# Impact of (NH4)2SO4 agricultural atmospheric pollutant on the degradation mechanisms of thin layer Cu(In,Ga)Se2 solar cells

Adèle Debono, Noor Fikree, Amelle Rebai, Nathanaelle Schneider, Jean-François Guillemoles, Polina Volovitch

► To cite this version:

Adèle Debono, Noor Fikree, Amelle Rebai, Nathanaelle Schneider, Jean-François Guillemoles, et al. Impact of (NH4)2SO4 agricultural atmospheric pollutant on the degradation mechanisms of thin layer Cu(In,Ga)Se2 solar cells. Corrosion Science, 2025, 249, pp.112829. 10.1016/j.corsci.2025.112829. hal-05267696

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# Impact of (NH₄)₂SO₄ agricultural atmospheric pollutant on the degradation mechanisms of thin layer Cu(In,Ga)Se₂ solar cells

Adèle Debono<sup>a,b,●</sup>, Noor Fikree<sup>a,b</sup>, Amelle Rebai<sup>b</sup>, Nathanaelle Schneider<sup>b</sup>,
Jean-François Guillemoles<sup>b</sup>, Polina Volovitch<sup>a,*</sup>

## ARTICLE INFO

### Keywords:

Thin film solar cells

Molybdenum

Al-doped ZnO

CIGS

Atmospheric corrosion

Agrivoltaics

## ABSTRACT

Chemical environment is usually disregarded in stability evaluations of photovoltaic devices but it can be critical for agrivoltaics (dual land use for energy production and agriculture). This work considers for the first time the effect of (NH₄)₂SO₄ agricultural pollutant on the chemical degradation mechanisms of thin layer solar cells. Cu (In,Ga)Se₂ (CIGS) cells with the architecture SLG/Mo/CIGS/Zn(O,S)/ZnMgO/Al:ZnO/NiAlNi and representative stacks were characterized at different times of temperature/humidity cyclic aging with and without (NH₄)₂SO₄. The pollutant strongly increased degradation rate and modified the key degradation mechanisms. Cu(In,Ga)Se₂-absorber and its interface with Mo-back contact were the most affected by accelerated aging without (NH₄)₂SO₄, while corrosion of front NiAlNi contacts and pitting of window Al-doped ZnO layer was the key degradation mechanism in the presence of (NH₄)₂SO₄. Chemical modifications of the cell and layers were coherent with optoelectrical characteristics loss. The work implies the necessity to take into account specific agricultural pollutants in reliability evaluation of new technologies for agrivoltaic applications.

## 1. Introduction

Agricultural photovoltaics (Agri-PV), or agrivoltaics, is a practice combining crops and electricity production on the same plot of land. Agri-PV can solve not only the main problems of the land use faced by the deployment of the photovoltaic (PV) energy in the EU's Solar Energy Strategy [1] and worldwide [2], but a good management of sun irradiation can be mutually beneficial for crops and photovoltaic installations in certain climates [3]. High humidity and the presence of chemicals used in agriculture (fertilizers, pesticides, adjuvants, and so on) can strongly impact long-term stability of the materials [4,5]. If the effects of temperature and humidity in the presence of the most common naturally present atmospheric gases are taken into account by the PV community [6], the impact of agricultural atmospheric aerosol pollutants on the stability of photovoltaic cells stays unexplored. This is true both, for the dominating the market crystalline-Si technology [7] and for the less present on the market thin film chalcopyrite Cu(InₓGa(1-x))Se₂ (CIGS)--based photovoltaics [8]. CIGS indeed offers for AgriPV applications the advantage of mobility and adaptability thanks to its light weight and

possible flexibility.

Development of new flexible solutions for Agri-PV requires understanding of the effects of specific agricultural environment on the solar cell stability. In the present work we propose for the first time to consider the impact of farms atmospheres on the chemical stability and degradation mechanisms of CIGS solar cells and relevant materials and interfaces. These mechanisms are very complex because of multiple interfaces and various materials, including metals, oxides, semiconductors, used in a complex architecture of the CIGS cell [9], as briefly presented below (Fig. 1). The charge carriers (electrons/holes pairs) are produced and separated thanks to a p-n junction between the p-type semiconductor (absorber made of a polycrystalline chalcopyrite Cu (InₓGa(1-x))Se₂) and a thin n-type semiconductor (buffer, typically made of CdS (Fig. 1a) or less toxic alternatives like Zn(O,S) (Fig. 1b)). Grain boundaries in CIGS absorber are usually passivated by Na⁺ ions, which migrate to CIGS from the soda lime glass (SLG) substrate during the production process [10]. Extraction and collection of the charge carriers requires top and back contacts made of highly conductive materials. Back contact in CIGS is usually made from metallic molybdenum. The

<footnote><sup>a</sup> Chimie ParisTech-CNRS, PSL Research University, Institut de Recherche de Chimie Paris (IRCP), 11 rue Pierre et Marie Curie, Paris 75005, France</footnote>

<footnote><sup>b</sup> Institut Photovoltaïque d'Ile-de-France (IPVF), UMR 9006, CNRS, Ecole Polytechnique - IP Paris, Chimie ParisTech - PSL, 18 Boulevard Thomas Gobert, Palaiseau 91120, France</footnote>

<footnote>* Corresponding author.</footnote>

<footnote>E-mail address: polina.volovitch@chimieparistech.psl.eu (P. Volovitch).</footnote>
top side should however allow the photons to reach the absorber, implying the replacement of a metallic layer by a thin “window” layer of a transparent conductive oxide like aluminum-doped zinc oxide (AZO). Shunt paths (e.g. pin holes) can occur between the window layer and the absorber because the n-type buffer is usually very thin and/or does not cover the entire surface. To avoid pin holes, a very thin resistive layer (usually zinc oxide) is applied between the n-type buffer and the window layer (i:ZnO in Fig 1.a and ZnMgO in Fig 1.b). A top metallic grid contact (made of nickel/aluminum or silver) is sometime added to improve the electrons collection (Fig. 1b) while to interconnect CIGS cells in a module, 3 scribing steps (P1, P2 and P3 in Fig. 1a) are often applied.

Chemical evolution of CIGS cells leading to performance loss in hot humid atmospheres with the presence of oxygen gas and carbon dioxide, was previously correlated with the redistribution of Na⁺ during aging [10]. Carbon dioxide led to hydrozincite formation at grain boundaries of a ZnO window layer while the presence of oxygen enhanced corrosion of molybdenum back contacts [11–13]. Chemical evolution of the CIGS cells in the presence of other pollutants has not been studied yet. In the present work, we consider in detail the impact of ammonium sulfate (NH₄)₂SO₄, one of the most known aerosol pollutants in agricultural atmospheres [14], on the chemical stability and degradation mechanisms of all the layers and interfaces in CIGS solar cells. This pollutant is introduced in alkaline soils as N and S fertilizer allowing also a pH decrease and can volatilize into the atmosphere, especially if the soil has a high calcium carbonate content [15], leading to the aerosol formation. It is also present as a spray adjuvant for water-soluble herbicides, insecticides, or fungicides [16]. Furthermore, (NH₄)₂SO₄ can be formed as a secondary aerosol during the neutralization of ammonia gas (NH₃) in acid rains [17–19]. The agricultural sector accounts for more than 90 % of ammonia emissions due to animal manure, urea and ammonium-based fertilizers in Europe [15,20]. Understanding the effect of (NH₄)₂SO₄ on the PV systems installed in rural areas, especially for Agri-PV applications, is therefore very relevant. Several studies reported the accelerated failure of the encapsulation on the modules in the presence of gaseous ammonia [21,22] or sprayed ammonium salts (NH₄Cl or (NH₄)₂SO₄, see [23,24] respectively). These observations highlight that in ammonium containing environments encapsulation can be insufficient. It is necessary therefore to consider ammonia-based species for the stability studies of all the layers and interfaces in PV cells. High corrosivity of ammonium sulfate has been reported for some structural metals (zinc [25,26], copper [27,28] or aluminum [29]) which could be present in CIGS solar cells and is coherent with the high solubility of (NH₄)₂SO₄ in water (754 g·L⁻¹ at 20 °C), its hygroscopic properties and its acidifying effect [30]. Nevertheless, to the best of our knowledge, no research devoted to the chemical mechanisms leading to the cell performance loss in the presence of ammonia or ammonium ions has been published, except the effect on the encapsulation in the works mentioned below. At the same time, our previous work [31] has demonstrated strong detrimental effect of ammonium sulfate on the optoelectrical properties of the non-encapsulated CIGS cells during

