterogeneities[2,25,26]. Generally, the inner part of the oxide film consists mainly of Cr oxide with a minor amount of other elements, and it is considered as the barrier layer endowing the covered
surface with high corrosion resistance[13,14,21,22,25]. The outer part of the oxide film
can contain a mixture of oxides and hydroxides (Cr, Fe, Mo, Ni) depending on the
substrate alloy composition[13,14,21,22,25]. Therefore, a simplified bilayer model
with an inner layer enriched in Cr and an outer layer with diverse species is usually
used for illustrating the oxide film formed on Fe- or Ni-based Cr-containing
alloys[13,14,22,25]. The content of Cr and Mo species in the oxide film are always
higher than their nominal content in the bulk alloy[12,21,25,27].

With the idea to design new alloys, multi-principal element alloys (MPEAs),
forming a new class of alloys composed of several major elements instead of only one
base element, have been developed over the last two decades[28,29]. The compositional
flexibility endows MPEA systems with huge potential in property optimization, which
triggered intensive research soon after their emergence. For designing highly corrosion-
resistant MPEAs, CrFeCoNiMo is one of the prominent and most explored alloy
systems[30-34]. Studies have proven the feasibility of producing single-phase fcc
CrFeCoNiMo MPEAs with Mo content up to 15 at. % (vs. 2-2.5 at. % in 316L stainless
steel), and have demonstrated the high corrosion resistance of such alloys[30-34].
Detailed surface analysis revealed that the surface oxide films formed on these MPEAs
also present the typical stratified structure with Cr oxide concentrated in the inner part
and other (hydr)oxides in the outer part[35,36]. Recently, XPS and ToF-SIMS
measurements have suggested a combined effect of Mo and Cr for healing the Fe-rich
weak sites of the barrier layer while density functional theory (DFT) simulation
provided atomistic insight to explain the high corrosion resistance obtained with Cr,
Mo-containing MPEAs[35,37,38].

Since the protection provided by these passive oxide films involves dynamic
equilibrium between oxide formation and oxide dissolution, one may interrogate the
ion transport mechanisms within the film and across the passivated interface. Isotopic
labelling is an effective approach for investigating elemental behavior during oxide film
formation. ¹⁸O is typically used for tracing the behavior of oxygen during oxidation
[39–42]. Correspondingly, deuterium (D) can be used as isotopic marker for hydrogen
or hydroxides. However, most hydrogen-related isotopic studies for alloys were
launched in response to industrial demands concerning hydrogen embrittlement [43–
45], focusing mainly on revealing the distribution of hydrogen atoms and their
interactions with the bulk alloy microstructure[44,46–49]. In contrast, refined
characterization of hydrogen and/or hydroxyl transport behavior between electrolyte
and surface oxide film[47] for a fundamental understanding of the oxide film built-up
process[43] is less emphasized. This relative lack of attention may be because
hydroxides exist mainly in the outer part of passive films, while the high corrosion
resistance is typically attributed to the inner barrier oxide layer [3,4]. Besides, the
thinness of the oxide film, the difficulty in detecting hydrogen, and the residual
hydrogen in ultra-high vacuum (UHV) analytical chambers[43] make such study
difficult.

Being in direct contact with the external medium, the behavior of the outer oxide
layer is essential for understanding the ion transport mechanisms across the passivated
interface and must play a role in the penetration of aggressive species in the oxide film.
Several published works have already demonstrated the existence of deuterium/hydrogen in the oxide film [43,47,50] and an inward OH diffusion through oxide film in pressurized water reactors primary water, which is high-temperature high-pressure[43,50]. However, such studies are still missing for passivation mechanisms at room temperature, and the chemical environment of hydrogen atoms was not specified.

In this context, this work aims to investigate the ion transfer mechanisms between electrolyte and surface oxide film during dissolution and electrochemical passivation of single-phase fcc Cr<sub>15</sub>Fe<sub>10</sub>Co<sub>5</sub>Ni<sub>60</sub>Mo<sub>10</sub> (at.%) MPEA surfaces in sulfuric acid solution at room temperature. ToF-SIMS depth profiling using deuterium labelling for tracing hydroxyl, and XPS and ICP-OES analyses for monitoring cations were combined. Due to the similarity in mass of one deuterium atom D and two hydrogen atoms 2H, signals representing deuterated hydroxides are seriously overlapped with 2H-containing signals representing non-deuterated hydroxides at the current resolution of 8000-10000 (m/Δm for <sup>28</sup>Si) of ToF-SIMS analysis. Therefore, a methodology for extracting the deuterated hydroxide signals was developed and adopted in this work.

This methodology is an improvement for isotopic study applied to nanometer-thick surface oxide films, enabling the investigations of diverse deuterated hydroxides. Based on the obtained results, the deuterium surface uptake during dissolution and passivation in D<sub>2</sub>O is discussed, providing deeper insight into the cation and hydroxyl transport process across the oxide/electrolyte interface during electrochemical passivation at room temperature.
# 2 Experimental and methods

## 2.1 Sample preparation

Homogeneous single-phase fcc Cr<sub>15</sub>Fe<sub>10</sub>Co<sub>5</sub>Ni<sub>60</sub>Mo<sub>10</sub> (at.%) MPEA samples with a grain size of 61 ± 5 µm were studied in this work. The manufacturing of the alloy has been reported elsewhere[34]. The specimen surfaces were mechanically polished successively with 1200, 2400 and 4000 grade SiC paper, followed by diamond paste finish of 3, 1 and 0.25 µm. Samples were then ultrasonically cleaned for 5 min in acetone, ethanol and deionized water (Merk Millipore, resistivity of 18.2 MΩ.cm), and dried in compressed air flow.

The anodically passivated surfaces were prepared at room temperature with a 3-electrode cell comprising a saturated mercury sulfate reference electrode (MSE, $E_{SHE} = E_{MSE} + 0.64$ V), a Pt wire counter electrode and the MPEA sample surface constituting the working electrode. The working area was delimited to 1.13 cm² by a Viton O-ring. A VersaSTAT4 potentiostat (AMETEK) was used as controller. The electrolyte for passivation was 5 mL 0.05 M H<sub>2</sub>SO<sub>4</sub> solution prepared from sulfuric acid (VMR, reagent ACS grade) and deionized water or D<sub>2</sub>O (Eurisotop, deuterium purity of 99.97%). Anodic passivation was performed by stepping the potential to 0 V<sub>MSE</sub> for 1 h after 15 min of stabilization at Open Circuit Potential (OCP). The surfaces passivated at OCP, hereafter referred to as immersed surfaces, were prepared in either 5 mL of pure water or 5 mL 0.05 M H<sub>2</sub>SO<sub>4</sub> solution, one using deionized H<sub>2</sub>O and the other one using D<sub>2</sub>O. The duration of the immersions was 1 hour and 15 min, corresponding to the total time that the passivated surfaces were exposed to electrolyte. Prior to anodic passivation
or OCP immersion, the electrolytes were deaerated by bubbling Ar gas for 30 min to minimize dissolved oxygen. After immersion or passivation in sulfuric acid solution, the samples were rinsed with deionized water during several seconds and dried with compressed air. The samples were then introduced into XPS or ToF-SIMS UHV chamber for analysis. The air exposure duration during the transfer process was about 20 min.

