# Deuterium retention in tungsten oxide: the role of oxide damage

Mykola Ialovega, Celine Martin, Cédric Pardanaud, Elodie Bernard, Martiane Cabié, Andrea Campos, Christian Grisolia, Marco Minissale, Thomas Neisius, Thierry Angot, et al.

► To cite this version:

Mykola Ialovega, Celine Martin, Cédric Pardanaud, Elodie Bernard, Martiane Cabié, et al. Deuterium retention in tungsten oxide: the role of oxide damage. Nuclear Fusion, 2025, 65 (7), pp.076024. 10.1088/1741-4326/addfeea. hal-05504924

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# Deuterium retention in tungsten oxide: the role of oxide damage

To cite this article: Mykola Ialovega et al 2025 *Nucl. Fusion* **65** 076024

View the [article online](#) for updates and enhancements.

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# Deuterium retention in tungsten oxide: the role of oxide damage

Mykola Ialovega<fnref n="1" />,<fnref n="2" />,<fnref n="3" /> ID, Celine Martin<fnref n="2" /> ID, Cédric Pardanaud<fnref n="2" /> ID, Elodie Bernard<fnref n="1" /> ID, Martiane Cabié<fnref n="4" />, Andrea Campos<fnref n="4" />, Christian Grisolia<fnref n="1" />, Marco Minissale<fnref n="2" /> ID, Thomas Neisius<fnref n="4" />, Thierry Angot<fnref n="2" /> and Régis Bisson<fnref n="2" />,\* ID

E-mail: regis.bisson@univ-amu.fr

Received 22 December 2024, revised 1 April 2025

Accepted for publication 3 June 2025

Published 13 June 2025

<figure><img src="image_2.png" /></figure>

CrossMark

## Abstract

Tungsten (W) is a common plasma-facing material in nuclear fusion devices. It readily oxidizes in the presence of oxygen, forming tungsten oxides, particularly $WO_3$, which may modify deuterium retention. This experimental study investigates mechanisms of the evolution of deuterium retention in a thermally stable ~100 nm $WO_3$ layer grown by thermal oxidation of a W substrate at 1073 K under low pure oxygen pressure of 7 Pa. 500 eV $D_2^+$ deuterium implantation and Temperature Programmed Desorption (TPD) were used to explore deuterium trapping and release from the oxide as a function of incident ion fluence and storage time in ultra high vacuum. Upon deuterium irradiation, the formation of a W-rich layer on the surface of the oxide is evidenced with x-ray photoelectron spectroscopy. After successive implantation/TPD cycles corresponding to an accumulated deuterium fluence >$10^{21}$ D m$^{-2}$, the appearance of an amorphous oxide encapsulated between two W-rich layers is observed with transmission electron microscopy. Following this high fluence deuterium implantation, deuterium retention increases by a factor of 10. The findings provide insight into the behavior of tungsten oxides under deuterium implantation, emphasizing the importance of considering its thermal stability and its structural modifications at the surface of tungsten plasma facing components.

Keywords: tungsten, oxide, retention, deuterium, TEM

(Some figures may appear in colour only in the online journal)

## 1. Introduction

Tungsten (W) is the preferred plasma-facing material (PFM) in fusion devices due to its exceptional properties such as

the high melting point, excellent thermal conductivity, resistance to sputtering and erosion, good mechanical properties and low retention of hydrogen isotopes (deuterium (D) and tritium) as compared with previously used carbon PFM [1]. However, a strong chemical affinity of tungsten with oxygen (O) raises concerns on the change of the PFM properties since oxide layers on W surfaces may be formed due to the presence of oxygen and water impurities in the vacuum vessel, the introduction of $O_2$ during tritium removal operations [2], or in case of an accident such as a loss of coolant or emergency shutdown [3, 4]. The presence of an oxide layer on tungsten surface changes the erosion yield [5, 6],

<figure><img src="image_3.png" /></figure>

Original Content from this work may be used under the terms of the Creative Commons Attribution 4.0 licence. Any further distribution of this work must maintain attribution to the author(s) and the title of the work, journal citation and DOI.

<footnote n="1">¹ CEA, IRFM, F-13108 Saint Paul-lez-Durance, France</footnote>

<footnote n="2">² Aix Marseille Univ., CNRS, PIIM, UMR 7345, 13013 Marseille, France</footnote>

<footnote n="3">³ University of Wisconsin-Madison, Department of Engineering Physics, 3DPSI, 1500 Engineering Drive,53706 Madison, WI, United States of America</footnote>

<footnote n="4">⁴ Aix Marseille Univ., CNRS, Centrale Marseille, FSCM (FR1739), CP2M, 13397 Marseille, France</footnote>

<footnote>* Author to whom any correspondence should be addressed.</footnote>
the reflectance [7] and the retention of D [8–10] within the PFM.

In addition, the effect of the high surface temperature on oxidation of W PFM during plasma operations should be considered in the next generation fusion devices like ITER, DEMO, SPARC etc. For example, the surface temperature at the ITER divertor strike-points will exceed 1000 K during the discharges due to the high heat loading up to 20 MW m⁻² and then might drop down to 340 K between the discharges due to the active cooling of the W PFM [1, 11, 12]. Such thermal cycles in presence of oxygen impurities, together with erosion and re-deposition effects, hlcan lead to the deposition of a complex multilayer structure composed of layers of tungsten and tungsten oxide, as observed in the WEST tokamak on actively cooled ITER-like plasma facing units [13] and on iner- tially cooled, W-coated divertor [14].

Temperature and oxygen partial pressure play an important role in the oxidation kinetics which determines the film thickness grown over a period of time [15]. Linear oxidation rate is usually obtained for surface temperature above 1073 K [16, 17] however sometimes parabolic oxidation rate is observed up to 1273 K [18]. When the surface temperature of W exceeds 1473 K, oxide starts to evaporate as quickly as it is formed [19]. Among the large number of tungsten oxides, WO₂, WO₂.₇₂, WO₂.₉₀, WO₂.₉₆ and WO₃ are the stable stoichiometric solid phases commonly observed in experimental works [17, 20, 21]. At the high temperatures relevant to ITER tungsten divertor, oxide usually consists of a mixture of several phases, mainly WO₂ and WO₃ [22, 23]. However, if the oxygen pressure is sufficiently high, oxidation of W will result in formation of the most stable WO₃ phase which has the lowest Gibbs energy in the range 300–1600 K. Thus WO₃ phase with a simple cubic perovskite-like structure is the most commonly investigated [24–26]. In this structure, the central site is not occupied and the corners share regular octahedra, with tungsten atoms located in the centers of each octahedron while oxygen atoms are located in the corners. Various allotropic forms of this ideal structure have been identified, and there are several WO₃ crystallographic structures that are stable within defined temperature ranges. For example, monoclinic and orthorhombic structures are stable from ambient temperature to 573 K and from 573 K to 973 K, respectively. Since the WO₃ crystal structure remains stable at high temperatures, one can expect that this oxide, which may form on the W divertor surface at elevated temperatures, will remain stable when the PFM cools down.

Unfortunately, there is still a lack of comprehensive data regarding W oxidation under low oxygen partial pressure conditions and the effect of these oxides on D retention. Recent experiments have demonstrated a decrease of D retention with an increase of oxygen impurity concentration on the clean W surface associated with the reduction of the available surface sites for the D adsorption [10]. However, these experiments also revealed a significant increase of fusion fuel retention in W when the thickness of the oxide exceeds one atomic layer. This result is consistent with other experimental [27–32] and modeling [33, 34] works which highlighted that native oxide

film, as well as thicker WO₃ oxide, act as a trapping layer for D, remarkably increasing retention as compared to clean W. Moreover, the synergistic effects of ion irradiation and temperature cause the structural changes of tungsten oxides even at the low kinetic energy of the incident D ions (≤50 eV) [6, 35, 36]. Finally, the enhanced retention of D in tungsten oxides, in additon to D trapping in defects, may also be associated with chemical bonding of D atoms to O atoms and the formation of tungsten bronze [6, 15, 37].