temperature and humidity cycling. Without the pollutant, the aged cells lost about 37 % of the initial efficiency but partly recovered after rinsing; in the presence of the pollutant the cells lost 58 % of their initial efficiency and no recovery was observed. Moreover, optoelectrical properties of the individual layers or model stacks evolved differently when aged with and without pollutant, suggesting that the presence of the pollutant modifies the degradation mechanisms. More details of the measured optoelectrical characteristics of the cells can be found in **Supplementary Information** and in reference [31]. The chemical evolutions during aging with and without ammonium sulfate need therefore to be verified experimentally by the individual evaluation of chemical and physical modifications for each material with and without the pollutant. On the basis of the observed evolutions, the degradation mechanisms consistent with both chemical and optoelectrical characterizations need to be proposed.

In this work, for the first time we propose a systematic methodology allowing to survey and understand the mechanisms of chemical and physico-chemical modifications of CIGS solar cells in the presence of atmospheric aerosol pollutants relevant for agricultural application and to correlate these mechanisms with the performance parameters on the example of ammonium sulfate. For the selected agricultural pollutant, the work aims the following scientific objectives:

- Experimental verification if the selected pollutant affects the chemical structure of CIGS solar cell and if yes – experimental identification of the layers or interfaces for which the chemical stability is the most impacted by the presence of the salt and the experimental description of their chemical evolution.

- Suggestion of the chemical degradation mechanisms and verification of their coherence with the degradation of optoelectrical properties and performance of the cell.

## 2. Experimental

### 2.1. *Samples architecture and fabrication*

Both complete CIGS cells and representative model systems consisting of a few layers and interfaces, schematically represented in Fig. 1, were investigated. The substrate of all samples was a 3-mm soda lime glass (SLG). 600-nm molybdenum (Mo) back contacts were deposited by direct current magnetron sputtering. 2-µm Cu(In, Ga)Se₂ (CIGS) absorbers were co-evaporated (atomic ratios of CGI = Cu/(Ga+In) ~ 0.80 and GGI = Ga/(Ga+In) ~ 0.39 corresponding to a band gap of E<sub>g</sub> ~ 1.3 eV) [32]. 50-nm zinc oxide (i:ZnO) and 480-nm aluminum-doped zinc oxide (AZO) transparent conductive oxides were deposited by atomic layer deposition [33] (typical for CdS architecture presented in Fig. 1a) and 50/600/50-nm nickel/aluminum/nickel (NiAlNi) front contact by evaporation (typical for cells in Fig. 1b and for some modules). The cells (Fig. 1b) were manufactured on 5 × 5 cm² SLG substrate, and then scribed to isolate 10 cells per sample. To produce them, a 20

<figure><img src="image_2.png" /><figcaption>**Fig. 1.** Schematics of the sample configurations: (a) monolithically-integrated module, (b) complete cell on the left (SLG/Mo/CIGS/Zn(O,S)/ZnMgO/Al:ZnO/NiAlNi), (c) representative model stacks of layers corresponding to the configurations outlined in red in (a) and (b): SLG/Mo, SLG/Mo/CIGS, SLG/i:ZnO/AZO, SLG/i:ZnO/AZO/NiAlNi and lift-off procedure to reveal the Mo/MoSe₂/CIGS interface.</figcaption></figure>
nm- thick Zn(O,S) layer were chemical bath deposited, the 40 to 70 nm-ZnMgO and the 300 nm-AZO were sputtered. Finally, buried interface between Mo and CIGS (for SLG/Mo/CIGS samples) was dissociated thanks to a lift off process. The CIGS layer was glued with an ethyl cyanoacrylate glue on a piece of SLG for 2 min. A hammer blow broke the samples at the most fragile interface between MoSe₂ and CIGS. More details of the film preparation procedures can be found in *Supporting Information* (SI) and in the previous article [31].

## 2.2. Accelerated degradation protocol

Samples were aged in a Weiss WKL 100 climatic chamber between 8 and 14 days using alternating cycles of temperature and humidity between 25°C at 85 % Relative Humidity (RH) and 80°C at 30 % RH modelling day/night and seasonal variations. One cycle consisted of eight periods covering 24 h (programmed as 25°C and 30 % RH for 2 h; 25°C and 85 % RH for 2 h; 80°C and 30 % RH for 4 h; 25°C and 85 % of RH for 4 h; 25°C and 30 % RH for 2 h; 25°C and 85 % RH for 2 h; 80°C and 30 % RH for 4 h; 25°C and 85 % of RH for 4 h). Measured evolution of T°C and RH in this cycles and used climatic chamber can be found in our previous publications [23,24]. To evaluate the effect of humidity on the degradation, dry heat test ("DrH") with temperature at 80°C and 15 % RH was also performed during 7 days on SLG/Mo/CIGS samples.

Wet deposition of atmospheric aerosol was simulated by the daily deposition of a 100 µL drop of an aqueous solution containing 0.78 mmol.L⁻¹ of (NH₄)₂SO₄, representing about one year of NH₄⁺ deposition [34–37]. The amount of sulfate ions was dictated by the stoichiometry of (NH₄)₂SO₄. Drops were deposited in the first period of the 24 h dry–wet cycle (25°C and RH 30 % for 2 h) during the five first days of the ageing procedure. Each day the new drop was deposited on the same location as the previous day. For each material, the volume of the drop was fixed and the area of the surface occupied by the deposited drop was measured. This value varied between materials but stayed constant for the same material (0.5 ± 0.1 cm² for AZO, 1.1 ± 0.3 cm² for Mo and 0.8 ± 0.2 cm² for CIGS). The variation of the drop size between the materials can be related to the variation of their hydrophilicity. Additional samples from the same fabrication batch were used for comparison. The "Reference" samples were stored in a vacuum desiccator at room temperature, the "Blank" samples were aged under the same conditions as the "(NH₄)₂SO₄" but without drop deposition and the "H₂O" samples were aged with a drop of deionised water of the same volume as the pollutant drop for the "(NH₄)₂SO₄". (NH₄)₂SO₄ will also be referred as "pollutant" in the rest of the article. Samples name and aging conditions are summarized in Table 1.

## 2.3. Characterization of cells and layers

For each architecture and each aging condition at least three samples were aged in one experiment to verify the reproducibility. One series of aging was made for the full cells and for the model stacks two series of aging were reproduced. Characterizations of the aged samples were replicated in multiple points which was possible because of a relatively large surface area of each sample (at least 1.5 × 1.5 cm²).

The chemical evolution of individual layers and complete cells was studied by Raman spectroscopy, X-ray diffraction (XRD), scanning

electron microscopy (SEM) coupled with electron dispersive X-Ray spectroscopy (EDS) and glow discharge optical emission spectroscopy (GD-OES).

The visual aspect of the samples was inspected using a VXH Keyence microscope equipped with a VH-Z100R/W/T lens.

The Raman spectra were recorded at room temperature on a Renishaw InVia Raman microscope with green laser (λ=532 nm=2.33 eV – Nd-YAG double, maximum power 100 mW), with 1800 lines/mm grating and a Leica DM2500 x50 objective (NA 0.75). The laser power selection depended on the studied sample and is specified in the results section for each figure. The spectra were acquired for 1 min with accumulations of 15 or 30 s to avoid saturation of the detector as the laser power depends on the material absorption of light. The whole surface was examined in the optical microscope and hundreds of individual spectra and multiple maps were collected in multiple areas of the sample for each configuration and aging condition in order to verify correlations between typical surface appearances and their chemical compositions in different places, to ensure sufficiently representative results. Data acquisition and analysis (cosmic ray removal, background subtraction) were carried out using the software WiRE 4.2.