## 2.2 Solution analysis

The sulfuric acid solution was analyzed before passivation (i.e., after 15 min of stabilization at OCP) and after passivation by ICP-OES to calculate the quantity of dissolved cations during passivation. The measurements were performed using an ICP-OES iCAP 6300 (Thermo Fisher Scientific) after passivation in 0.05 M H₂SO₄ prepared with H₂O. The solutions used for ICP-OES analysis were acidified with 5 wt.% HNO₃ solution. Calibration solutions from 0.03 ppm to 0.5 ppm were prepared for Cr, Fe, Co, Ni and Mo with corresponding standard solution of 1000 µg/mL (Chem-Lab NV), 5 wt.% HNO₃ solution for acidification, and 0.05 M H₂SO₄ for ensuring the same matrix effect as the sample solutions. The quantity of dissolved cations was normalized to the working electrode area (1.13 cm²).

## 2.3 Surface analysis

The stratified structure and composition of the surface oxide films after immersion at OCP or anodic passivation were determined by XPS and ToF-SIMS analysis.

XPS analysis was carried out in a Thermo Electron ESCALAB 250Xi spectrometer
with a pressure lower than 10⁻¹⁰ mbar in the analysis chamber. A monochromatic Al Kα X-ray source (hν = 1486.6 eV) was used. High-resolution spectra of the Cr 2p, Fe 2p, Co 2p, Ni 2p, Mo 3d, S 2p, O 1s, and C 1s core levels were recorded with a pass energy of 20 eV and a step size of 0.1 eV. The take-off angle of the photoelectrons was 90°. The fitting of the spectra was performed using Casa XPS software. The decomposition procedure and parameters were the same as previously established for the Cr₁₅Fe₁₀Co₅Ni₆₀Mo₁₀ MPEA, reported in details elsewhere [35,37,51]. In summary, Shirley type background was used and constraints on binding energy (BE), full width at half maximum (FWHM), lineshape, intensity ratio between spin-orbit doublets, and intensity ratio between photoelectron peak and Auger peak were applied for curve fitting. Reference lineshapes obtained on corresponding pure metals or pure Cr oxide were used to consider the asymmetry of metallic peaks and the multiplet splitting structure of Cr oxide peak. Other oxide or hydroxide peaks were fitted with symmetric Gaussian/Lorentzian product lineshape GL(x), where x defines the percentage of Lorentzian character. For the Cr 2p, Fe 2p, Co 2p, Mo 3d core level spectra, the entire region containing both 2p₃/₂ and 2p₁/₂ (or 3d₅/₂ and 3d₃/₂ for Mo 3d) peaks was decomposed due to the overlap between the spin-orbit doublets. For the Ni 2p spectra, only the Ni 2p₃/₂ region was decomposed since it is not overlapping with Ni 2p₁/₂ peaks.

ToF-SIMS depth profile elemental analysis was performed in an IONTOF ToF-SIMS 5 spectrometer with a pressure of 10⁻⁸ mbar or lower and operated in dual-beam operation mode. A 25 keV Bi⁺ primary LMIG source delivering 1.2 pA target current over a 100 × 100 µm² area was used for static SIMS analysis. The static analysis was
interlaced with sputtering realized by a 0.5 keV $Cs^+$ sputter beam delivering 17 nA target current over a 300 × 300 µm² area. The two beams were at an incidence angle of 45° with respect to sample surface and well-aligned to ensure analysis from the center of the sputtered crater. The time acquisition interval of the data was 1.3 s. Negative ions were recorded because they are more sensitive to ion fragments coming from oxide matrices. For each surface, at least two depth profiles in two different zones were recorded to verify the uniformity of the surface. Data acquisition and processing were performed with the IONSPEC 7.2 software.

$D^-$, $OD^-$, $CrO_3D^-$, $NiO_2D^-$, and $^{100}MoO_4D^-$ were selected as representative ion fragments for D, OD, and deuterated Cr, Ni, Mo hydroxides, respectively. However, the similarity in mass of one D (2.0141 amu) and two H (2.0157 amu) results in the overlap between the hydroxide signals containing one D ($CrO_3D^-$, $NiO_2D^-$, or $^{100}MoO_4D^-$) and the signals containing $H_2$ ($CrO_3H_2^-$, $NiO_2H_2^-$, or $^{100}MoO_4H_2^-$) at the mass resolution of 8000-10000 ($\frac{m}{\Delta m}$ for $^{28}$Si). Therefore, based on the findings of changeless intensity ratios $\frac{MO_xH_2^-}{MO_xH^-}$ for M = Cr, Ni or $^{100}Mo$ for each kind of surface, and the consistency in nature between the deuterated or non-deuterated oxide films, equation (1) was used to extract the signal for deuterated hydroxides:

$$MO_xD^- = (MO_xH_2^-)_{raw} - MO_xH^- \times \left(\frac{MO_xH_2^-}{MO_xH^-}\right)_{reference} \quad (1)$$

Where:

* $MO_xD^-$ is the characteristic signal for deuterated hydroxide calculated for the deuterated surface.

* $(MO_xH_2^-)_{raw}$ is the raw signal measured on the deuterated surface. It is in
fact the sum of deuterated hydroxide signal $MO_xD^-$ and non-deuterated hydroxide signal $MO_xH_2^-$.

* $MO_xH^-$ is the characteristic signal for non-deuterated hydroxide measured on the deuterated surface.

* $(\frac{MO_xH_2^-}{MO_xH^-})_{reference}$ is the reference ratio obtained on the non-deuterated surface, prepared following the same procedure but in corresponding non-deuterated electrolyte.

In order to minimize the overlap with other signals, a specified mass interval was taken for each ion fragment instead of the whole peak detected in the mass spectrum. The mass intervals were maintained the same throughout this work in order to ensure the validity of the reference ratios.