In the present paper, deuterium retention properties in ~100 nm-thick thermally stable tungsten oxide layer grown at low pure oxygen pressure (7 Pa) and high W temperature (1073 K) are investigated using laboratory experiments involving deuterium ion implantation and Temperature Programmed Desorption spectrometry (TPD). The analysis of desorbed deuterated species is coupled with characterizations of the surface and near-surface structure, such as chemical composition and morphology of the oxide. Deuterium retention in the WO₃ oxide is evaluated as a function of the ion fluence and storage time in vacuum, then directly compared with the results obtained for the case of the polycrystalline tungsten (PCW) with a native oxide layer. Finally, the synergistic effects of thermal cycling and ion irradiation on D retention in the oxide are discussed.

## 2. Experiment

### 2.1. Growth of a thermally stable tungsten oxide

A square-shaped PCW sample was provided by Toho Kinzoku Co. Ltd with a purity of 99.9% and dimensions of 7 × 7 × 1 mm³. The sample was mechanically polished and annealed at 1473 K for 2 h in vacuum by the company and then delivered to the PIIM laboratory in Marseille (France). The surface morphology of the pristine PCW sample was examined using confocal laser scanning microscopy (CLSM) and scanning electron microscopy (SEM). The surface roughness of the sample measured by CLSM was found to be $R_a \sim 90$ nm prior to oxidation.

The pristine sample was introduced into the high vacuum setup CAsPEr for high-temperature oxidation. Firstly, the sample was outgassed at room temperature (RT) at a base pressure of 10⁻⁶ Pa for 24 h. Secondly, the sample was heated with the linear heating ramp of 1 K s⁻¹ up to 1073 K and maintained for 1 h at this temperature. Then, the sample was exposed to a pure oxygen atmosphere at the initial pressure of 7 Pa and temperature of 1073 K. These conditions allow growing a thermally stable WO₃ crystalline layer with controlled thickness, enabling fundamental studies on D retention using TPD. The chosen high surface temperature exceeds the maximum temperature during thermodesorption, ensuring that TPD did not induce vacuum annealing of the sample. Additionally, such high surface temperatures are expected in the divertor region of a tokamak [1]. The low oxygen pressure (7 Pa) was selected to depart from usual high pressure studies (10⁴ to 10⁵ Pa [38]) and to try to approach more relevant tokamak conditions where the partial pressure of oxygen should be well below 0.1 Pa. The
linear drop of oxygen pressure as a function of time was followed to estimate the oxide thickness growth [15]. Oxidation was terminated after 146 s by evacuating the residual oxygen from the system using a turbo pump and cooling down the oxidation chamber with active cooling. Once the oxidized sample cooled down to RT within 20 min, it was taken in atmospheric condition for the initial surface analysis using CLSM and SEM as well as for the structural analysis of the formed oxide using Raman microscopy technique. SEM and transmission electron microscopy (TEM) cross-sectional analysis on a thin lamellae cut from the sample using focused ion beam technique (FIB) allowed investigation of the thickness and structural properties of the oxide at sub-nanometer scale.

Additionally, another identical pure tungsten sample was oxidized at the same temperature but at a more standard [38] pure oxygen pressure of 700 Pa which led to a formation of a perfectly ordered $WO_3$ oxide on its surface. This sample was only used as a reference in non-destructive Raman analysis [22, 39].

## 2.2. Deuterium ion implantation and retention measurements by thermodesorption

After the initial characterization, the oxidized sample was introduced into the ultra-high-vacuum (UHV) setup Camiter (base pressure <2.3 × 10⁻⁷ Pa) for cycles of in situ low-energy $D_2^+$ implantation followed by thermodesorption experiments (implantation/TPD). Kinetic energy of the impinging molecular ions was set to 500 eV, which corresponds to a kinetic energy per deuteron of about 250 eV/D in order to be comparable with previous experiments conducted on PCW samples with their native oxide [29]. Deuterium implantation was realized with an Omicron ISE 10 sputter ion gun at RT (i.e. 310 ± 15 K). Typical flux of D ions measured by a Keithley picoammeter was about 2.8 × 10¹⁶ D m⁻²s⁻¹. The incident D ion fluence was obtained by multiplying the duration of the implantation to the monitored ion beam flux and the accuracy of the fluence measurement was estimated to be of the order of 10%.

TPD experiments were performed with a temperature ramp of 1 K s⁻¹ up to 720 ± 5 K in order to avoid any modification of the oxide due to the annealing effect of TPD.

Two D implantation/TPD experimental campaigns for low and high accumulated fluence were performed on the sample. Before each experimental campaign, four initial implantation/TPD cycles were realized at a D fluence of 2.8 × 10¹⁹ D m⁻² in the as-introduced oxide after which a reproducible desorption of D from the sample was obtained and the tungsten oxide was considered to be cleaned from air contaminant and stabilized.

The first D implantation/TPD campaign was realized with incident D fluences varied between 0.7 × 10¹⁹ and 3.0 × 10²⁰ D m⁻². The total accumulated fluence on the sample reached about 2.8 × 10²¹ D m⁻². Finally, the sample was implanted with a penultimate fluence of 1.8 × 10²¹ D m⁻² and the TPD showed a retention in line with the previous measurements (so called, low accumulated fluence (LAF) state).

The last implantation of the first campaign, which was a standard 2.8 × 10¹⁹ D m⁻², however exhibited a retention that was 10 fold higher than the previous implantation/TPD measurements realized with the same fluence. The sample seemed to have evolved because of the large fluence of the penultimate implantation/TPD. We removed the oxide sample from UHV to check with the reference oxide sample that the retention measured in the LAF-state was reproducible, which was indeed the case. Therefore, we confirmed that the oxide sample in the LAF-state has evolved into a high accumulated fluence (HAF) state with higher D retention, once exposed to a total fluence of 4.6 × 10²¹ D m⁻². The oxide sample in the HAF-state was then re-introduced in the UHV setup Camiter and the second campaign of implantation/TPD cycles was realized on the HAF state oxide for various incident D fluences until the accumulated fluence reached about 9.0 × 10²¹ D m⁻².

In all implantation/TPD experiments, the quadrupole mass spectrometer (QMS) was recording simultaneously desorption of HD, $D_2$, HDO and $D_2O$. The collection slit of the differentially-pumped QMS was located 2 mm above the sample allowing for a high sensitivity line-of-sight detection of desorbing products. The QMS calibration was performed for various $D_2$ gas leaks with the assumption that the angular distribution of desorbing species follows a cosine law [29]. The sensitivity of detection of HD, HDO and $D_2O$ signals was calculated by applying appropriate correction factors to the measured sensitivity for $D_2$ molecules according to the formula:

$$S_X = S_{D_2} \times \frac{\sigma_X}{\sigma_{D_2}} \sqrt{\frac{m_{D_2}}{m_X}} \quad (1)$$

where $S_{D_2}$ is the sensitivity coefficient for $D_2$ molecules, $\sigma_X$ and $\sigma_{D_2}$ are the electron-impact ionization cross sections for molecules X and $D_2$ respectively for the selected ionization energy [40–42]. Since D desorption as HDO was in the noise level for the first experimental campaign (the noise level at $m/z = 19$ is about 100 times higher than at $m/z = 20$), HDO contribution was not considered in the data analysis of the LAF state. The contribution of $D_2O$ signal was found to be about 10% for both lowest and highest implanted D fluences. We stress that the present TPD setup has a line-of-sight QMS which allows to quantify heavy water desorption in an absolute manner [43]. The time delay (storage time) between the end of D ion implantations and the subsequent TPD measurements was 1.9 ± 0.2 h and the sample was kept in UHV during this time. D retention in the sample was quantified by time integration of the collected desorption spectrum of each implantation/TPD experiment and the uncertainties were estimated as the standard deviation from replicate measurements. In addition, the effect of the storage time (or the isothermal desorption) at RT was investigated for two different incident fluences of 2.8 × 10¹⁹ D m⁻² and 8.8 × 10¹⁹ D m⁻².