Bragg-Brentano XRD data were collected using a PANalytical X'Pert Pro X-ray diffractometer equipped with Cu Kα source, in the 2θ range of 8°-75° (step size 0.0263°, scan speed 0.058°/s). Grazing Incidence X-Ray Diffraction (GIXRD) patterns were taken with a Cu Kα sourced PANalytical Empyrean diffractometer (grazing incidence angle at 0.3°, 2θ range of 8°-75° with a step size of 0.2° at a scan speed of 0.04°/s).

Scanning Electron Microscopic (SEM) images were acquired with a Supra 35 (Zeiss) microscope at working distance of 10.5 mm at 15 kV. Energy Dispersive X-Ray Spectroscopy (EDS) was coupled with the SEM and data were recorded with a Bruker Xflash 6/60 detector at 15 kV and working distance of 10.5 mm by Quantax Esprit software. Quantification was done with atomic number effects (Z), absorption (A) and fluorescence (F) correction (ZAF) method and an average value was considered from the analyses made on at least 3 different zones with similar morphology.

Glow-discharge optical emission spectroscopy (GD-OES) profiles were recorded with a Horiba G-Profiler 2TM with a 2 mm² and 4 mm² anode after purging with nitrogen to improve detection in the UV region. The Argon plasma conditions were: 450 Pa, 20 W with asynchronous pulse at 3000 Hz, 0.3125 duty cycle to reduce heating effect on insulating substrates. The signal evolution for the elements of interested was recorded at 589.600 nm for Na, 386.416 nm for Mo, 324.759 for Cu, 451.138 nm for In, 294.368 nm for Ga, 196.094 nm for Se, 180.738 nm for S, 285.217 for Mg, 481.060 nm for Zn, 130.223 nm for O and 288.162 nm for Si. The signals were normalized by the total light emitted by the plasma (Fi). The signals of sulfur presented in Fig. 8 were smoothed (moving average of 10 units). At least two depth profiles were measured for each aged configuration.

X-Ray Fluorescence (XRF) was performed with a Fluorescence X Fisherscope® X-Ray XDV®-SDD spectrometer, driven by WinFTM software. The system was calibrated with a Cu sample and measurements were performed with a 1 mm-diameter collimator for 30 s. 5 measurements per configuration were acquired to have statistical data.

Open-Circuit Potential (OCP) measurements were performed using

<table><caption>**Table 1**

Samples' name with the corresponding aging conditions.</caption><thead><tr><th>Name of sample</th><th>Temperature</th><th>RH (%)</th><th>Drop of pollutant and concentration</th><th>Daily deposition [µg.cm<sup>-2</sup>.day<sup>-1</sup>]</th><th>Modelled environment</th></tr></thead><tbody><tr><td>Reference</td><td>RT</td><td>&lt; 30 %</td><td>No</td><td>-</td><td>Initial</td></tr><tr><td>Blank</td><td>Cycle (25°C/80°C)</td><td>Cycle (30%/85 %)</td><td>No</td><td>-</td><td>Control</td></tr><tr><td>H<sub>2</sub>O</td><td>Cycle (25°C/80°C)</td><td>Cycle (30%/85 %)</td><td>Yes<br>H<sub>2</sub>O</td><td>-</td><td>Control</td></tr><tr><td>(NH<sub>4</sub>)<sub>2</sub>SO<sub>4</sub></td><td>Cycle (25°C/80°C)</td><td>Cycle (30%/85 %)</td><td>Yes<br>(NH<sub>4</sub>)<sub>2</sub>SO<sub>4</sub> = 0.78 mmol.L<sup>-1</sup></td><td>[SO<sub>4</sub><sup>2-</sup>] = 15.0<br>[NH<sub>4</sub><sup>+</sup>] = 11.2</td><td>Rural</td></tr><tr><td>DrH</td><td>80°C</td><td>15 %</td><td>No</td><td>-</td><td>Dry and hot climate</td></tr></tbody></table>
<figure><img src="image_3.png" /><figcaption>**Fig. 2.** Optical micrographs of SLG/Mo and SLG/Mo/MoSe₂ samples: initial state (Day 0), during aging (Day 4) and final aspect (Day 8 or Day 10). Aging conditions are indicated: without drop (“Blank”), with H₂O drop (“H₂O”) or with “(NH₄)₂SO₄” drop. Images obtained without rinsing if not specified or after rinsing when specified. The scale is the same for all the images in each line.</figcaption></figure>

<figure><img src="image_4.png" /><figcaption>**Fig. 3.** Examples of (a) macroscopic pictures, (b) optical micrographs and (c) corresponding Raman spectra (laser power: 5 mW) of SLG/Mo samples before aging (“Reference”), aged for 8 days without drop (“Blank”, spectrum (3)) or with water drop (“H₂O”, spectra (1) and (2)), before rinsing.</figcaption></figure>

Gamry 3000 potentiostat equipped with the Gamry Framework and Gamry Echem Analyst software after a stabilization time of 2 hours. A two-electrode set-up consisted of the sample as a working electrode and

a silver/silver-chloride (Ag/AgCl in saturated KCl) as a reference electrode immersed in a (NH₄)₂SO₄ solution at 0.78 mmol.L⁻¹. The OCP values shown in the results and discussion sections are converted to the
<figure><img src="image_5.png" /><figcaption>Fig. 4. Examples of (a) macroscopic pictures, (b) optical micrographs and (c) corresponding Raman spectra (laser power: 1 mW) of SLG/Mo samples with (NH₄)₂SO₄ pollutant after 14 days aging, before and after rinsing, as indicated.</figcaption></figure>

reference with the standard hydrogen electrode (SHE) considering the potential of the Ag/AgCl potential $E_{Ag/AgCl} = +0.197$ V vs SHE.

**Contact conductivity** was evaluated with in-line 4-point probe (Lucas Labs Pro 4 test system) to measure the sheet resistance, $R_{sheet}$. For greater precision, measurements were performed at 20 different locations and the statistics of the results were expressed as the mean values and standard deviations.

## 3. Results

### 3.1. Effect of ammonium sulfate on the metallic Mo back-contact aging

#### 3.1.1. Effect of water drop on Mo oxides growth without pollutant

Visual evolution of the surface of molybdenum samples during aging in different conditions is presented in Fig. 2. Some typical appearances at higher resolution are also shown in Fig. 3. After 4 days of degradation (Fig. 2b), “Blank” sample is covered with brown species (previously identified as m-Mo<sub>8</sub>O<sub>23</sub> [24]) while the “H<sub>2</sub>O” sample presents some blue corrosion products at the border of the initial drop (Fig. 2f). After 8 days in Blank conditions (Fig. 2c), the sample is largely covered with rainbow iridescent areas while after the same time with H<sub>2</sub>O drop (Fig. 2g), brown is visible on the edge of the drop.

Raman spectra measured for zones of the Blank (upper figures in Fig. 3 a-c) and H<sub>2</sub>O drop (lower figures in Fig. 3 a-c) samples after 8 days of the test allow to identify different corrosion products. The iridescent areas reveal the presence of orthorhombic α-Mo<sub>7/2</sub>V<sub>1/2</sub>O<sub>3</sub> (main B<sub>1</sub> g peak at 808 cm<sup>-1</sup> [38]), together with MoO<sub>2</sub> (main peak at 728 cm<sup>-1</sup> [38]) and amorphous MoO<sub>3</sub> (main peaks at 864 and 945 cm<sup>-1</sup> [39]). Presented in Fig. 3c spectrum 3 looks very similar to the spectra reported for the Mo surface after damp heat test, which was ascribed in the literature as MoO<sub>2.88</sub> on the basis of XPS [40]. With the water drop (Fig. 3a down image), most of the surface looks metallic with some yellowish/iridescent areas showing the peaks typical for monoclinic m-MoO<sub>2</sub> and hexagonal h-MoO<sub>3</sub> (spectrum 2 in Fig. 3c). While the peaks positions of MoO<sub>2</sub> are slightly shifted for the “Blank” (spectrum 3 in Fig. 3c) they match perfectly in the presence of water drop (spectrum 2 in Fig. 3c). More details about the spectra identification can be found in Supporting information (SI) - Table S1.