Due to the existence of multiple laboratory-scale stable Mo isotopes of important relative abundance (9.2 % $^{94}$Mo, 15.9 % $^{95}$Mo, 16.7 % $^{96}$Mo, 9.6 % $^{97}$Mo, 24.3 % $^{98}$Mo, 9.7 % $^{100}$Mo), there were overlaps between hydroxylated fragments of lighter Mo isotope and non-hydroxylated Mo fragments of heavier Mo isotope, for example $^{98}$MoO$_4$H$_2^-$ and $^{100}$MoO$_4^-$.$^{1}$ Therefore, $^{100}$Mo, the heaviest stable Mo isotope, was selected instead of the principal isotope $^{98}$Mo so that no overlap with heavier isotope could occur.

# 3 Results and Discussion

## 3.1 Deuterium detection in the oxide film

Two types of ToF-SIMS representative signals of interest were distinguished for
characterizing D-containing species in the oxide films covering the immersed and passivated Cr<sub>15</sub>Fe<sub>10</sub>Co<sub>5</sub>Ni<sub>60</sub>Mo<sub>10</sub> surfaces prepared in acidified D<sub>2</sub>O. Figure 1 presents raw or treated ToF-SIMS negative ion depth profiles obtained on non-deuterated or deuterated passivated surfaces as an example for illustrating the detection of D-containing species in the surface oxide film.

<figure><img src="image_2.png" /><figcaption>Figure 1 ToF-SIMS negative ion depth profiles of raw H<sup>-</sup> and <sup>18</sup>OH<sup>-</sup> signals (a), and of the raw intensity ratio CrO<sub>3</sub>H<sub>2</sub><sup>-</sup>/CrO<sub>3</sub>H<sup>-</sup> (b) for anodically passivated Cr<sub>15</sub>Fe<sub>10</sub>Co<sub>5</sub>Ni<sub>60</sub>Mo<sub>10</sub> surfaces prepared in non-deuterated or deuterated solutions.</figcaption></figure>

Figure 1(a) shows the depth profiles of the first type of characteristic signals, H<sub>2</sub><sup>-</sup> [43,47,52] and OH<sub>2</sub><sup>-</sup> [47]. The ordinate of the plots is the intensity of the signals normalized to the nickel signal in the alloy (I/I<sub>Ni2(bulk)</sub>). The abscissa is the sputtering time, reflecting the depth from the surface. The dashed lines mark the positions of the outer/inner oxide layer interface (at around 18 s) and oxide/alloy interface (at around 38 s). The definitions of the interfaces are explained in section 3.2. Since D and 2H are not distinguishable at the present mass resolution, the measured raw H<sub>2</sub><sup>-</sup> and OH<sub>2</sub><sup>-</sup> signals were in fact (H<sub>2</sub><sup>-</sup>+D<sup>-</sup>) and (OH<sub>2</sub><sup>-</sup>+OD<sup>-</sup>), respectively. However, as shown in Figure 1(a), the intensities of raw (H<sub>2</sub><sup>-</sup>+D<sup>-</sup>) and (OH<sub>2</sub><sup>-</sup>+OD<sup>-</sup>) signals recorded on the deuterated surface were much higher than those recorded on the non-deuterated surface. This significant difference in intensity suggests that there was deuterium originating
from the electrolyte in the oxide film after passivation, and that in practice the H₂⁻ and
OH₂⁻ contributions to the signals are negligible compared to D⁻ and OD⁻. Therefore, the
recorded raw (H₂⁻+D⁻) and (OH₂⁻+OD⁻) signals were used directly as characteristic ions
for reconstructing the distributions of D and OD hereafter.

The second type of characteristic ions consists of CrO₃D⁻, FeO₂D⁻, NiO₂D⁻, and
¹⁰⁰MoO₃D⁻, representing deuterated hydroxides of Cr, Fe, Ni, and Mo, respectively.
These signals also overlap with corresponding 2H-containing signals MOₓH₂⁻ in the
mass spectra. These MOₓH₂⁻ signals, characterizing the non-deuterated hydroxides, are
not negligible according to their intensity measured on non-deuterated surfaces.
However, it was observed that the intensity ratio $\frac{\text{raw CrO}_3\text{H}_2^-}{\text{raw CrO}_3\text{H}^-}$ for one surface is
consistent and even insensitive to experimental conditions. Besides, the $\frac{\text{raw CrO}_3\text{H}_2^-}{\text{raw CrO}_3\text{H}^-}$
ratio obtained for deuterated surfaces is always higher than that for corresponding non-
deuterated surfaces in the superficial region where hydroxides existed. The extra part
should result from deuterated species. Figure 1(b) presents the depth profile of the
$\frac{\text{raw CrO}_3\text{H}_2^-}{\text{raw CrO}_3\text{H}^-}$ ratio for the deuterated passivated surface as an example. The difference
between the two curves, marked with gray shadowing, must be related to the CrO₃D⁻
contribution to the signal which indicates the existence of deuterated Cr hydroxide. At
deeper location in the oxide inner layer and bulk alloy regions, the fluctuating signals
resulted only from background noise enlarged by the division operation and are not
characteristic of the species distribution. Similar higher $\frac{\text{raw MO}}{\text{raw CrO}}$ ratio in the
superficial region on deuterated surface was also observed for Fe, Ni and Mo.

Figure 1 proves with the D⁻ and OD⁻ signals that there is a certain quantity of
deuterium in the surface oxide film after passivation in deuterated electrolyte. It also elucidates with raw intensity ratios a potential way to characterize specific deuterium species by ToF-SIMS. However, since isotope distribution should tend towards equilibrium, it cannot be concluded whether the deuterium surface uptake was principally due to passivation or to isotopic exchange[53]. Therefore, the same analyses on Cr<sub>15</sub>Fe<sub>10</sub>Co<sub>5</sub>Ni<sub>60</sub>Mo<sub>10</sub> surfaces immersed in pure D<sub>2</sub>O or 0.05 M H<sub>2</sub>SO<sub>4</sub> prepared with D<sub>2</sub>O are necessary to differentiate the effects of passivation from other parasitic processes. For the surfaces immersed in pure or acidified D<sub>2</sub>O, direct usability of raw (H<sub>2</sub><sup>-</sup>+D<sup>-</sup>) and (OH<sub>2</sub><sup>-</sup>+OD<sup>-</sup>) signals for representing D and OD and the higher (M=Cr, Ni or <sup>100</sup>Mo) ratio in the presence of deuterium were also observed.