After the implantation/TPD experiments, the oxidized PCW sample was analyzed *ex situ* with CLSM, SEM and Raman spectroscopy. Additionally, cross-sectional observations were performed using the TEM on the FIB-cut lamella. Since the sputtering coefficient of tungsten oxide for 500 eV
$D_2^+$ ions is about $4 \times 10^{-4}$ [44] and the implantation range of the ions at an angle of 45° is about 12 nm [45], it is expected that the near surface ~12 nm layer of the oxide might be modified with increasing deuterium irradiation fluence. This modification might be not detected by the Raman spectroscopy technique which probes the whole thickness of the oxide. That is why, the elemental composition modification in the oxidized sample with and without implanted D was performed using x-ray photoelectron spectroscopy (XPS) [46]. XPS measurements were realized between the two experimental campaigns, that is when the oxide sample just switched from the LAF-state to the HAF-state.

### 3. Locally disordered crystalline $WO_3$ oxide

Figure 1(a) shows a CLSM image of the surface of the $WO_3$ sample grown at 7 Pa with four oxide grains, labeled G1, G2, G3, and G4, which are representative of the sample. The coloration of the oxide grains results from interference phenomena, which are determined by the thickness of the layer and its refractive index [15]. Consequently, the color also depends on the oxygen content, the phase of the oxide, and the orientation of light propagation, as $WO_3$ is an anisotropic material. Figure 1(b) presents typical SEM micrographs of the surfaces of grains G1–G4, showing oxide nanorods structures. The shape and size of these structures are likely determined by the variations in oxidation kinetics across different grain orientations of the W substrate. CLSM measurements of the roughness, averaged across the surface of the formed oxide layer, yield a value of approximately $R_a \sim 80$ nm, consistent with the initial roughness of the W substrate. The black curve in figure 1(c) depicts the Raman spectrum of the reference oxide grown at 700 Pa, with two distinct peaks at $703 \text{ cm}^{-1}$ and $805 \text{ cm}^{-1}$, characteristic of the monoclinic structure of an ideal crystalline $WO_3$ oxide. The colored spectra in figure 1(c) indicate a similar $WO_3$ stoichiometry among grains G1–G4; however, the broadening of the peaks in the $500–1050 \text{ cm}^{-1}$ domain suggests local disorder in the oxide compared to the ideal structure of a perfect $WO_3$ oxide [22, 37, 39], which should be the results of the lower oxygen pressure growth. Figure 1(d) shows that after all implantation/TPD cycles realized on the sample, no clear evolution of the Raman spectra is noticed in the $500–1050 \text{ cm}^{-1}$ domain. While Raman microscopy is not sensitive to possible modifications of the subsurface of the oxide generated by D implantation, this measurement shows that annealing up to 720 K did not cause structural changes through the whole thickness the oxide grains. Interpretation of small modifications of the Raman spectra in the $100–500 \text{ cm}$ domain is not under the scope of this work as this would require to significantly increase D implantation volume in the oxide for a meaningful conclusion.

Cross-sectional TEM observations, shown in figure 1(e), allowed for the measurement of the pristine oxide thickness, which varies between 45 and 160 nm, highlighting the non-uniform oxidation kinetics even within a single grain. As seen in the inset image in figure 1(e), presenting a

high-magnification cross-sectional view of the oxide, the $WO_3$ layer exhibits a crystalline structure in the bulk. The contrast difference in the TEM image is due to the curtaining effect related to the thin lamella specimen preparation with FIB.

In summary, we conclude that the $WO_3$ layer grown on the PCW surface at 1073 K and a pure oxygen pressure of 7 Pa has a non-uniform thickness ranging from 45 to 160 nm composed of a nanorods monoclinic crystal oxide structure with some local disorder in the bulk. We note that despite the high specific area of the oxide layer, we did not observe any $D_2$ dissociative chemisorption onto the sample, similarly to the native oxide behavior [10].

### 4. Deuterium retention in $WO_3$: LAF state

#### 4.1. Deuterium desorption from the as-introduced LAF oxide

Figure 2(a) shows the TPD spectra for D desorption as HD and $D_2$ during the initial four implantation/TPD cycles for an incident fluence of $2.8 \times 10^{19} \text{ D m}^{-2}$. In the first experimental cycle, D release from the as-introduced $WO_3$ oxide occurs in the temperature range between 350 K and 680 K, consisting of two desorption peaks at approximately 450 K and 600 K with similar desorption rate maxima of about $3 \times 10^{15} \text{ D m}^{-2}\text{s}^{-1}$. In the subsequent three experimental cycles, the low-temperature desorption peak at 450 K remains unchanged, while the high-temperature peak shifts to 560 K and shows a gradual increase in desorption rate, reaching a maximum of about $2.4 \times 10^{16} \text{ D m}^{-2}\text{s}^{-1}$. As a result of this evolution in the TPD spectra, D release as HD and $D_2$ increases from $6 \times 10^{17} \text{ D m}^{-2}$ in the first cycle, saturates at approximately $2.3 \times 10^{18} \text{ D m}^{-2}$ by the fourth cycle and remains consistent for all subsequent eleven implantation/TPD cycles, each performed with the same incident D fluence of $2.8 \times 10^{19} \text{ D m}^{-2}$, as shown in figure 2(b). The observed changes in the D thermodesorption spectra from the sample indicate that the surface and bulk conditions of $WO_3$ were unstable and involved modifications of the available trapping sites and de-trapping energy for D atoms.

In previous experiments on PCW samples [29] conducted with the same setup, cycles of D ion implantation/TPD experiments were performed after a couple annealing cycles at 1200–1350 K for 10 min. Annealing facilitates the desorption of loosely bound impurities, such as water molecules and hydrocarbons, from the surface and some impurities trapped in the bulk of PCW samples. However, when implantation/TPD experiments are conducted on an ‘as-introduced’ PCW sample without annealing at 1200–1350 K, D release might occur differently due to the presence of surface impurities. Additionally, the near-surface layer of the oxide might be modified due to the physical sputtering of O by the impinging $500 \text{ eV } D_2^+$ ions, as well as chemical sputtering. Consequently, the sub-surface oxide layer becomes more resistant to compositional changes with each subsequent implantation/TPD cycle due to the increasing fraction of W atoms in the layer. Once all water molecules and hydrocarbons are removed from the surface of the oxide, and the sub-surface region of the
<figure><img src="image_4.png" /><figcaption>**Figure 1.** (a) Confocal laser scanning microscopy image, (b) SEM micrographs of four typical oxide grains G1–G4 on the surface of the PCW sample oxidized at 1073 K. (c) and (d) Raman spectra for the grains G1–G4 of the as-oxidized sample and after D implantation/TPD annealing cycles at 720 K. (e) Oxide thickness measurement with high-angle annular dark-field TEM (HAADF-TEM) on the as-oxidized sample. The inset shows a closer TEM view of the oxide.</figcaption></figure>

oxide becomes resistant to irradiation with 500 eV $D_2^+$ ions, the retention mechanisms stabilize, and the oxide can be considered clean and stable. These factors contribute to the reproducibility of our implantation/TPD measurement cycles in the examined fluence range of (0.7 – 178.4) × 10¹⁹ D m⁻². Henceforth, all D desorption spectra will be presented as average spectra of replicate measurements (for example, see the squared green points for the case of 2.8 × 10¹⁹ D m⁻² in figure 2(a).