The sheet resistance of the sample aged in the presence of a water drop (0.19 ± 0.01 Ω.sq<sup>-1</sup>) was less important than for the “Blank” sample aged without water drop (0.21 ± 0.07 Ω.sq<sup>-1</sup> [31]) but higher than for the reference (0.18 ± 0.01 Ω.sq<sup>-1</sup>, [31]). After rinsing the

electrical resistance was only slightly affected (0.21 ± 0.01 Ω.sq<sup>-1</sup> for “H<sub>2</sub>O” and 0.20 ± 0.01 Ω.sq<sup>-1</sup> for “Blank”). The decrease of the resistance after rinsing for Blank sample is consistent with partial dissolution of the oxides. The increase of the resistance after rinsing of “H<sub>2</sub>O” sample can be due to the thickness of conducting layer; the significance of these changes is however questionable because they are at the limit of the error bar. Rinsed surfaces (Fig. 2d and h) appear more homogenous, their initial visual aspect is partly restored, which could be also attributed to the dissolution of corrosion products. The hypothesis of formation of soluble corrosion products in Blank sample is also confirmed by the evolution of the ratios $R_{O/Mo}=at\%(O)/at\%(Mo)$ and $R_{Na/Mo}=at\%(Na)/at\%(Mo)$ before and after rinsing, calculated based on the results of the EDS measurements. Before rinsing, the Blank sample was enriched in oxygen ($R_{O/Mo}$ ranging from 0.75 to 2.5) and sodium ($R_{Na/Mo}$ reaching 0.8) compared to the Reference ($R_{O/Mo} = 0.05$ and $R_{Na/Mo}=0$). After rinsing, no sodium ($R_{Na/Mo}=0$) and less oxygen ($R_{O/Mo} = 0.50$ to 0.65) were detected on the surface.

#### 3.1.2. Effect of ammonium sulfate on aging of Mo

The presence of ammonium sulfate modified the surface chemistry of molybdenum. As shown in Fig. 4a-b, the visual aspect was different (dark edges where the drop was deposited with a wide variety of aspects inside the drop) and little affected by rinsing. Raman analysis demonstrated that the presence of (NH<sub>4</sub>)<sub>2</sub>SO<sub>4</sub> on the sample modified the corrosion species compared to the Blank sample (Fig. 4c). Over the entire sample, a narrowing of the 800–900 cm<sup>-1</sup> peak and a shift of its maximum from 850 to 870 cm<sup>-1</sup> could be observed after rinsing. The peak positions at 850–870 and 940–950 cm<sup>-1</sup> correlate well with ν(Mo-O-Mo) and ν(O=Mo) vibrations of hydrated molybdenum (VI) oxides MoO<sub>3</sub>·xH<sub>2</sub>O although distinguishing between all possibilities of hydrates is not easy [41,42]. For a very similar spectrum, Ajito *et al.* suggested a film of hydrated MoO<sub>3</sub> and a framework-type structure of anhydrous MoO<sub>3</sub> [39]. The oxides were identified as mainly MoO<sub>3</sub>·0.7 H<sub>2</sub>O before rinsing and as mainly MoO<sub>3</sub>·1.4 H<sub>2</sub>O after rinsing. The proposed identification is detailed in SI - Table S2. Increase in hydration of MoO<sub>3</sub> after rinsing is also coherent with the slight resistivity increase reported in [31] (from 0.20 ± 0.01 Ω.sq<sup>-1</sup> before rinsing to 0.22 ± 0.01 Ω.sq<sup>-1</sup> after rinsing).

GD-OES depth profiling presented in Supplementary Information (SI - Figure S1) confirmed the Mo surface oxidation and its recovery upon rinsing of Blank while “(NH<sub>4</sub>)<sub>2</sub>SO<sub>4</sub>” samples did not recover. For the Blank sample, the depth of the oxide layer was greatly reduced by
<figure><img src="image_6.png" /><figcaption>**Fig. 5.** Examples of visual aspect and corresponding GD-OES elemental profiles of SLG/Mo/CIGS samples. (1) –reference before aging; (2–4) - after 14 days of T°C - RH cycling without or with drop as indicated: (2)- blank; (3)- with (NH₄)₂SO₄ drop; (4) - with H₂O drop; sample (5) Dry heat - after 7 days at 80°C/15 % RH dry heat. Analyzed regions and associated anode size (4 or 2 mm) are schematically indicated by red circles. The signal of Na is scaled to the right Y-axis. Cu and Mo signals are normalized by their maximum intensities and correspond to the left Y-axis. Buried Mo/CIGS interface is highlighted in pink.</figcaption></figure>

<figure><img src="image_7.png" /><figcaption>**Fig. 6.** Normalized GIXRD spectra of samples aged in different conditions, as indicated (a) SLG/Mo/CIGS aged for 14 days (see peaks identification in SI – Table S3) and (b) full cells aged for 9 days (see peaks identification in SI – Tables S4–5).</figcaption></figure>
<figure><img src="image_8.png" /><figcaption>**Fig. 7.** (a, b) Optical micrograph of yellow zones on the aged for 9 days and rinsed cell and (c) representative Raman spectra, collected at laser power of 5 mW in the zones shown in (b) by the same number (spectrum (1) on a sample aged with (NH₄)₂SO₄ pollutant, spectrum (2) on a Blank, spectrum (1) of a reference of unaged AZO is given for comparison).</figcaption></figure>

<figure><img src="image_9.png" /><figcaption>**Fig. 8.** (a) Normalized GD-OES profiles of Zn, Mg, Cu, Mo and Si and (b) absolute intensity Na of complete cells (all data normalized on total light Fi), before aging (1.“Reference”) and after aging without (2.“Blank”) and with (3.“(NH₄)₂SO₄”) the pollutant, after rinsing. The p-n junction position is indicated by a green and the back Mo/CIGs interface by a pink bar.</figcaption></figure>

rinsing but it did not change after the rinsing of “(NH₄)₂SO₄” samples. Additionally, GD-OES demonstrated Na⁺ accumulation at the Mo/oxide interface and an effective removal of accumulated Na⁺ by rinsing for all the samples (SI - Figure S1).

In conclusion, when subjected to temperature and humidity cycling, the Mo surface is oxidized. Ammonium sulfate modifies the nature of corrosion products favoring higher oxidation states and more hydrated

structures but decreasing their solubility. Mo corrosion products can impact conductivity at the interface Mo/CIGS and sodium migration from SLG. According to the literature, Mo corrosion products and the oxide layer can affect the formation of MoSe₂, limit the migration of Na⁺ and decrease electrical conductivity at Mo/CIGS interface [43,44]. All these effects are detrimental for the good performance of the cell.
### 3.1.3. Alteration of the MoSe₂ layer

The previous section described the evolution of SLG/Mo samples during aging. As an intermediate layer of MoSe₂ is formed at the interface with CIGS during the samples production, it can modify the behavior of Mo. It is especially important when directly exposed to atmosphere at P3 scribing. To verify the effect of MoSe₂ on the Mo stability, SLG/Mo/MoSe₂ stacks were prepared by lift-off procedure described in Section 2.1 and aged for 14 days. The visual evolution of the surface without pollutant is presented in Fig. 2i-j. MoSe₂ is clearly much less affected by aging than bare Mo. In the areas covered by a drop containing (NH₄)₂SO₄ (Fig. 2k), no corrosion products were detected, with the exception of the (NH₄)₂SO₄ crystals formed by drying, which were eliminated by rinsing (Fig. 2l). No visible surface degradation in the presence of MoSe₂ layer is also coherent with the corrosion potentials of the SLG/Mo and SLG/Mo/MoSe₂ samples, measured as 0.2 V vs SHE and 0.9 V vs SHE respectively. MoSe₂ is therefore much nobler than Mo and can be considered as a passive film, acting as a barrier if the layer is intact.