## 3.2 The structure and composition of the oxide films

The prerequisite to using the $\frac{MO_xH_2^-}{MO_xH^-}$ ratio obtained on non-deuterated surfaces as a reference for extracting D-containing signals for deuterated surfaces is to verify that the $\frac{MO_xH_2^-}{MO_xH^-}$ ratio obtained on non-deuterated surfaces is still valid for the corresponding species on deuterated surfaces. The validation relied on examining the consistency of the oxide film nature in the absence and presence of deuterium. The structure and the composition of non-deuterated oxide films and corresponding deuterated oxide films were investigated and compared to examine whether the reactivity difference between D and H [54] leads to modifications on the oxide film except for the substitution of H by D.

The composition of the oxide film was determined by XPS analysis using the fitting method described previously[35]. Figure 2 shows the XPS high-resolution spectra of
the Cr 2p, Ni 2p<sub>3/2</sub>, Fe 2p, Mo 3d, and Co 2p core level regions recorded on non-deuterated or deuterated passivated surfaces as an example for illustrating the consistency of identified species.

<figure><img src="image_3.png" /><figcaption>Figure 2 XPS Cr 2p (a), Ni 2p<sub>3/2</sub> (b), Fe 2p (c), Mo 3d (d) and Co 2p (e) core level spectra for Cr<sub>15</sub>Fe<sub>10</sub>Co<sub>5</sub>Ni<sub>60</sub>Mo<sub>10</sub> surfaces anodically passivated in 0.05 M H<sub>2</sub>SO<sub>4</sub> solution prepared with H<sub>2</sub>O or D<sub>2</sub>O.</figcaption></figure>

The decomposition of these XPS spectra identified the same species for both surfaces, namely Cr(III) oxide and hydroxide, Fe(II) and Fe(III) (hydr)oxides, Ni hydroxide, Mo(IV) and Mo(VI) (hydr)oxides. The same components were also identified for the Cr<sub>15</sub>Fe<sub>10</sub>Co<sub>5</sub>Ni<sub>60</sub>Mo<sub>10</sub> surfaces immersed in pure or acidified H<sub>2</sub>O (D<sub>2</sub>O).

Figure 3 displays the depth profiles illustrating the structure of the surface layers for surfaces immersed in pure H<sub>2</sub>O (Figure 3(a)), pure D<sub>2</sub>O (Figure 3(b)), acidified H<sub>2</sub>O (Figure 3(c)), acidified D<sub>2</sub>O (Figure 3(d)), or anodically passivated in acidified H<sub>2</sub>O
(Figure 3(e)) or acidified D₂O (Figure 3(f)). The representative secondary ions determined in previous work for the Cr₁₅Fe₁₀Co₅Ni₆₀Mo₁₀ MPEA were used for reconstructing the depth profiles [35]. The H-free signals CrO⁻ (67.9354 amu), FeO₂⁻ (87.9248 amu), MoO₂⁻ (129.8952 amu), and MoO₃⁻ (145.8902 amu) represent Cr, Fe, Mo oxides, respectively. The Ni₂⁻ signal (115.8707 amu) represents metallic nickel. The CrO₃Hₓ⁻ and NiO₂Hᵧ⁻ series of signals is typically used as characteristic ions for Cr hydroxide and Ni hydroxide in literature[13,14,35]. Here, CrO₃H₂⁻ (101.9409 amu) and NiO₂H⁻ (88.9330 amu) are shown. The dashed lines indicate the inner/outer oxide layer interface and the oxide/metal interface, located using the convention of the position where the CrO⁻ and Ni₂⁻ signals reach 80% of their maximum intensity[35,37]. The profiles have been normalized by the intensity of Ni₂⁻ in the bulk alloy.

<figure><img src="image_4.png" /></figure>

<figure><img src="image_5.png" /></figure>

<figure><img src="image_6.png" /></figure>

<figure><img src="image_7.png" /></figure>

<figure><img src="image_8.png" /></figure>

<figure><img src="image_9.png" /></figure>
**Figure 3** ToF-SIMS negative ion depth profiles for the Cr<sub>15</sub>Fe<sub>10</sub>Cos<sub>5</sub>Ni<sub>60</sub>Mo<sub>10</sub> surfaces immersed at OCP in H<sub>2</sub>O (a) or D<sub>2</sub>O (b), immersed at OCP in 0.05 M H<sub>2</sub>SO<sub>4</sub> prepared with H<sub>2</sub>O (c) or D<sub>2</sub>O (d), or anodically passivated in 0.05 M H<sub>2</sub>SO<sub>4</sub> prepared with H<sub>2</sub>O (e) or D<sub>2</sub>O (f).

As shown in Figure 3(a)-(f), among the characteristic signals of (hydr)oxides, two distinct regions exist. CrO<sub>3</sub>H<sub>2</sub><sup>-</sup>, NiO<sub>2</sub>H<sup>-</sup>, FeO<sub>2</sub><sup>-</sup>, MoO<sub>2</sub><sup>-</sup>, and MoO<sub>3</sub><sup>-</sup> signals always peak at sputtering times earlier than the CrO<sup>-</sup> signal, showing that the hydroxide of Cr, Ni, and (hydr)oxides of Fe, Mo, are concentrated at a position more superficial than Cr oxide. As explained in previous work, the stratified distribution of these species is consistent with a bilayer structure of the oxide film. The inner layer of the oxide film is enriched in Cr oxide, and the outer layer is composed of Cr, Fe, Mo oxides or hydroxides.

For each of the three pairs of surfaces, no significant modification on the stratified structure of the oxide film was observed between non-deuterated (Figure 3a, c, e) or deuterated (Figure 3b, d, f) state, indicating that the deuteration of the electrolyte did not lead to structural modifications of the surface oxide film.

Based on the bilayer model, quantification was then carried out using XPS intensities and the intensity attenuation equations for an electron flux passing through surface layers [22,55]. The values of photoionization cross-section of each element for Al Kα emission at 1486.6 eV and the inelastic mean free paths of the emitted photoelectrons were taken from J. H. Scofield [56] and S. Tanuma *et al.* [57], respectively. The alloy density is 8.3 g.cm<sup>-3</sup> and an assumed oxide film density of 4.5 g.cm<sup>-3</sup> was used. The XPS intensity of Cr(III) oxide component was attributed to the inner layer and the intensities of the other species to the outer layer[35]. The calculated
composition of the outer oxide layer and thickness of the inner and the outer layers are compiled in Table 1.