## 4.2. Deuterium retention from the stable LAF oxide

### 4.2.1. Deuterium desorption as a function of incident fluence.
Figure 3(a) shows the obtained TPD spectra for D atoms retained in the $WO_3$ as a function of deuterium fluence. Generally, the TPD spectra exhibit two desorption peaks whose intensities depend on the incident D fluence. The high-temperature peak is dominant for fluences below 15 × 10¹⁹ D m⁻², after which the lower-temperature peak becomes
<figure><img src="image_5.png" /><figcaption>**Figure 2.** TPD measurements for the WO₃ sample sequentially irradiated with 500 eV D₂⁺ ions with a fluence of 2.8 × 10¹⁹ D m⁻² at RT and heated with TPD up to 720 K (implantation/TPD cycles) for (a) the LAF and (c) the HAF state of the oxide. D retention for the four first implantation/TPD cycles in (b) the LAF state and (d) the HAF state of the oxide. The TPD with square symbols represent the average spectra obtained over replicate implantation/TPD experiments (12 replicates in the LAF-state, 7 replicates in the HAF-state) with a storage time of 1.9 ± 0.2 h. Circled numbers denote the cycle number of the deuterium implantation/TPD experiments.</figcaption></figure>

dominant. This shift in D desorption behavior results in a gradual decrease in the temperature position of the TPD spectrum desorption rate maxima as the incident fluence increases, up to 178.4 × 10¹⁹ D m⁻² (see figure 3(b)).

As shown in figure 3(c), the absolute contributions of HD and D₂ to the total quantity of D released from WO₃ increase similarly with fluence. This results in a consistent proportion of HD and D₂ signals contributing to the TPD spectra across all fluences: 70%–80% of D is desorbed as D₂ and 15%–25% as HD. As previously mentioned, about 10% of deuterium is desorbed as D₂O molecules 2 h after implantation, while the HDO signal remains at the noise level of m/z = 19. It is important to note that the large error bars at the lowest fluence of 0.7 × 10¹⁹ D m⁻² as well as for the data points representing the D₂O signal are due to the difficulty of background subtraction from the experimental TPD spectra.

Figure 4 shows the desorbed amount of D from WO₃ as a function of impinging D ion fluence, with the black dashed line presenting the results obtained for similar in-situ retention

measurements on pristine PCW samples with their native oxide [29]. Pristine PCW samples exhibit a simple relationship between the implantation fluence and the TPD desorption retention<sub>PCW</sub> ~ fluence<sup>0.65±0.03</sup>, typical of a D retention influenced by both diffusion into the bulk and trapping at defects. However, D release from the thick WO₃ layer can be described by two power law functions presented in equation (2):

$$ 
\begin{array}{l}
\text{retention}_{\text{WO}_3\text{LAF}} \propto \text{fluence}^{1.28} \\
\text{for fluence} < 15 \times 10^{19} \text{D m}^{-2}, \\
\text{and} \\
\text{retention}_{\text{WO}_3\text{LAF}} \propto \text{fluence}^{0.53} \\
\text{for fluence} \ge 15 \times 10^{19} \text{D m}^{-2}.
\end{array} 
\quad (2) $$

Thus, D release behavior in the WO₃ oxide is different from PCW for low incident fluences, pointing to a dominant D trapping on defects with potentially an increasing density of trap with increasing fluence. However, for high incident fluences, D release from the oxide follows almost a square
<figure><img src="image_6.png" /><figcaption>**Figure 3.** TPD measurements for the WO₃ sample irradiated with 500 eV D₂⁺ ions at RT at different incident fluences with a storage time of 1.9 ± 0.2 h in the case of (a) the LAF state and (e) the HAF state of the oxide (1 ML = 10¹⁹ D m⁻²). (b) (resp. (f)) Evolution of the temperature at which the desorption rate is maximum. (c) (resp. (g)) Absolute release of D in molecular species of HD, D₂ and D₂O. (d) (resp. (h)) Release of D in HD, D₂ and D₂O in fractions.</figcaption></figure>

root proportionality as for the case of PCW: *retention*<sub>PCW</sub> ∝ fluence<sup>0.65</sup> [29], showing the increased influence of diffusion into the bulk on D retention.

The solid lines in figure 4 represent the theoretical retention maxima calculated with SRIM [45] for our implantation conditions for the pure tungsten (black line) and for the tungsten
<figure><img src="image_7.png" /><figcaption>Figure 4. D retention in the WO₃ sample (colored squares) can be represented by two desorption regimes (squared numbers) in the LAF state: 1) retention ∝ fluence¹.²⁸ and 2) retention ∝ fluence⁰.⁵³. D retention in the oxide in the HAF state is represented by the regime 3) retention ∝ fluence⁰.⁸¹. D retention in pristine PCW is represented by the dashed black line retentionPCW ∝ fluence⁰.⁶⁵±⁰.⁰³ [29]. The solid lines represent the theoretical retention maxima calculated with the SRIM code for the WO₃ oxide (red line) and for pure tungsten (black line).</figcaption></figure>

trioxide (red line) materials. The SRIM prediction of deuterium ion implantation probability is quite similar for pure W and for WO₃. However, our experimental results on the WO₃ oxide in the LAF state exhibits a fluence dependence markedly different than PCW sample with a native oxide layer. In particular, this fluence dependence has a kink when the ion fluence is above 10²⁰ D m⁻². It is important to note that an implantation duration in the 2–20 h range is needed for deuterium fluence in the 10²⁰ – 10²¹ D m⁻² range. On this time scale, PCW samples with a native oxide are known to exhibit dynamic retention [29] i.e. deuterium desorbs naturally at RT. Thus, in order to check if a difference in dynamic retention could explain a difference in fluence dependence between a pristine PCW and a WO₃ sample, an experimental study on the dynamic deuterium retention of WO₃ at RT is presented in the next section.

### 4.2.2. Deuterium retention as a function of storage time.

Figures 5(a) and (b) present the evolution of the TPD spectra from the oxide sample in the LAF state as a function of storage time for two incident fluences of 2.8 × 10¹⁹ D m⁻² and 8.8 × 10¹⁹ D m⁻². For both D ion fluence, it is observed that the measured D desorption rate strongly decreases as the storage time increases from 2 h to more than 130 h, evidencing dynamic retention in WO₃.

It is surprising that such dynamic retention is linked with a strong decrease of the TPD peak at high-temperature (550 K)

because usually such a high-temperature desorption peak is attributed to traps in the material with a high detrapping energy for deuterium $E_d > 1.4$ eV [34, 47] which should limit D release from W at RT. The dynamic retention results summarized in figure 6 highlight that D release from the oxide occurs at RT for at least 330 h. Storing the WO₃ sample in UHV during this time, about 90% of the initial D retention was released from the oxide at RT for both incident fluences. The loss of implanted D from the oxide occurs in a non-linear fashion and can be approximated with an exponential decay of the form:

$$retention(t) = B + S \times \exp\left(-\frac{t}{\tau}\right), \quad (3)$$

where $(B+S)$ is the retention just after implantation, $t$ is the storage time and $\tau = 97 \pm 8$ h in the LAF state. Experiments by Bisson et al [29] performed in the same conditions on clean PCW samples showed that the storage time dependency in that case is well adjusted with a significantly shorter exponential decay $\tau \approx 19$ h and only 50% of PCW’s initial retention was released at RT.