## 3.2. Effect of the pollutant on SLG/Mo/CIGS interface

### 3.2.1. Chemistry and distribution of corrosion products generated by SLG/Mo/CIGS

When SLG/Mo/CIGS samples are aged in Blank conditions, dark dots were observed on the surface (see example of Picture 2 in Fig. 5a). A collection of corrosion products was detected inside the dots by Raman, including sodium molybdates Na₂MoO₄, Na₂Mo₂O₇ and copper selenides Cu(2-x)Se, similar to our previous work [24]. By comparison of Blank samples with the dry heat samples aged at 15 % RH and 85°C for 14 days (Fig. 5.5), which look intact, one can conclude that the presence of humidity is a necessary condition for the formation of black spots. Our previous work [24] reported that black dots formed on SLG/Mo/CIGS but not on SLG/CIGS in Blank conditions and suggested a degradation mechanism of formation of black dots in Blank conditions as a result of galvanic coupling between Mo and CIGS with the formation of molybdate species. The results of dry heat test in which no dots were formed corroborate with the electrochemical nature of this mechanism. The formation of sodium molybdates is consistent with the results of GD-OES (Fig. 5) and GIXRD (Fig. 6a). To facilitate the interpretation of the GD-OES profiles, one characteristic element of each layer was selected as a marker to define the layers of the structure: Cu for CIGS, Mo for back contact and Na for SLG. One can clearly see in Fig. 5 that inside the dot (Fig. 5b sample (2)) more Na is accumulated at the surface while the concentration of Na at the Mo/CIGS interface is reduced compared with outside the dot. The interface between Mo and CIGS is less defined inside the dot, as if the Mo had partly mixed with the CIGS layer. These observations are in agreement with the presence of Na and Mo in the corrosion products on the top surface. Additionally, the peak at 8.6° in the XRD spectra of the SLG/Mo/CIGS samples (Fig. 6a) matches well with Na₂Mo₃O₁₀·3 H₂O (JCPDS: 01-087-1785) or Na₄Mo₈O₂₆·12 H₂O (JCPDS: 00-035-1465). XRD peaks identification is consistent with the literature [45] and detailed in SI - Table S3-5.

The spots formation and growth are reduced in the areas in which a drop of water (sample 3) - Fig. 5b or of a solution containing pollutant was deposited (sample 4) - Fig. 5). Moreover, only Cu(2-x)Se (and no sodium molybdate) was detected by Raman in the spots developed inside the H₂O drop (not shown). These observations suggest that corrosion of Mo buried under CIGS layer is diminished in the presence of the electrolyte, in accordance with Section 3.1.1 for non-buried Mo. However, Na seems to be removed from CIGS layer in the zones inside the drop (sample 3) or sample (4) in Fig. 5b), indicating its diffusion to the surface in the form of soluble salts and washing by a thick liquid layer.

GIXRD results demonstrate (Fig. 6a) that the signals of both buried Mo and Mo corrosion products are more pronounced after temperature-humidity cycling in the absence of the liquid drop, supporting the

hypothesis that degradation of SLG/Mo/CIGS stack can be decreased due to the presence of a thick liquid layer. No additional crystalline corrosion products is formed in the presence of ammonium sulfate (the additional peak at 23.7° for “(NH₄)₂SO₄” sample corresponds to the residual (NH₄)₂SO₄ (JCPDS: 00-001-0363)).

The SEM cross sections did show neither significant modification of the CIGS layer nor formation of a thick corrosion products layer (SI - Figure S2). XRF results (SI - Table S6) demonstrated that the variation of the atomic content of Cu, Ga, In and Se was below 1 % and no clear trend could be identified. The average thickness of the CIGS layer was also similar for all the samples. These observations are coherent with a very localized degradation of CIGS.

For the full cells, the SLG/Mo/CIGS interface can be exposed near the scribing used to isolate each cell on the sample (Fig. 7a). In these regions some dark species were present even before aging and expanded during aging in Blank conditions. The Raman peaks associated with these species (spectrum 2 in Fig. 7c) were attributed to Na₂MoO₄·2 H₂O: the peaks at 333 cm⁻¹ corresponding to the vibration ν₄ (MoO₄), 818 and 854 cm⁻¹ to ν₃ (MoO₄), and 916 cm⁻¹ to ν₃ (MoO₄) [46]. The same peaks were observed on the top of SLG/Mo/CIGS sample after degradation before and after rinsing, which is consistent with such an identification. Compound Na₂MoO₄·2 H₂O is also coherent with the composition of the degradation products reported in the literature for SLG/Mo samples [47]. Other peaks correspond to CIGS (176 and 218 cm⁻¹ peaks), ZnO and AZO (505 and 573 cm⁻¹ peaks) [48–50] present for the full cells both before and after aging far from the scribing (spectrum 1 in Fig. 7c, identification in SI – Table S7).

### 3.2.2. Migration of elements through SLG/Mo/CIGS in model and full cell stack

Depth distribution of the chemical elements in SLG/Mo/CIGS structures and the full cells was obtained by GD-OES profiling and presented in Figs. 5b and 8 respectively. To define the interfaces and correlate the depth profiles with the layered structure, one characteristic element of each layer was selected as a marker of this one: Zn for AZO, Mg for ZnMgO, S for Zn(O,S), Cu for CIGS, Mo for back contact and Si for SLG. Once the layer's markers are defined sodium and molybdenum redistribution by aging can be seen. Na⁺ diffused from the SLG substrate during the samples preparation; for SLG/Mo/CIGS (Fig. 5) the Na signal is higher in Mo than in CIGS and increases at Mo/CIGS interface and CIGS surface. In the presence of a deposited drop (Fig. 5b 3–4), little Na signal is present on the surface and inside the CIGS layer if compared with the Reference or Blank samples. Therefore, a thick water layer on the surface certainly attracted and dissolved part of the Na⁺ initially accumulated in the CIGS layer or on the top surface. For the full cells, two important interfaces are marked in Fig. 8 in color boxes: the p-n junction in green and the buried CIGS/Mo interface in pink. Without pollutant the positions of Mg and S are recovering and indicate the p-n junction. For the “(NH₄)₂SO₄” sample, a very strong S signal is coming from the surface of the cell, due to the sulfur contained in the pollutant, thus Mg signal is used as a marker of the p-n junction.

The analysis of the figure shows that the aging promoted the migration of sodium from the Mo/CIGS interface to the p-n junction (Fig. 8b – Blank). To quantify this accumulation, the peak areas A of sodium at p-n junction Aₚₙ (in green in the figure) and Mo/CIGS interface Aₐ₉₆₄ (in pink in the figure) were calculated and their ratio Aₚₙ/Aₐ₉₆₄ was considered. This ratio was 0.11 for the Reference and 0.27 for the Blank sample, confirming that Na⁺ accumulation at p-n junction relatively increased after Blank aging. Nevertheless, compared to SLG/Mo/CIGS configurations (Fig. 5), no Mo migrated inside the CIGS layer. In the presence of the pollutant, sodium accumulated on the surface or inside the AZO layer rather than at the other interfaces (Fig. 8b – (NH₄)₂SO₄). Focusing on the CIGS elements, no significant change was observed on the ratio between Cu, Ga, In and Se (SI – Figure S2) which is also consistent with the XRF results (Section 3.2.1).

In conclusion, the drop limited the formation of dark dots on the top
surface and less corrosion products were found inside the zone on which a liquid (either water or the salt solution) was deposited. Sodium migration during aging was impacted by the presence of the drop: in Blank conditions $Na^+$ accumulated at the CIGS surface (in SLG/Mo/CIGS) and at the p-n-junction (in the cell) while in the presence of a drop it was removed from the surface of CIGS (or accumulated in the AZO layer for full cell stack).

### 3.3. Effect of pollutant on chemical and morphological characteristics of AZO window layer

Surface morphologies of the SLG/i:ZnO/AZO models and complete cells are presented in Fig. 9. The “Reference” samples (Fig. 9a) differ between the models and cells. The cell surface is rougher than the models due to the morphology of the substrate (a polycrystalline CIGS for cells compared to a flat SLG for models). The surface seems to be little impacted by aging in the absence of the pollutant (Fig. 9b). After aging with the pollutant both structures are covered by needle-like products (Fig. 9c). The EDS analysis confirmed that the corrosion products contain a high amount of oxygen as well as some sulfur (coming from the sulfate salt) and sodium (maybe coming from migration through the cell) in this product. The products also contain Zn. The product composition however cannot be properly quantified because its layer is too thin.