<table><caption>Table 1 Composition and thickness of the surface oxide films after immersion of the Cr<sub>15</sub>Fe<sub>10</sub>Co<sub>5</sub>Ni<sub>60</sub>Mo<sub>10</sub> surfaces in pure H<sub>2</sub>O or D<sub>2</sub>O, in acidified H<sub>2</sub>O or D<sub>2</sub>O, or anodic passivation in acidified H<sub>2</sub>O or D<sub>2</sub>O. Values were obtained from the XPS intensities using a bilayer model of intensity attenuation.</caption><thead><tr><th rowspan="2"></th><th colspan="5">Cationic constitution of the outer layer</th><th colspan="2">Thickness (±0.1 nm)</th></tr><tr><th>Cr (±2 at%)</th><th>Fe (±3 at%)</th><th>Co (±3 at%)</th><th>Ni (±2 at%)</th><th>Mo (±2 at%)</th><th>Outer layer</th><th>Inner layer</th></tr></thead><tbody><tr><td>Native oxide film[35]</td><td>17</td><td>13</td><td>3</td><td>46</td><td>21</td><td>1.3</td><td>0.2</td></tr><tr><td>Immersion in H<sub>2</sub>O</td><td>43</td><td>10</td><td>2</td><td>22</td><td>23</td><td>1.5</td><td>0.2</td></tr><tr><td>Immersion in D<sub>2</sub>O</td><td>41</td><td>11</td><td>2</td><td>24</td><td>22</td><td>1.5</td><td>0.3</td></tr><tr><td>Immersion in 0.05 M H<sub>2</sub>SO<sub>4</sub> (H<sub>2</sub>O)</td><td>32</td><td>8</td><td>2</td><td>29</td><td>29</td><td>1.3</td><td>0.2</td></tr><tr><td>Immersion in 0.05 M H<sub>2</sub>SO<sub>4</sub> (D<sub>2</sub>O)</td><td>29</td><td>8</td><td>3</td><td>30</td><td>30</td><td>1.3</td><td>0.2</td></tr><tr><td>Passivation in H<sub>2</sub>O with 0.05 M H<sub>2</sub>SO<sub>4</sub> [35]</td><td>45</td><td>6</td><td>2</td><td>16</td><td>31</td><td>1.8</td><td>0.3</td></tr><tr><td>Passivation in D<sub>2</sub>O with 0.05 M H<sub>2</sub>SO<sub>4</sub></td><td>42</td><td>7</td><td>4</td><td>13</td><td>34</td><td>1.7</td><td>0.3</td></tr></tbody></table>

The composition and thickness of the oxide films obtained for surfaces prepared following the same procedure with or without deuteration reveal that the deuteration of the electrolyte did not lead to any significant modifications in the composition nor thickness.

The structural consistency (supported by the depth profiles in Figure 3) and the compositional consistency (proven by Figure 2 and Table 1) between the non-
deuterated and deuterated surfaces suggest that the use of D₂O in place of H₂O does not
induce significant modification of the passivation mechanism and kinetics, possibly due
to the extremely small depth dimension of the passivation reaction. The usability of
non-deuterated surfaces as a reference for extracting D-containing signals is thus
justified.

Regarding surface composition after different preparation procedures, comparison
between the native oxide film and the oxide film after immersion in pure H₂O or D₂O
reveals that the immersion in water led to slight thickening and Cr enrichment of the
oxide film, with a decrease in Ni content. The decrease in Ni content and the increase
in Cr content can be explained by the solubility of Ni(OH)₂ in water (0.0015 g/L)[58]
and the stability of Cr hydroxide according to the Pourbaix diagram. However, after
immersion in acid for the same duration, the Cr enrichment was less pronounced, while
Mo species were further enriched. The impediment to Cr enrichment may be related to
the dissolution of Cr hydroxide in acidic media. Both the decrease in equivalent oxide
film thickness in Table 1 and the interface position at earlier sputtering times in Figure
3(c and d) also confirm that dissolution occurred to the oxide film during immersion in
acid solution, which has been reported for stainless steels[21,59] and
CrFeCoNi(Mn)(RuW) MPEAs[60,61]. The Mo enrichment caused by immersion at
OCP has also been observed for NiCrFeRuMoW MPEA[60], which may be related to
the difference in stability of Mo oxides in acid compared to other oxides. Regarding
anodic passivation, evident thickening of the outer oxide layer was observed,
accompanied by pronounced Cr and Mo enrichments and Ni depletion. These
observations substantiate that the electrochemical polarization in the passive range enhances the oxide film by promoting the Cr and Mo enrichments and the Ni dissolution [12,14,21,25,59–65].

### 3.3 Deuterated signals distribution

After verifying the oxide film consistency in the absence or presence of deuterium in the solution, the $MO_xD^-$ (M = Cr, Ni, $^{100}$Mo, Fe) signals were calculated using equation (1) in Section 2.3 for the three $Cr_{15}Fe_{10}Co_5Ni_{60}Mo_{10}$ surfaces treated in deuterated environment.

Figure 4 presents the depth profiles of the $D^-$, $OD^-$, and $MO_xD^-$ signals normalized to bulk $Ni_{2}^-$ intensity obtained for the three surfaces. As explained above, the $D^-$ and $OD^-$ signals reflect the general distribution of D atoms and OD hydroxyl groups, respectively. The $CrO_3D^-$, $NiO_2D^-$, and $^{100}MoO_3D^-$ signals represent deuterated Cr hydroxide, deuterated Ni hydroxide and deuterated Mo hydroxide, respectively. The overlap in mass between $FeH_2^-$ (57.9506 amu) and $Ni^-$ (57.9353 amu), and the low Fe content in the oxide[35] made the extracted deuterated Fe signal too noisy to develop reliable discussion. Therefore, the in-depth distribution of deuterated Fe hydroxides will not be discussed. The dashed lines mark the positions of the outer/inner oxide layer interfaces and oxide/alloy interfaces defined in Figure 3(b), (d), and (f). The fluctuations of the thus normalized intensities obtained during repetitive measurements performed at different locations on the same sample were less than 8 %.
<figure><img src="image_10.png" /><figcaption>Figure 4 ToF-SIMS negative ion depth profiles of the deuterated signals for the Cr<sub>15</sub>Fe<sub>10</sub>Co<sub>5</sub>Ni<sub>60</sub>Mo<sub>10</sub> surface immersed at OCP in D<sub>2</sub>O (a), in 0.05 M H<sub>2</sub>SO<sub>4</sub> prepared with D<sub>2</sub>O (b), or passivated at 0 V<sub>MSE</sub> in 0.05 M H<sub>2</sub>SO<sub>4</sub> prepared with D<sub>2</sub>O (c).</figcaption></figure>

For the Cr<sub>15</sub>Fe<sub>10</sub>Co<sub>5</sub>Ni<sub>60</sub>Mo<sub>10</sub> surface immersed in pure D<sub>2</sub>O for 1 h 15 min (Figure 4(a)), all deuterated representative signals peak in the outer oxide layer region before reaching the inner/outer oxide interface. However, slight differences can be noticed in the sputtering time corresponding to their maximum intensity values. Among the general signals D<sup>-</sup> and OD<sup>-</sup>, the D<sup>-</sup> signal peaks at earlier sputtering time than the OD<sup>-</sup> signal. Among metal hydroxide signals, the NiO<sub>2</sub>D<sup>-</sup> signal peaks earlier than CrO<sub>3</sub>D<sup>-</sup> and <sup>100</sup>MoO<sub>4</sub>D<sup>-</sup> signals.