In summary, the WO₃ oxide in the LAF state exhibited a surprising dynamic retention as compared to PCW samples: WO₃ in the LAF state lose twice more its initial retention at RT than PCW, but at a slower rate, while its temperature of maximum desorption rate is 100 K higher than the one of PCW. These features appear contradictory in the frame of
<figure><img src="image_8.png" /><figcaption>**Figure 5.** Storage time dependency of the deuterium TPD spectra for WO₃ sample irradiated with 500 eV D₂⁺ ions at RT: (a) 2.8 × 10¹⁹ D m⁻² and (b) 8.8 × 10¹⁹ D m⁻² for the LAF state. Respectively, (c) and (d) for the HAF state.</figcaption></figure>

usual 1-dimensional rate equation interpretations. Such surprising behavior may find an explanation by taking into consideration the complex WO₃ microstructure seen in figure 1, in particular the interfaces between the nanorods.

**5. Oxide transition from LAF state to HAF state**

As described in the experimental section 2.2, D retention properties of the oxidized tungsten sample WO₃ evolved after the completion of the first D implantation/TPD campaign, when the accumulated fluence reached 4.6 × 10²¹ D m⁻², the oxide transitioned from the LAF to HAF state which has different D retention properties.

The red curve ① in figure 7(a) shows the XPS spectrum of the WO₃ sample that transitioned in the HAF state. The sample was transferred from the implantation/TPD setup to the XPS setup through air. The W4f region (41–28 eV) of the WO₃ sample in the HAF state exposed to air (for about a month) resembles the spectrum typical for the tungsten oxides [48]. The intensity maximum of the spectrum is located at 35.7 eV signifying the dominant contribution of the WO₃ (W⁶⁺) with

the shoulder ending at 30 eV signifying the presence of small fractions of WO₂ (W⁴⁺) and metallic tungsten. Heating of the sample in UHV up to 680 K (a temperature slightly lower than the maximum temperature reached during TPD experiments) does not induce any noticeable evolution of the XPS spectrum (see spectrum ②). An evolution of the XPS spectrum occurs after the sample was irradiated with 500 eV D₂⁺ ions at RT (D fluence ~3 × 10¹⁹ D m⁻²). The increased contribution of the metallic tungsten in spectrum ③ suggests that the surface and sub-surface layer of the oxide changed due to the sputtering by the impinging D ions. It is important to note that the depth resolution of the XPS measurement is well below 10 nm while the implantation range of D ions is about 12 nm for WO₃ which is the expected depth of the damaged sub-surface layer. Annealing of the D-implanted oxide up to 680 K (spectrum ④) prompts a release of the trapped D (see figure 5). However, release of D as well as the subsequent cooling of the sample down to RT (spectrum ⑤) do not modify the composition of the XPS-probed layer highlighting its stability. Further irradiation of the oxide with 500 eV D₂⁺ ions at RT realized with incident fluence of ~1 × 10²² D m⁻² reveals a significant increase in
<figure><img src="image_9.png" /><figcaption>**Figure 6.** D retention in the WO₃ sample as a function of storage time in UHV at RT for two incident fluences of 2.8 × 10¹⁹ D m⁻² and 8.8 × 10¹⁹ D m⁻² in the LAF and HAF states.</figcaption></figure>

<figure><img src="image_10.png" /><figcaption>**Figure 7.** X-ray photoelectron spectroscopy in the W4f region (41–28 eV) (a) spectra and (b) evolution of oxide composition during six experimental steps.</figcaption></figure>

intensity of the metallic W peak highlighting that the surface and sub-surface layer loose oxygen to a greater extent (specrum ⑥).

XPS spectra can be fitted with six Gaussian peaks attrib-
uted to the contributions of W⁶⁺ (WO₃), W⁴⁺ (WO₂), and
pure tungsten metal W⁰ species, all having 5/2 and 7/2 com-
ponents [36, 48]:

(1) 38.0 eV: W⁶⁺ 5/2,

(2) 35.7 eV: W⁶⁺ 7/2 and W⁴⁺ 5/2,

(3) 33.4 eV: W⁴⁺ 7/2 and W⁰ metal 5/2,

(4) 31.4 eV: W⁰ metal 7/2.

Decomposition of the species in each peak is done by using
a branching ratio between 5/2 and 7/2 components of 0.75. It is
important to note that the peaks at 35.7 eV and 33.4 eV can be
attributed to both, a WO₂ pure oxide as well as to a defectuous
WO₃ oxide. Figure 7(b) presents the evolution of WO₃, WO₂
and W in each XPS experimental spectra ①–⑥ calculated from
the area of the corresponding peaks. The results show that the
surface and sub-surface of the as-introduced oxide are mainly composed of two contributions, WO₂ and dominating WO₃ which is coherent with the Raman measurements presented in figure 1(c) and discussed in section 3. Comparison of the elemental composition of the sample between the XPS experiments ① and ⑥ reveals that irradiation with D ions creates a sub-surface oxide layer with a dominant contribution of WO₂ mixed with metallic W. It is fair to suggest that irradiation of the as-oxidized sample with high fluence of D facilitates preferential physical sputtering of O as well as chemical erosion of O and causes the sub-surface layer stoichiometry to evolve from WO₃ to a significantly sub-stoichiometric damaged tungsten layer characteristic to the HAF state of the oxide.

## 6. Deuterium retention in WO₃: HAF state

The WO₃ sample in the HAF state was exposed to the ambient air at RT after XPS measurements for several months before it was re-introduced in the Camiter setup for the second experimental implantation/TPD campaign. The next sections describe the evolution of the retention properties of the sample and highlight the differences between the LAF and HAF oxide states.

### 6.1. Deuterium desorption from the as-introduced HAF oxide

D release from the HAF oxide after the air exposure occurs in the temperature range between 350 and 680 K and generally resembles the release from the as-introduced LAF unstable oxide. However, figure 2(c) shows that while the position of the lower temperature desorption peak remains at 450 K, the position of the higher temperature peak in the HAF state is already shifted closer to 560 K for the first three implantation/TPD cycles as compared to the LAF oxide state. Figure 2(d) shows that D release from the HAF oxide state gradually increases from $0.2 \times 10^{19} \text{ D m}^{-2}$ in the first cycle, saturates at $1.9 \times 10^{19} \text{ D m}^{-2}$ after the fourth cycle and remains consistent for all subsequent six implantation/TPD experiments realized at the same incident fluence of $2.8 \times 10^{19} \text{ D m}^{-2}$. The gradual increase in desorption rate of the second desorption peak of the HAF oxide saturates at approximately $1.3 \times 10^{17} \text{ D m}^{-2} \text{ s}^{-1}$ which is almost ten fold higher than compared to the LAF oxide state (figure 2(a)) and consistent with the last implantation of the first implantation/TPD experimental campaign, i.e. just after the transition from the LAF to HAF state and before air exposure. Similarly to the ‘as-introduced’ LAF oxide, the processes of desorption of loosely bound water molecules and hydrocarbons at the surface, as well as chemical and physical sputtering of O by the impinging D ions, seen by XPS in section 5, should explain the observed evolution of D release from the ‘as-introduced’ HAF oxide until the retention mechanisms stabilize.

### 6.2. Deuterium retention from the stable HAF oxide

#### 6.2.1. Deuterium desorption as a function of incident fluence.
The TPD spectra for D atoms retained in the WO₃ sample in

the HAF state exposed to different incident fluences exhibit two desorption peaks whose positions and desorption rate depends on the incident fluence (figure 3(e)). Similarly to the case of the LAF oxide, a shift in D desorption behavior occurs at high incident fluence when the low-temperature desorption peak becomes dominant over the high-temperature desorption peak resulting in a gradual decrease in the temperature position of the desorption rate maxima as the incident fluence increases (figure 3(f)). The absolute desorption rate maxima of each TPD spectra shown in figure 3(e) is significantly higher compared to the LAF case (figure 3(a)) signifying the increased D retention in the HAF oxide. This change occurs due to the increased absolute contributions of HD, D₂, HDO and D₂O from the HAF oxide shown in figure 3(g). In the same time, the relative contributions of the species evolved between the LAF and HAF oxide, with the HD product fraction being smaller and the HDO product fraction increasing on par with the D₂O product, as presented in figure 3(d) and in figure 3(h).