On the aged surface of SLG/i:ZnO/AZO samples, grey circular defects are visible only in the presence of pollutant (Figure 10.1). Inside these defects, the Raman peak at 490 cm⁻¹ induced by the doping of ZnO [50] was absent and only glass was detected [51], meaning that the oxide layer was locally completely dissolved (identification in SI – Table S8). Image analysis with Image J software was used to quantify the fraction of the surface on which the TCO was completely dissolved. About 4 % of the surface was dissolved, with a hole density of 37 holes/mm², each hole covering in average 700 µm².

The nature of corrosion products on top of SLG/i:ZnO/AZO was verified by Raman spectroscopy (Fig. 10). On the reference samples (Fig. 10 spectrum Ref), only soda lime glass substrate [51] and the zinc oxide layer [52] were detected (identification in SI – Table S8). On the Blank samples, no Zn(OH)₂ could be detected using Raman, although its formation was previously reported by other teams in damp heat conditions [53,54]. Corrosion products are visible on the samples aged with the pollutant (Fig. 10 zone and spectrum 2), especially at the border of the drop where the pollutant should be the most concentrated because of a thinner electrolyte layer. These products were not removed by rinsing,

thus they are insoluble in water. New Raman peaks appeared after aging, they were attributed to zinc hydroxysulfate (ZHS) $Zn_4(SO_4)(OH)_6 \cdot nH_2O$ with n = 1, 3, 4 or 5 [55], see details in SI – Table S8. The hexagonal platelets formed by the corrosion products shown in Fig. 9 are consistent with this identification [55,56]. The Raman spectrum of $Zn_4(SO_4)(OH)_6 \cdot 5H_2O$ best matches the observed peaks [55]. Complementary XRD and GIXRD spectra presented in SI – Figure S3 and Table S9 suggest however that the species could have a hydration number between n = 4 and 5 (peaks of $Zn_4(SO_4)(OH)_6 \cdot 4H_2O$ mixed with $Zn_4(SO_4)(OH)_6 \cdot 5H_2O$ peaks). EDS performed at the same place as Raman point (2) (not shown) revealed the presence of sulfur and an increased level of oxygen, consistent with sulfate species. As no nitrogen was detected, the possibility of $(NH_4)_2SO_4$ is ruled out, meaning that the EDS results are compatible with the proposed identification. The strong accumulation of charges during SEM observation of this product seems to demonstrate the insulating nature of ZHS, in coherence with its harmful effect on contact conductivity. Other sulfates were observed in areas where AZO was dissolved. The dark product in Fig. 10.1 was identified as a hydrated sulfate by the peaks at 983 cm⁻¹ ($v_1(A_1)$ of $SO_4^{2-}$) and 3450 cm⁻¹ (O-H vibration). These positions differ from that expected for the original $(NH_4)_2SO_4$ salt (975 and 3150 cm⁻¹). This sulfate could correspond to $Na_2SO_4 \cdot 10H_2O$ or possibly $K_2SO_4$. Sodium was detected by EDS but no potassium (SI – Figure S4), thus the first possibility seems to be the most probable. Formation of $Zn_4(SO_4)(OH)_6 \cdot nH_2O$ was also confirmed for the complete cell aged with $(NH_4)_2SO_4$ (spectra 1 and 3 in Fig. 7, peaks at 396, 461, 615, 967 and 1023 cm⁻¹ [57] and GIXRD presented in Fig. 6b). The level of hydration of ZHS remains between n = 4 and 5 (see SI – Table S5).

In conclusion, AZO stability in both model systems and full cells is therefore strongly compromised by $(NH_4)_2SO_4$ because of the formation of very different corrosion products if compared with aging without the pollutant.

### 3.4. Impact of $(NH_4)_2SO_4$ on the corrosion of NiAlNi contacts

A thin layer of oxides and hydroxides was present before degradation on the NiAlNi contact of both cells (Fig. 12) and SLG/i:ZnO/AZO/NiAlNi models (Fig. 11a) and identified by Raman as a mixture of bunsenite NiO [58] and α-Ni(OH)₂ [59] (detailed identification in SI – Table S10). This thin layer of oxides and hydroxides initially formed on the surface of nickel is consistent with air-formed passive layer observed by Hall et al. [60]. After degradation in the absence of the pollutant, the composition of the nickel oxide layer and its visual aspect were unaltered as

<figure><img src="image_10.png" /><figcaption>**Fig. 9.** SEM images taken on the top surface of the SLG/i:ZnO/AZO and cells (a) before aging, “Reference”, and (b) after aging in “Blank” conditions and (c) “$(NH_4)_2SO_4$” conditions after rinsing.</figcaption></figure>
<figure><img src="image_11.png" /><figcaption>**Fig. 10.** Raman spectra (laser power: 50 mW and 5 mW for (NH₄)₂SO₄) of species detected on SLG/i:ZnO/AZO with (NH₄)₂SO₄ pollutant in points 1 and 2 (identified as Zn₄(SO₄)(OH)₆·nH₂O), compared to "Reference", "SLG" substrate and "(NH₄)₂SO₄" salt, the intensity was normalized on value at 2000 cm⁻¹.</figcaption></figure>

<figure><img src="image_12.png" /><figcaption>**Fig. 11.** Microscopic images, SEM images and associated Raman spectra (laser power: 10 mW) of species detected on SLG/i:ZnO/AZO/NiAlNi samples (a) before aging, "Reference", (b) after 14 days of aging without pollutant, "Blank" and (c) after 14 days of aging with (NH₄)₂SO₄ pollutant after rinsing.</figcaption></figure>

<figure><img src="image_13.png" /><figcaption>**Fig. 12.** (a) Optical micrographs and (b) Raman spectra detected on the aged and rinsed cells in different areas as indicated. Spectra obtained on the samples aged for 9 days with (NH₄)₂SO₄ pollutant (spectrum 3) compared to its Reference and Blank (spectra 1 and 2). All the spectra were recorded at laser power of 5 mW.</figcaption></figure>
illustrated in spectrum 2 of Fig. 11c, as a proof of a successful passivation. In contrast, the “(NH₄)₂SO₄” samples have different visual aspect and corrosion products. In spectrum 3 of Fig. 11c, the main peak vibrating around 984 cm⁻¹ corresponds to nickel hydrates retgersite Ni (SO₄)·6 H₂O or morenosite Ni(SO₄)·7 H₂O [61,62]. An additional peak at 1028 cm⁻¹ can be explained by the formation of dwornikite Ni(SO₄)·H₂O [63,64]. This peak was systematically accompanied by the peak at 984 cm⁻¹ of retgersite or morenosite. The SEM image of the zone 3 in Fig. 11 shows a surface hollowed by pitting corrosion. Indeed, atomic ratios between Ni and Ni+Al measured by EDS as 8.3 ± 0.1 % outside this zone and 7.5 ± 0.3 % inside it, confirming a small thinning of the Ni layer. Moreover, sulfur was detected by EDS only inside the degraded zone. Sulfate species are generally soluble and should be in principle removed by rinsing. This removal can be however incomplete if the species are accumulated in deep pits.

The 620, 982 and 1060 cm⁻¹ Raman peaks attributed to Ni(SO₄)·6 H₂O or Ni(SO₄)·7 H₂O were also detected on the NiAlNi contact of the complete cells aged with (NH₄)₂SO₄ (spectrum 2 in Fig. 12) confirming the consistency between the models and the cells. To sum up, the presence of pollutant destabilized the passivation layer, inducing the formation of soluble nickel sulfates.

# 4. Discussion

## 4.1. Impact of pollutant on the corrosion of CIGS contacts

### 4.1.1. Molybdenum back contact

Depending on the aging conditions, different Mo oxides were detected on the surface, see Table 2. The difference in the Mo behavior under the pure water drop and under the (NH₄)₂SO₄ aqueous solution could be understood considering the properties of the Mo oxides and the electrochemical mechanisms of Mo corrosion.