These depth profiles indicate that there is deuterium in the oxide film even after immersion in pure D<sub>2</sub>O, and that the deuterated species concentrate principally in the outer oxide layer. Previous works on Cr oxide [66,67] and Fe oxides[68] suggest that pure water in gaseous or liquid state can interact with metal oxides via dissociative chemisorption[67–69]. On the oxide surface, water molecules dissociate into OH and
H, followed by OH adsorption on metal cations via cation-O bonding and H adsorption on oxygen anions initially forming the oxide surface. The dissociative chemisorption of water molecules could be one reason for having deuterated species in the oxide film. Besides, considering the D and H concentration gradient between the water and the oxide film, the H/D isotopic exchange[70,71] cannot be ignored, which could be another important reason for having deuterium into the outer oxide layer.

The maximal intensity of the D⁻ signal observed at the very beginning of sputtering, earlier than for OD⁻, suggests that globally the preferential location of D atoms is more superficial than that of O, making OD fragments sputtered out later than D alone. This superficial distribution of D may indicate the existence of some protruding D terminations on the oxide surface[68].

Regarding the deuterated metal hydroxides characterized by CrO₃D⁻, NiO₂D⁻ and ¹⁰⁰MoO₄D⁻, the difference in sputtering times for reaching their maximum intensity indicates that deuterated Ni hydroxide is more superficial than deuterated Cr and Mo hydroxides. This fine stratification exists also between normal Ni, Cr, and Mo species in non-deuterated oxide film as shown in Figure 3, meaning that the deuterated hydroxides followed the same distribution as the non-deuterated species of corresponding elements in the outer oxide layer.

Figure 4(b) shows the depth profiles of deuterated species for the Cr₁₅Fe₁₀Co₅Ni₆₀Mo₁₀ surface immersed at OCP in acidified D₂O. Again, the peaks of the deuterated signals are all located in the outer oxide layer region, and the D⁻ ion profile peaks earlier than that of OD⁻. The difference in peak locations of the NiO₂D⁻,
$CrO_3D^-$, and $^{100}MoO_4D^-$ signals decreases. It is noteworthy that the relative location of the $OD^-$ peak in the outer layer becomes slightly deeper in Figure 4(b), whereas it was relatively more superficial in Figure 4(a). Besides, the intensities of all D-containing signals in Figure 4(b) are lower than in Figure 4(a).

The profiles in Figure 4(b) confirm the existence of deuterated species in the outer oxide layer of the surface oxide film after immersion at OCP in acidified deuterated water. The more superficial distribution of D than O is retained as supported by the respective location of the $D^-$ and $OD^-$ peaks, whereas the stratification between deuterated Ni, Cr, Mo hydroxides became less evident, possibly due to thinning of the oxide film as discussed above. Since the $OD^-$ profile is determined by the contribution of various (oxy)hydroxides and adsorbed hydroxyl/water, its relative deeper location in the outer layer could result from enhanced dissolution of some hydroxylated components at extreme surface in the acidified solution[66,72]. Regarding the intensity decrease compared to neutral solution, it cannot be excluded that the compositional changes induced by dissolution resulted in ion yield changes. However, the lower intensity level measured for deuterated species agrees well with the proposition that dissolution of the surface hydroxides occurred when the water was acidified.

Figure 4(c) shows the depth profiles of D-containing components for the surface anodically passivated in acidified $D_2O$. Expectedly, the profiles of the deuterated species still peak in the outer oxide layer region. However, compared to Figure 4(b), the intensities of all deuterated signals further decrease, especially that of the $NiO_2D^-$ signal. Broadening of the $CrO_3D^-$ profile is observed, whereas it is not observed for
$^{100}$MoO$_4$D$^-$.

The location of the peaks in Figure 4(c) demonstrates that deuterated species also exist for the passive oxide film and are located principally in the outer layer. The relative peak positions of the $^{100}$MoO$_4$D$^-$ and CrO$_3$D$^-$ signals indicate that the deuterated Mo species are distributed in the outer part of the outer layer of the passive film, while the deuterated Cr species are concentrated at a slightly deeper position in the outer oxide layer. The significantly lower intensity of NiO$_2$D$^-$ signal is in agreement with the diminution of Ni content observed after anodic passivation (Table 1), indicating the preferential dissolution of Ni hydroxide induced by anodic electrochemical polarization in the passive range[13,35,60,63]. No evident intensity increase was observed for all these deuterated signals after anodic passivation compared to immersion at OCP in water or acid solution. Under the assumption that the ion sputtering yields were not markedly modified, it can be inferred that the re-growth of the oxide film induced by anodic passivation is not accompanied by a pronounced surface uptake of deuterium from the electrolyte.

**3.4 Hydroxyl transport mechanisms during passivation**

To further understand the intensity variations of the deuterated-related signals, the correlation between elemental content of the surface oxide film and corresponding ToF-SIMS signal intensities was investigated. Based on the exploration and comparison of the oxide film cationic composition obtained by XPS and the ToF-SIMS signal intensities measured for the same CrFeCoNiMo MPEA surfaces after different preparation [34], a linearity was observed between the XPS-measured content of Cr
hydroxide in the outer oxide layer and the maximum intensity of the ToF-SIMS signal representing Cr hydroxide in equivalent measurement conditions. Figure 5 compares the intensity variation tendency of ToF-SIMS depth profile for non-deuterated (CrO₃H⁻ and CrO₃H₂⁻) and deuterated (CrO₃D⁻) Cr hydroxides with the content of Cr hydroxide in the outer oxide layer obtained by XPS for the three Cr₁₅Fe₁₀Co₅Ni₆₀Mo₁₀ surfaces, in order to elucidate the different behavior of OH⁻ and OD⁻ upon OCP and anodic passivation.