Contrary to the results obtained for the LAF oxide state, the total desorbed amount of D from the WO₃ in the HAF state as a function of incident fluence (figure 4) exhibits a behavior qualitatively similar to the pristine PCW samples and can be described by one power law function within the whole range of impinging fluence of the study:

$$retention_{WO_3 HAF} \propto fluence^{0.81}. \quad (4)$$

This highlights a significant effect of diffusion in the retention mechanisms of the WO₃ oxide exposed to $500 \text{ eV D}_2^+$, though the total amount of D released from the oxide in the HAF state is almost an order of magnitude higher and comes close to the implantation probability for the tungsten trioxide (red line in figure 4).

#### 6.2.2. Deuterium retention as a function of storage time.
If the effect of bulk diffusion in the HAF state is still present like in the LAF state, it appears in figures 5(c) and (d) that diffusion to the surface and desorption at RT is strongly reduced in the HAF state. Dynamic retention appears to be slowed down and D retention in the $10^{19} \text{ D m}^{-2}$ range is still observed even after 200 h of storage time. There are two important remarks to be highlighted from the experiments. First, similarly to the LAF case, the temperature position of the desorption rate maxima slightly shifts to about 550 K after couple dozen hours of storing the HAF oxide in UHV upon D implantation. Second, diversely to the LAF case, the desorption rate maxima of the spectra almost does not evolve in the case of implantation at $2.8 \times 10^{19} \text{ D m}^{-2}$ while it decreases by about 30% in the case of implantation at $8.8 \times 10^{19} \text{ D m}^{-2}$ after more than 137 h of storage in UHV condition. Figure 6 shows that the amount of D desorbed from the oxide in the HAF state almost did not change for the implantation realized at $2.8 \times 10^{19} \text{ D m}^{-2}$ during 200 h of storage time while a slow decrease in retention is evidenced during 137 h of storage time for the implantation realized at $8.8 \times 10^{19} \text{ D m}^{-2}$. These remarks could be indicative of the difference in isothermal release of D due to the higher degree of trap saturation in the case of the higher
impinging fluence on the HAF oxide. Indeed, recent stud-
ies [34, 49] showed that detrapping energy for D in a defect in
tungsten decreases with an increase of the number of trapped
D atoms. Thus, we can suggest that at the low incident fluence
the detrapping energy for D is high enough to prevent D release
at RT. In the case of exposures realized at the elevated fluence,
the same type of traps may contain larger amounts of D which
may be released from the sample even at RT. Finally, one
should notice that differences of dynamic retention between
the LAF and HAF oxide states are observed despite quite
similar position and shape of desorption peaks. Again, this is
going against the simplistic picture of dynamic retention being
related to low-binding energy within 1-dimensional models.
As we proposed at the end of section 4, an interfacial effect
could be again at play, as suggested by the structural observa-
tion presented in the following section.

7. Morphological evolution of tungsten oxides

7.1. *High-resolution observation of oxide evolution*

This systematic D retention study on WO₃ oxide layer grown
on PCW demonstrates several important results. First, Raman
(figures 1(c) and (d)) and XPS (figure 7) analysis revealed that
a WO₃ oxide is stable when heated below the temperature of
oxidation, in our case WO₃ oxides were grown at 1073 K and
were found to be thermally stable up to 750 K. Second, after
three initial D implantation/TPD experiments (for an accu-
mulated fluence of ~8 × 10¹⁹ D m⁻²), realized on the oxid-
ized PCW sample, D release behavior remains unchanged for
500 eV D₂⁺ ion implantation and subsequent thermodesorp-
tion experiments up to 720 K ('stable' oxide in the LAF state)
until irradiated with large incident fluence on the order of
2.8 × 10²¹ D m⁻² ('stable' oxide in the HAF state). Obtaining
such an oxide stability, as seen in TPD, should be related to
the removal of the surface impurities (e.g. water molecules)
and a modification of the surface and near-surface of the oxide
by preferential sputtering of O atoms, thus increasing resist-
ance of the implantation zone to sputtering with increasing
D fluence. The oxygen depletion sputtering mechanism from
D ions is corroborated by the XPS measurements (figure 7)
showing that this happens at RT. In order to better character-
ize this sputtering mechanism, expected to extend over a few
surface atomic planes (a couple of nm), we performed cross-
sectional measurements on the HAF WO₃ sample after the
second experimental campaign.

Figure 8(a) presents the cross-sectional TEM view of thin lamellae cut from a part of the WO₃ sample which was not exposed to the D ion beam. The micrographs were obtained with the energy filtered TEM (EF-TEM) which allows to acquire the oxygen and carbon maps of the area of interest. The surface of the oxide is protected by the coating mixture of platinum and carbon carefully deposited during the preparation of the lamellae which preserves the surface condition and prevents a possible damage by FIB cutting. It can be seen that the non D-exposed part of the HAF WO₃ is a >45 nm thick homogeneous crystalline oxide grown on W substrate. Figure 8(b)

presents the zero-loss and EF-TEM maps acquired on the HAF WO₃ sample in the D-ion exposed part. Such a cross-sectional view reveals that the bulk part of the oxide is similar to the original crystalline WO₃ structure observed in figure 8(a). However, the sub-surface of the oxide has drastically evolved and features two 2 nm-thick W-rich layers separated by an amorphous 4 nm-thick oxide layer. These striking modifications are likely formed by the combination of high D ion fluence and TPD thermal cycles and should be linked with the transition of the sample from LAF to HAF oxide state. A recent work by Zavasnik et al [50] indicated stronger binding of D within an amorphous oxide matrix compared to the crystalline WO₃.

It is clear that preferential sputtering of O by the impinging
500 eV D<sub>2</sub><sup>+</sup> ions can induce the formation of a 2 nm thick W-
rich layer on the oxide surface as evidenced by the XPS exper-
iments (figure 7). Assuming the sputtering yield of tungsten
oxide by 250 eV D to be around 0.002 O atom per ion [5], in
total 9.1 × 10<sup>18</sup> O m<sup>-2</sup> were sputtered during the first exper-
imental campaign and about 1.8 × 10<sup>19</sup> O m<sup>-2</sup> atoms were
sputtered during the second experimental campaign. These
values are too low to explain the formation of the 2 nm W-rich
layer, which suggests the contribution of chemical erosion.
The effect of formation of metallic tungsten layer on the sur-
face of tungsten trioxide was also recently reported for expos-
ures to D ions with lower kinetic energy of 50 eV and a 100
times higher fluence [6, 36].

However one needs to find an interpretation for the second
oxygen-depleted layer which is buried 6 nm below the O-
depleted layer at the surface.

7.2. *A model for the evolution of WO₃ oxide structure upon*
*deuterium irradiation and thermal cycling*

There are two possible explanations for the observed changes
of the sub-surface layer of the WO₃ oxide during our
experiments.