The difference between the Blank and the H₂O samples can be explained by the thick water layer formed by the drop, blocking O₂ diffusion (Fig. 13.a-b). The importance of O₂ as an oxidant agent on Mo was previously demonstrated [13]. The presence of oxygen in the electrolyte has multiple roles. First, at potential close to the OCP of Mo, the major cathodic reaction is oxygen reduction (reaction 1)

$$O_2 + 2H_2O + 4 e^- \rightleftharpoons 4 OH^- \quad (1)$$

meaning that higher the O₂ concentration, higher corrosion rate. Second, both dissolved oxygen and pH increase contribute to the oxidation of Mo and favoring the highest oxidation state of Mo (in the oxide film, because Mo(VI) becomes soluble at high pH, thus also increasing corrosion rates by chemical reactions 2.a and 2.d) or electrochemical reactions 2.b and 2.c.

$$Mo + O_2 \rightleftharpoons MoO_2 \quad (2.a)$$

$$MoO_2 + 2 (1-x)OH^- \rightleftharpoons MoO_{3-x} + (1-x)H_2O + 2 (1-x)e^- \quad (2.b)$$

$$MoO_2 + 4 OH^- \rightleftharpoons MoO_4^{2-} + 2 H_2O + 2e^- \quad (2.c)$$

<table><caption>**Table 2**

Synthesis of experimental and literature data about Mo aging.</caption><thead><tr><th>Aging condition of samples</th><th>Ref</th><th>H<sub>2</sub>O</th><th>Blank</th><th>(NH<sub>4</sub>)<sub>2</sub>SO<sub>4</sub></th></tr></thead><tbody><tr><td>Experimental surface chemistry (major species are shown in bold)</td><td>Mo</td><td>MoO<sub>2</sub><br/>+ h-<br/>MoO<sub>3</sub></td><td>α-MoO<sub>3</sub><br/>+ MoO<sub>2</sub> + a-<br/>MoO<sub>3</sub></td><td>MoO<sub>3</sub>·<br/>xH<sub>2</sub>O</td></tr><tr><td>Experimental R<sub>sheet</sub> before rinsing [Ω·sq<sup>-1</sup>] [31]</td><td>0.18<br/>± 0.01</td><td>0.19<br/>± 0.01</td><td>0.21 ± 0.07</td><td>0.20<br/>± 0.01</td></tr><tr><td>Experimental R<sub>sheet</sub> after rinsing [Ω·sq<sup>-1</sup>] [31]</td><td>-</td><td>0.21<br/>± 0.01</td><td>0.20 ± 0.01</td><td>0.22<br/>± 0.04</td></tr></tbody></table>

$$MoO_3 + 2 OH^- \rightleftharpoons MoO_4^{2-} + H_2O \quad (2.d)$$

In H₂O sample, thick water layer caused by the drop limits the diffusion of atmospheric O₂ and can therefore slow down corrosion of Mo, while in Blank sample the oxygen diffusion from the atmosphere to the surface is faster. This is schematically presented in Fig. 13 b by two arrows showing oxygen reduction reactions for the Blank sample in place of only one arrow shown in Fig. 13 a and Fig. 13 c for the samples with a drop. Indeed, Mo(IV) species (MoO₂) was present on the Mo exposed with water drop (Table 3), but the peaks were shifted and only Mo(VI) oxides were clearly detected for Blank.

Interestingly, sodium molybdates were found on the Blank SLG/Mo/CIGS samples. In these samples Mo was buried under CIGS layer. Galvanic coupling with CIGS, confined zone and dissolution of Na⁺ accumulated at the p-n junction and in CIGS grain boundaries could probably contribute to the conductivity and pH increase of the electrolyte, thus favoring transformation of MoO₃ into MoO₄²⁻ by reaction 2.d. The aqueous solubility being in orders of magnitude higher for molybdates than for MoO₃ (at room temperature it is 0.49 g·L⁻¹ for MoO₃ [65], 2.9 g·L⁻¹ for Na₂Mo₂O₇ [66] and 840 g·L⁻¹ for Na₂MoO₄ [67]), Mo contact corrosion could be locally accelerated in the cell configuration in Blank condition once the upper layers allow the access of humidity till Mo.

The value of the open circuit potential of Mo measured as 0.2 V vs SHE in (NH₄)₂SO₄ solution is consistent with the literature [68]) and confirms that the major cathodic reaction could be oxygen reduction. In the presence of (NH₄)₂SO₄, the medium is however expected to be slightly acidic because of the hydrolysis of ammonium salt formed from a strong acid and a weak base (reaction 3).

$$[NH_4]^+ + 2 H_2O = NH_3·H_2O + [H_3O]^+ \quad (3)$$

According to the Pourbaix diagrams of Mo reported in the literature [69,70], Mo(VI) can form a hydrated MoO₃·H₂O (molybdic acids) in acidified solutions. This product was indeed detected with Raman spectroscopy (see Fig. 4). Finally, the differences between (NH₄)₂SO₄ and Blank samples could be explained by two opposite effects of the drop containing a soluble salt (Fig. 13.a and c). Although O₂ ingress is slowed down because of the liquid layer, as previously discussed, the presence of SO₄²⁻ and NH₄⁺ ions the solution results in a higher conductivity and lower pH in a solution with the dissolved salt than in a thin electrolyte layer in the blank condition or inside of a pure water drop, thus accelerating corrosion. One can therefore explain why, in the presence of (NH₄)₂SO₄, the surface is more degraded than with a deionized water drop and only MoO₃ and no MoO₂ are detected. Taking into account that MoO₃·H₂O is less conductive than MoO₃ [41], stronger conductivity loss can be also understood. The results shown in Section 3.1 have also demonstrated that after selenisation of the Mo, the oxidation was much slower (Fig. 2,j-l). This can be explained by the much nobler OCP potential of MoSe₂ (0.9 V vs SHE) than metallic Mo (0.2 V vs SHE). MoSe₂ can be hence considered as a barrier passivating layer for Mo in humid condition. In the presence of defects in this layer localized corrosion could be expected.

To conclude on the chemical stability of the back contact, the corrosiveness of the aging conditions for Mo back contact increased in the order “H₂O” < “Blank” < “(NH₄)₂SO₄”. The barrier layer of MoO₂ is the main product in the absence of oxygen, the diffusion of which may be limited by the drop (“H₂O”). It can be oxidized into soluble Mo(VI) species if oxygen access is not limited (“Blank”). With ammonium sulfate (“(NH₄)₂SO₄”), oxidation is accelerated because of the electrolyte conductivity increase and the pH decrease.

Difference in the formed corrosion products can explain the observed differences in the performance loss of the layer between aged samples. R<sub>sheet</sub> increase after aging with a pure water drop, reported in Table 3, can be explained by the formation of the oxide layer, composed mainly of MoO₂ which has lower conductivity (3.10⁵ S·m⁻¹ [71,72]) than bare
<figure><img src="image_14.png" /><figcaption>**Fig. 13.** Schematics of proposed corrosion mechanisms relevant for CIGS solar cells contacts ((a-c) Mo, (d) AZO and (e-f) NiAlNi) aged with and without $NH_4_2SO_4$ as indicated.</figcaption></figure>

<table><caption>**Table 3**
Synthesis of data about Al-doped ZnO aging.</caption><thead><tr><th>Aging condition of samples</th><th>Ref</th><th>Blank</th><th>(NH<sub>4</sub>)<sub>2</sub>SO<sub>4</sub></th></tr></thead><tbody><tr><td>Experimental surface chemistry</td><td>Al doped ZnO</td><td>Al doped ZnO</td><td>Al doped ZnO partially dissolved, presence of zinc hydroxysulfates</td></tr><tr><td>Experimental R<sub>sheet</sub> before rinsing [Ω. sq<sup>-1</sup>] [31]</td><td>50.1 ± 0.7</td><td>59.7 ± 0.7</td><td>72.9 ± 3.0</td></tr><tr><td>Experimental R<sub>sheet</sub> after rinsing [Ω.sq<sup>-1</sup>] [31]</td><td>-</td><td>58.7 ± 0.5</td><td>76.8 ± 9.8</td></tr></tbody></table>