<figure><img src="image_11.png" /><figcaption>Figure 5 XPS-measured of Cr content in the outer oxide layer and ToF-SIMS depth profiles of representative ions for non-deuterated (CrO₃H⁻, CrO₃H₂⁻) and deuterated (CrO₃D⁻) Cr hydroxide for Cr₁₅Fe₁₀Co₅Ni₆₀Mo₁₀ surfaces treated in three deuterated solutions.</figcaption></figure>

Comparing the ToF-SIMS depth profiles after immersion at OCP in acidified D₂O to that in pure D₂O, the maximum intensity of all Cr characteristic signals CrO₃H⁻,
CrO₃H₂⁻, and CrO₃D⁻ decreased consistently with the Cr content in the outer layer measured by XPS. The ratio between maximal intensities CrO₃D⁻max/CrO₃H₂⁻max was calculated to be 0.6 for both surfaces immersed at OCP. However, after anodic passivation, a divergence can be observed between non-deuterated and deuterated Cr hydroxide. Compared to immersion at OCP, the XPS-measured Cr content in the outer layer of passive oxide film increased to 42 at. %. The maximal intensity of CrO₃H⁻ and CrO₃H₂⁻ increased correspondingly, whereas the maximal intensity of CrO₃D⁻ decreased. The maximum ratio CrO₃D⁻max/CrO₃H₂⁻max was reduced to 0.3. This discrepancy demonstrates clearly that anodic passivation resulted in a higher proportion of non-deuterated Cr hydroxide among Cr hydroxide in the final passive film.

The only source of non-deuterated hydroxyl OH⁻ in these experiments was the OH⁻ coming from hydroxide species (Ni(OH)₂, Fe(OH)₂, and Fe(OH)₃) originally present in the native oxide film. Therefore, one may ask whether the higher non-deuterated Cr hydroxide proportion is due to the limitation of dissolution upon anodic passivation or related to the ion transport mechanism. Table 2 compiles the amounts of cations measured in the electrolyte by ICP-OES after passivation, and the amount of cations contained in the native and passive oxide films measured by XPS.

<table><caption>Table 2 Quantity of cations contained in the electrolyte as measured by ICP-OES analysis after anodic passivation, and in the native and passive oxide films on Cr₁₅Fe₁₀Co₅Ni₆₀Mo₁₀ surfaces as obtained from XPS</caption><thead><tr><th rowspan="2">Cation</th><th colspan="3">Cationic amount (10<sup>-9</sup> mol·cm<sup>-2</sup>)</th></tr><tr><th>In electrolyte<br>(by ICP-OES)</th><th>In the outer layer of<br>the native oxide film<br>(by XPS)</th><th>In the outer layer of the<br>passive oxide film<br>(by XPS)</th></tr></thead><tbody><tr><td>Cr</td><td>1.7(±0.5)</td><td>0.9(±0.1)</td><td>2.6(±0.1)</td></tr><tr><td>Fe</td><td>4.1(±0.7)</td><td>0.7(±0.1)</td><td>0.5(±0.1)</td></tr><tr><td>Co</td><td>1.1(±0.2)</td><td>0.2(±0.1)</td><td>0.2(±0.1)</td></tr><tr><td>Ni</td><td>17.5(±0.7)</td><td>2.5(±0.1)</td><td>0.8(±0.1)</td></tr><tr><td>Mo</td><td>1.8(±0.4)</td><td>1.1(±0.1)</td><td>2.1(±0.1)</td></tr></tbody></table>
analysis.

The amounts of cations in the electrolyte and in the oxide films confirm again the enhanced dissolution of Ni, Fe species and the enrichment of Cr, Mo species in oxide film induced by passivation. Besides, comparing the quantities of cations dissolved into the electrolyte to those contained in the oxide films, the quantity of dissolved cations was not of minor extent. Therefore, the higher non-deuterated Cr hydroxide proportion in the passive oxide film should be related to enhanced formation of non-deuterated Cr hydroxide and/or hindered formation of deuterated Cr hydroxide upon anodic passivation.

To form new non-deuterated Cr hydroxide in deuterated solution, OH⁻ could only come from the hydroxide species initially contained in the native oxide film, i.e., in Ni hydroxide, and possibly in Fe and Mo hydroxides if there were such species. Since less Ni and Fe are measured in the final passive oxide film, Ni and Fe hydroxides could be considered as suppliers of OH⁻. Thus, it is reasonable to propose that under the electrical field applied for anodic passivation, the mobility of OH⁻ originating from Ni (and Fe) hydroxides is reduced. These OH⁻ are retained at the interface between electrolyte and oxide layer instead of being released into electrolyte congruently with the cations. These OH⁻ participate then in forming Cr hydroxide, thus reducing the OD uptake from the electrolyte. As a result, the proportion of non-deuterated Cr hydroxide increases in the passive oxide film.

To support this hypothesis and to estimate whether the OH⁻ initially present in the native oxide film were enough for forming all newly grown Cr hydroxide, a balance
calculation was performed for $OH^{-}$. For simplification, species were considered in the form of $Cr(OH)_{3}$, $Ni(OH)_{2}$, $Fe(OH)_{2}$, and $Fe(OH)_{3}$. It was also assumed that all $OH^{-}$ in dissolved species can be utilized. Due to the difficulty in identifying Fe oxide and hydroxide, the estimation of $OH^{-}$ quantity was given for two extreme cases: all Fe species were hydroxides, and all Fe species were oxides. Co species were neglected due to their low content. Mo species were not considered since Mo quantity increased and its chemical nature was still uncertain. The compositional difference between the native and passive oxide films gives that the maximum $OH^{-}$ quantity that could be given by dissolved $Ni(OH)_{2}$, $Fe(OH)_{2}$, and $Fe(OH)_{3}$ were $3.5 \times 10^{-9} mol \cdot cm^{-2}$, $0.3 \times 10^{-9} mol \cdot cm^{-2}$, and $0.3 \times 10^{-9} mol \cdot cm^{-2}$, respectively. In the case of all Fe species being hydroxides, the formation of $Cr(OH)_{3}$ could thus re-use $4.1 \times 10^{-9} mol \cdot cm^{-2} OH^{-}$ at the most. However, the $OH^{-}$ contained in newly formed $Cr(OH)_{3}$ was $5.2 \times 10^{-9} mol \cdot cm^{-2}$. Therefore, at least $1.1 \times 10^{-9} mol \cdot cm^{-2} OD^{-}$ should still be taken from the electrolyte, amounting to 14 % of the total hydroxyl quantity in Cr hydroxide. If all Fe species were oxides and could not provide $OH^{-}$, the $OD^{-}$ uptake from the electrolyte among all hydroxyl in Cr hydroxide would be 22 %.