The first explanation is linked with the history of the D
irradiation/TPD experimental campaigns and the exposure to
ambient air between and after them (prior the FIB-TEM ana-
lysis). After the release of surface impurities and stabilization
of the oxide (three first implantation/TPD cycles with an accu-
mulated D ion fluence of ~8 × 10¹⁹ D m⁻²), D implantations
to a ~4.5 × 10²¹ D m⁻² fluence may have resulted in the cre-
ation of a thick (~10 nm) W-rich layer by preferential sput-
tering of O atoms from the oxide and the sample evolves to a
higher retention state. Then, the sample was exposed for one
month to ambient air and later introduced in an UHV setup
for XPS characterization and D implantation. During this one
month of air exposure, the 10 nm thick W-rich layer could
have partially re-oxidized, as suggested by XPS. Further D
irradiation up to ~10²² D m⁻² could lead to the recreation of
a W-rich layer due to the preferential sputtering of O from the
oxide. After the XPS experiment the sample was exposed to
ambient air for several month which again could have resulted
in the partial oxidation of the top W-rich layer and the forma-
tion of an amorphous native oxide. After that, the sample was
<figure><img src="image_11.png" /><figcaption>**Figure 8.** EF-TEM observations of thin lamellae cut from (a) non-irradiated part of the oxidized sample which only underwent multiple heating cycles up to 720 K, and, (b) the D-irradiated part of the oxidized sample.</figcaption></figure>

introduced again in the UHV setup for the second D implantation/TPD campaign. Multiple D ion exposures to a fluence of ~4.5 × 10²¹ D m⁻² could lead to the formation of a new W-rich layer on the surface of the amorphous native oxide evidenced by the TEM observations (figure 7). However, this version of the modification of the oxide does not account for the desorption of water molecules in every TPD experiment nor it explains why the absolute high D retention observed at the beginning of the second experimental campaign (HAF state) was identical to the absolute retention measured at the end of the first campaign (when the sample just transitioned to the LAF state). Finally, this explanation does not explain why the third air exposure (for more than a year) between the second D implantation/TPD campaign and the FIB-TEM measurements would not have re-oxidized the surface, since we do see O-depleted layers in TEM even at the surface.

The second explanation is pictured in figure 9. It considers that the first three implantation/TPD experiment realized with the incident fluence of 2.8 × 10¹⁹ D m⁻² induced release of the surface contamination and, due to the preferential sputtering of O from the oxide, the creation of the first thin W-rich layer. Further implantation/TPD experiments realized with moderate fluence (<15 × 10²⁰ D m⁻²) evidenced the stability of D

release from the sample which probably can be associated with D trapping mainly in the W-rich layer with numerous vacancy-type defects and some fraction of D atoms reaching below and trapping in the original WO₃ layer. This is the oxide in the LAF state. Irradiation of the sample with high D fluence (exceeding 15 × 10²⁰ D m⁻²) filled all available trapping sites for D in the W-rich layer, and a major fraction of implanted D atoms diffused and trapped in the bulk of WO₃ layer. Subsequent thermal desorption prompted D release from the W-rich layer as well as D and O desorption from WO₃, producing a defective oxide layer below the surface tungsten layer and the oxide evolved from the LAF to the HAF state characterized by a significant increase of D retention in the sample. In the HAF oxide state, O migrates from the bulk of the oxide during thermal desorption annealing of the sample, and fills some of the vacancies below the W-rich layer at the surface. This process generates the formation of a second W-rich layer within the maximum implantation depth of 500 eV D₂⁺ of 12 nm. Diffusion of O up to the surface and recombination with D surface atoms explains the detected desorption of water presented in figure 3.

Generally, we believe that the high retention of D in tungsten trioxide as compared to the retention levels in PCW can
<figure><img src="image_12.png" /><figcaption>**Figure 9.** Schematic representation of the evolution of the WO₃ oxide during D implantation/TPD experiments. (a) As-oxidized sample containing some water and hydrocarbons at the surface; (b) surface sputtering of the oxide irradiated at low accumulated fluence (LAF); (c) creation of an oxygen-enriched amorphous layer encapsulated between two oxygen-depleted W-rich layers when WO₃ is irradiated at high accumulated fluence (HAF), as seen by the TEM observations shown in (d).</figcaption></figure>

be explained by chemical bonding between O and D resulting in the formation of tungsten bronze which could be related to the reversible change of the oxide color, from purple to brown (not shown), in the implantation region of our sample.

To summarize, the observed TPD behavior of deuterium in the PCW sample oxidized at 1073 K and 7 Pa of pure oxygen atmosphere, can be rationalized with the evolution of the near-surface oxide structure due to the irradiation by the impinging D ions and subsequent annealing during the TPD experiments. First, the sample is a WO₃ oxide with some level of impurities on the surface which are released during first D implantation/TPD experiments and a thin W-rich layer is created on the surface of the oxide due to the oxygen erosion by impinging D ions. The retention properties of such oxide in the LAF state are stable (figure 9(a)) up to a fluence of a few 10 × 10²¹ D m⁻². With an increase of D irradiation in the oxide and subsequent migration of O during TPD annealing cycles, an amorphous oxide is formed below the surface enriched in W (figure 9(b)). Migration of O from the bulk of the original WO₃ induces creation of a second W-rich layer separating the crystalline and amorphous oxides (figures 9(c) and (d). The increased retention properties of such oxide in the HAF state are stable as a reduced dynamic retention at RT is observed.

## 8. Conclusions

This experimental study has investigated the retention properties of tungsten trioxide and its structural evolution under low-energy 500 eV $D_2^+$ irradiation.

A 45–160 nm-thick crystalline WO₃ oxide layer was grown at a temperature of 1073 K and pure oxygen pressure of 7 Pa. The formed trioxide is locally disordered and thermally stable when heated below the oxidation temperature (up to 750 K in this work).

Deuterium release up to 720 K from the as-oxidized sample is unstable until the oxide is cleaned from air contaminants. Deuterium retention in clean WO₃ as a function of incident

fluence is reproducible and contrary to pristine polycrystalline tungsten cannot be described by a single power law function. A transition to a ten-fold higher retention is evidenced upon irradiation of the trioxide with a fluence of deuterium higher than 4.6 × 10²¹ D m⁻² in our implantation conditions. The change in deuterium retention mechanisms is likely associated with near-surface structural modification of the oxide caused by sputtering of oxygen atoms by impinging deuterium ions and migration of oxygen from the bulk of WO₃ towards the oxygen-depleted surface. This results in a formation of two W-rich layers separated by an amorphous oxide in the sub-surface of the original WO₃.

This study provides essential data for the numerical modeling studies on hydrogen isotopes retention in tungsten plasma-facing components, particularly for high-fluence scenarios, where oxide-induced fuel retention could pose operational challenges. Our results indicate the importance of considering structural changes of the oxide on the surfaces of tungsten plasma facing components during plasma exposures, as well as subsequent to atmospheric oxidation during venting and maintenance of fusion devices.

## Acknowledgment

The project leading to this publication has received funding from the Excellence Initiative of Aix-Marseille University—A*Midex, a French ‘Investissements d’Avenir’ programme as well as from the ANR under grant ANR-18-CE05-0012. The work has been also carried out within the framework of the French Federation for Magnetic Fusion Studies (FR-FCM).

This work has been carried out within the framework of the EUROfusion Consortium, funded by the European Union via the Euratom Research and Training Programme (Grant Agreement No 101052200 - EUROfusion). Views and opinions expressed are however those of the author(s) only and do not necessarily reflect those of the European Union or the European Commission. Neither the European Union nor the European Commission can be held responsible for them.
Additionally, Mykola Ialovega was supported by the U.S. Department of Energy Award Number DE-SC00020428 and by discretionary funding of the Department of Nuclear Engineering & Engineering Physics of the College of Engineering at the UW-Madison. The authors are grateful to Prof. Oliver Schmitz affiliated with the UW-Madison for supporting the experimental work in Marseille that lead to this publication.