All the observations show that, for the $Cr_{15}Fe_{10}Co_{5}Ni_{60}Mo_{10}$ MPEA, anodic passivation in acid solution enhances the enrichment of Cr and Mo of the surface oxide film by preferentially dissolving Ni and Fe initially present in the native oxide film. They also indicate that, in this process, the hydroxyls initially contained in Ni (and Fe) hydroxides are retained by the electric field instead of being released congruently into the electrolyte. These retained $OH^{-}$ participate into the formation of Cr hydroxide, with
a small additional part complemented by $\text{OD}^{-}$ from the electrolyte. In the case of 1 h anodic passivation, Cr hydroxide formation uses mainly $\text{OH}^{-}$ originating from the dissolved hydroxides, and $\text{OD}^{-}$ surface uptake from the electrolyte for forming Cr hydroxide is only of minor extent. This hypothesis can be confirmed by the thermodynamic constant for chromium hydroxide precipitation ($\text{Cr}^{3+} + 3\text{OH}^{-}\rightarrow \text{Cr(OH)}_{3}$, $\text{Ks}(\text{Cr(OH)}_{3}=10^{-30})$ indicating that the formation of $\text{Cr(OH)}_{3}$ via precipitation of dissolved $\text{Cr}^{3+}$ species is not favoured. This observation corroborates the additional formation of $\text{Cr(OH)}_{3}$ in the outer layer by a conversion reaction in the solid state, retaining only the $\text{OH}^{-}$ initially present in the hydroxide layer and releasing $\text{Ni}^{2+}$ and $\text{Fe}^{2+}/3^{+}$. From these observations, a mechanism of formation of the outer oxide layer during anodic passivation is presented on Figure 6.

<figure><img src="image_12.png" /><figcaption>Figure 6: Mechanism of the outer oxide layer formation during anodic passivation in $\text{H}_{2}\text{SO}_{4}$ electrolyte prepared with $\text{D}_{2}\text{O}$.</figcaption></figure>

Previous work has shown that, when passivated, $\text{Cr}_{15}\text{Fe}_{10}\text{Co}_{5}\text{Ni}_{60}\text{Mo}_{10}$ MPEA can sustain at least 4.7 M $\text{Cl}^{-}$ without being corroded, whereas localized corrosion was already observed on SS 316L or Cantor alloys at 1 M or 0.3 M $\text{Cl}^{-}$, respectively [34].
The restrained hydroxyl uptake from the electrolyte may reflect a restrained anion uptake, which may contribute to the high corrosion resistance of the Cr<sub>15</sub>Fe<sub>10</sub>Co<sub>5</sub>Ni<sub>60</sub>Mo<sub>10</sub> alloy in chloride-containing medium [37]. In this previous work, it was clearly demonstrated that, for the same alloy, there is a limitation in the penetration of chloride anions into the outer layer during anodic passivation. In view of the results reported here, it seems established that there is competition between the formation of the new outer layer and the penetration of chlorides. Obviously, the action of the new outer layer alone is probably not sufficient to completely block the inward transport of anions towards the inner layer. Further work could investigate to which extent the penetration of chloride anions into the oxide film can be hindered by the abundance of reused hydroxyls.

Besides, ToF-SIMS depth profiling, both with or without isotopic labelling, evidences that there was almost no hydrogen in the inner layer of the surface oxide film. Hence, inwards OH<sup>-</sup> diffusion until the alloy/oxide interface, which has been observed for oxide film growth in high temperature high pressure water[43,46], is excluded from being primary during the oxide film growth at room temperature (and standard atmosphere).

## 4 Conclusion

The hydroxyl transport mechanisms between electrolyte and outer part of the nanometer-thick passive oxide film on Cr<sub>15</sub>Fe<sub>10</sub>Co<sub>5</sub>Ni<sub>60</sub>Mo<sub>10</sub> MPEA surfaces in 0.05 M H<sub>2</sub>SO<sub>4</sub> aqueous electrolyte were studied at room temperature and ambient pressure
using deuterium-labelled ToF-SIMS depth profiling, high-resolution XPS and ICP-OES.

The oxide film formed after immersion in pure water or sulfuric acid, or after anodic passivation in sulfuric acid retains the bilayer structure of the native oxide film. Thickness and composition of the oxide films show that although dissolution occurs during immersion in sulfuric acid, the oxide film is thickened with Cr and Mo enrichment during anodic passivation. Deuteration of the electrolyte has no effect on the passivation-induced alterations of the stratified structure and cationic composition of the nanometer-thick surface oxide film. A data treatment methodology for extracting the representative ToF-SIMS signals for deuterated species was developed, enabling the depth profile reconstruction for D-labelled species including deuterium atoms, deuterated hydroxyls, and deuterated Cr, Ni, Mo hydroxides. Surfaces immersed in pure or acidified D₂O exhibit deuterium uptake due to chemical interaction and isotopic exchange. The maximum intensities of the ToF-SIMS profiles for both deuterated and non-deuterated Cr hydroxides vary consistently with the content of Cr hydroxide in the outer layer measured by XPS. However, when being anodically passivated, the formation of non-deuterated Cr hydroxide is favored. Non-negligible extent of dissolution measured by ICP-OES and charge balance calculation suggested that the promoted formation of non-deuterated Cr hydroxide is related to the reutilization of hydroxyls initially in the native oxide film. It is thus suggested that during anodic passivation, the hydroxyls from the dissolving Ni (and Fe) hydroxides are retained by the electric field and form new Cr hydroxide, thus reducing the OD surface uptake from
the electrolyte.

### CRediT authorship contribution statement

**Xueying Wang:** Investigation, Validation, Methodology, Visualization, Writing - Original Draft. **Dimitri Mercier:** Conceptualization, Validation, Methodology, Supervision, Writing - review & editing. **Sandrine Zanna:** Investigation, Data Curation. **Antoine Syeux:** Investigation, Data Curation. **Loïc Perrière:** Resources.

**Mathilde Laurent-Brocq:** Resources. **Ivan Guillot:** Resources. **Vincent Maurice:** Funding acquisition, Writing - review & editing. **Philippe Marcus:** Funding acquisition, Project management, Conceptualization, Supervision, Writing - review & editing.

### Competing Interests

The authors 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

This work was supported by the European Research Council (ERC) under the European Union's Horizon 2020 research and innovation programme (ERC Advanced Grant agreement No. 741123).
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