## ORCID iDs

Mykola Ialovega https://orcid.org/0000-0003-0041-8039

Celine Martin https://orcid.org/0000-0002-4704-3273

Cédric Pardanaud https://orcid.org/0000-0001-8974-4396

Elodie Bernard https://orcid.org/0000-0002-9291-7654

Marco Minissale https://orcid.org/0000-0001-6331-1402

Régis Bisson https://orcid.org/0000-0002-8819-1563

## References

[1] Pitts R. et al 2013 *J. Nucl. Mater.* **438** S48-S56

[2] Philipps V., Lyssoivan A., Sergienko G., Schulz C., Esser H., Freisinger M. and Samm U. 2008 *Proc. 22nd Fusion Energy Conf. (Geneva, Switzerland, 13-18 October)* vol FT/4-2Ra (available at: www-naweb.iaea.org/napc/physics/FEC/FEC2008/html/node546.htm#104764)

[3] Federici G. et al 2001 *Nucl. Fusion* **41** 1967-2137

[4] Philipps V., Neu R., Rapp J., Samm U., Tokar M., Tanabe T. and Rubel M. 2000 *Plasma Phys. Control. Fusion* **42** B293

[5] Guseva M., Suvorov A., Korshunov S. and Lazarev N. 1999 *J. Nucl. Mater.* **266-269** 222-7

[6] Tu H., Li C., Ding W., Shi L. and Zhang Y. 2022 *J. Nucl. Mater.* **562** 153590

[7] Pappalardo F. et al 2022 *Optics* **3** 216-24

[8] Alimov V., Tyburska B., Balden M., Lindig S., Roth J., Isobe K. and Yamanishi T. 2011 *J. Nucl. Mater.* **409** 27-32

[9] Pezzoli A., Dellasega D., Russo V., Gallo A., Van Emmichoven P. A. Z. and Passoni M. 2015 *J. Nucl. Mater.* **463** 1041-4

[10] Dunand A., Minissale M., Faure J.B., Gallais L., Angot T. and Bisson R. 2022 *Nucl. Fusion* **62** 054002

[11] Federici G. et al 2003 *J. Nucl. Mater.* **313-316** 11-22

[12] Federici G., Loarte A. and Strohmayer G. 2003 *Plasma Phys. Control. Fusion* **45** 1523-47

[13] Martin C., Diez M., Campos A., Cabié M., Giacometti G., Balden M., Gallo A., Pegourié B., Bernard E. and Tsitrone E. (the WEST team) 2021 *Phys. Scr.* **96** 124035

[14] Balden M. et al (the WEST Team) 2021 *Phys. Scr.* **96** 124020

[15] Addab Y. et al 2016 *Phys. Scr.* **2016** 014036

[16] Cifuentes S., Monge M. and Perez P. 2012 *Corros. Sci.* **57** 114-21

[17] Ivanov V.E. 1969 *High Temperature Oxidation Protection of Tungsten* (NASA)

[18] Dunn J.S. 1929 *J. Chem. Soc.* 1149-50

[19] Gulbransen E.A. and Andrew K.F. 1960 *J. Electrochem. Soc.* **107** 619-28

[20] Wriedt H.A. 1989 *Bull. Alloy Phase Diagr.* **10** 368-84

[21] Lassner E. and Schubert W.D. 1999 *Tungsten: Properties, Chemistry, Technology of the Element, Alloys and Chemical Compounds* (Springer) (https://doi.org/10.1007/978-1-4615-4907-9_3)

[22] Pardanaud C. et al 2020 *Nucl. Fusion* **60** 086004

[23] Psoda M., Rubel M., Sergienko G., Sundelin P. and Pospieszczyk A. 2009 *J. Nucl. Mater.* **386-388** 740-3

[24] Ping Y., Rocca D. and Galli G. 2013 *Phys. Rev. B* **87** 165203

[25] Liu X. and Fan H.Q. 2018 *R. Soc. Open Sci.* **5** 171921

[26] Chatten R., Chadwick A.V., Rougier A. and Lindan P.J.D. 2005 *J. Phys. Chem. B* **109** 3146-56

[27] Matsunami N., Ohno N. and Tokitani M. 2009 *J. Nucl. Mater.* **390-391** 693-5

[28] Ogorodnikova O., Roth J. and Mayer M. 2003 *J. Nucl. Mater.* **313-316** 469-77

[29] Bisson R., Markelj S., Mourey O., Ghiorghiu F., Achkasov K., Layet J.M., Roubin P., Cartry G., Grisolia C. and Angot T. 2015 *J. Nucl. Mater.* **467** 432-8

[30] Ghiorghiu F., Minissale M., Hodille E., Grisolia C., Angot T. and Bisson R. 2019 *Nucl. Instrum. Methods Phys. Res. B* **461** 159-65

[31] Kremer K., Schwarz-Selinger T. and Jacob W. 2021 *Nucl. Mater. Energy* **27** 100991

[32] Kremer K., Brucker M., Jacob W. and Schwarz-Selinger T. 2022 *Nucl. Mater. Energy* **30** 101137

[33] Kong X.S., You Y.W., Fang Q., Liu C., Chen J.L., Luo G.N., Pan B. and Wang Z. 2013 *J. Nucl. Mater.* **433** 357-63

[34] Hodille E. et al 2017 *Nucl. Fusion* **57** 076019

[35] Kremer K., Schwarz-Selinger T. and Jacob W. 2023 *Nucl. Mater. Energy* **35** 101406

[36] Pappalardo F., Rayneau L., Martin C., Cabie M., Salomon E., Angot T., Cartry G., Bisson R. and Minissale M. 2024 *Nucl. Mater. Energy* **41** 101751

[37] Ialovega M. 2021 Surface conditions of tungsten components: impact on hydrogen inventory *PhD Thesis* Aix-Marseille University theses.fr energie Rayonnement, Plasma

[38] Fulton G. and Lunev A. 2020 *Corros. Sci.* **162** 108221

[39] Pardanaud C. et al (EUROfusion WP PFC contributors) 2023 *Mater. Res. Express* **10** 102003

[40] Kim Y.K. et al 2021 *Electron-Impact Cross Sections for Ionization and Excitation Database* (National Institute of Standards and Technology) (available at: www.nist.gov/pml/electron-impact-cross-sections-ionization-and-excitation-database)

[41] Yoon J.S., Kim Y.W., Kwon D.C., Song M.Y., Chang W.S., Kim C.G., Kumar V. and Lee B. 2010 *Rep. Prog. Phys.* **73** 116401

[42] Mark T. and Egger F. 1976 *Int. J. Mass Spectrom. Ion Phys.* **20** 11

[43] Ghiorghiu F., Aissou T., Minissale M., Angot T., Temmerman G.D. and Bisson R. 2021 *Nucl. Fusion* **61** 126067

[44] Gusev M.I., Korshunov S.N., Suvorov A.L. and Lazarev N.E. 1999 *Tech. Phys.* **44** 1123-7

[45] Ziegler J.F., Ziegler M. and Biersack J. 2010 *Nucl. Instrum. Methods Phys. Res. B* **268** 1818-23

[46] Salomon E., Minissale M., Lairado F.R., Coussan S., Rousselot-Pailley P., Dulieu F. and Angot T. 2021 *J. Phys. Chem. C* **125** 11166-74

[47] Hoen M., Tyburska-Püschel B., Ertl K., Mayer M., Rapp J., Kleyn A. and van Emmichoven P.Z. 2012 *Nucl. Fusion* **52** 023008

[48] Martin C., Hijazi H., Addab Y., Domenichini B., Bannister M.E., Meyer F.W., Pardanaud C., Giacometti G., Cabié M. and Roubin P. 2017 *Phys. Scr.* **2017** 014019

[49] Schmid K., von Toussaint U. and Schwarz-Selinger T. 2014 *J. Appl. Phys.* **116** 134901

[50] Zavašnik J., Shvalya V., Kremer K., Schwarz-Selinger T. and Jacob W. 2025 *Appl. Surf. Sci.* **692** 162712