Mo ($10^7 \text{ S.m}^{-1}$ [73]). As MoO$_2$ is insoluble [65], R$_{sheet}$ cannot recover by rinsing in H$_2$O condition. In Blank condition, MoO$_3$ was the main corrosion species. This oxide has lower conductivity ($10^{-5} - 10^{-7} \text{ S.m}^{-1}$ [41]) than MoO$_2$, resulting in stronger R$_{sheet}$ increase before rinsing for Blank than in H$_2$O conditions. Nonetheless, R$_{sheet}$ is partially restored by rinsing (Table 3) because MoO$_3$ is soluble in water ($0.49 \text{ g.L}^{-1}$ [65]). Considering the cell performance loss, the sensitivity of the back contact to humidity is less critical than for the front contact due to the architecture of the cell. Moreover, its sheet resistivity was very little impacted by aging even with pollutant (Table 2). The conductivity of the layer is affected by humidity because of the formation of insulating Mo oxides. If they are soluble, Mo can partly restore its initial conductivity (Figs. 2d and 2h). After aging with $(NH_4)_2SO_4$, the formed Mo oxides are less soluble and cannot be removed by rinsing (Fig. 4), thus the contact could not be restored.

### 4.1.2. Transparent conductive oxide (Aluminum-doped Zinc Oxide)

Considering the impact of TCO on the cell performance, the aging without pollutant is significantly shorter (336 h) than reported in the

literature damp heat tests (typically 1000 h) thus its surface chemistry was little impacted by temperature and humidity cycles. The small increase in the sheet resistance (Table 3) could not be explained by the transformation of ZnO into Zn(OH)$_2$, usually reported for longer exposure time [54,74]. Therefore, this test represents only an earlier stage of the degradation of the TCO.

In the presence of $(NH_4)_2SO_4$, two degradation phenomena were observed for TCO: AZO dissolution and its transformation into zinc hydroxysulfates $Zn_4(SO_4)(OH)_6 \cdot nH_2O$ (ZHS) (also illustrated in Fig. 13. d). ZHS are well known zinc corrosion products [14], penhydrate or tetrahydrate being the most stable phases at RH< 50 % and below 30°C. Nevertheless, they have a narrow stability domain (pH between 6.5 and 8 [56]) and are not chemically stable in some environments. ZHS are expected to be transformed into Zn(OH)$_2$ or ZnO if the pH increases or dissolved into soluble zinc complexes if $SO_4^{2-}$ concentration increases [56,75]. At advanced degradation stages, the adsorption of CO$_2$ can also lead to the formation of carbonates [11,12]. Thus, the duration of aging test (9 to 14 days) could be insufficient to reach a stable corrosion product. The presence of ZHS does not inhibit the dissolution of zinc oxide, thus recovery of the initial oxide is not expected [75]. Localized dissolution of AZO [23] is expected in the presence of ammonia (Fig. 13. d) because of the formation of soluble complexes (for instance by reaction 4):

$$ZnO + 2 [NH_4]^+ \rightleftharpoons [Zn(NH_3)_2]^{2+} + H_2O \quad (4)$$

Surface analysis of AZO evidenced not only oxide dissolution but also formation of new corrosion product, ZHS ($Zn_4(SO_4)(OH)_6 \cdot nH_2O$). This product can be formed for instance by dissolution-precipitation processes at cathodic zones with higher pH by a reaction sequence (4) - (5.a)

$$4 [Zn(NH_3)_2]^{2+} + 6OH^- + SO_4^{2-} + nH_2O \rightleftharpoons Zn_4(SO_4)(OH)_6 \cdot nH_2O + 8NH_3 \uparrow \quad (5.a)$$
or by transformation of the oxide in the presence of sulfate

detected (Fig. 11c, Fig. 12). This is consistent with the Pourbaix diagram of Ni ([14], see also Supplementary Information where E-pH diagram for

$$4\text{ZnO} + 2\text{H}_2\text{O} + 2[\text{NH}_4]^+ + \text{SO}_4^{2-} + \text{nH}_2\text{O} \rightleftharpoons \text{Zn}_4(\text{SO}_4)(\text{OH})_6 \cdot \text{nH}_2\text{O} + 2\text{NH}_3\uparrow \quad (5.b)$$

Finally, the inspection of the surfaces of the aged cells presented the same corrosion species as the aged materials alone, showing that the degradation of TCO occurs in the layer and the degradation mechanisms studied on a single layer are relevant for the full cells and mini-modules.

The observed chemical degradation mechanisms can explain difference in the evolution of the optoelectrical characteristics of the TCO in the cells, aged in different conditions. Both, local dissolution of the AZO and its transformation into insulating ZHS (low conductivity of which is demonstrated by the strong charge accumulation visible by white bril- liance in the SEM image) should be detrimental to the cell conductivity and transparency, which is coherent with the data presented in Table 3 and consistent with the literature [12,23]). The transformations also result in an increase of the surface roughness (Fig. 9) due to the for- mation of needle-like corrosion products together with local dissolution of AZO (Fig. 13.d). This corrosion - induced roughness can explain the optical losses (reduced amplitude of interference fringes of trans- mittance of the AZO) and therefore the loss of the external quantum efficiency (EQE) of the cells reported in the previous work [31]. The low solubility of ZHS in water (pKs=16.6 [76]) and the presence of holes prevent the recovery of the properties after rinsing. The local dissolution of the AZO layer could lead to shunt paths, resulting in a reduction in the R<sub>sh</sub> and fill factor (the ratio of maximum obtainable power to the product of the open-circuit voltage and short-circuit current).

### 4.1.3. NiAlNi front contact

In NiAlNi contacts, aluminum is responsible for the good conductive properties and nickel for the corrosion resistance. The standard potential of the couple Ni<sup>2+</sup>/Ni is -0.257 V against -1.662 V for the couple Al<sup>3+</sup>/Al [9]. Therefore, if the upper layer of Ni is perforated and the buried Al becomes exposed to an electrolyte, its corrosion is expected to be enhanced by galvanic coupling with Ni. Corrosion kinetics can be however influenced by the formation of an Al<sub>2</sub>O<sub>3</sub> passivation layer on the Al surface, making it less susceptible to galvanic corrosion. In practice, the surface of Ni seems to be passivated by insoluble [77] nickel oxides and hydroxides (NiO and α-Ni(OH)<sub>2</sub>) (Fig. 11a) which are well-known passive layer components for Ni alloys (Fig. 13.e) [60], see Pourbaix diagram of Ni [78]. After cycling only temperature and humidity, the surface was visually and chemically similar (Fig. 11b). The sheet resistivity remained stable during the test in blank conditions (Table 4), confirming the success of nickel passivation. Rinsing did not impact the conductivity nor the surface composition, consistent with the low solubility of nickel oxides and hydroxides.

In contrast, the behavior of “(NH<sub>4</sub>)<sub>2</sub>SO<sub>4</sub>” sample correlates well with the fact that nickel is very sensitive to sulfate ions [14]. In the tests with ammonium sulfate, highly soluble nickel sulfates Ni(SO<sub>4</sub>)·nH<sub>2</sub>O were

<table><caption>**Table 4**

Synthesis of data about Ni aging.</caption><thead><tr><th>Aging condition of samples</th><th>Ref</th><th>Blank</th><th>(NH<sub>4</sub>)<sub>2</sub>SO<sub>4</sub></th></tr></thead><tbody><tr><td>Experimental surface chemistry</td><td>Ni, NiO, α-Ni(OH)<sub>2</sub></td><td></td><td>NiSO<sub>4</sub>·6H<sub>2</sub>O</td></tr><tr><td>Experimental R<sub>sheet</sub> before rinsing [Ω·sq<sup>-1</sup>] [31]</td><td>0.046 ± 0.003</td><td>0.043 ± 0.002</td><td>*</td></tr><tr><td>Experimental R<sub>sheet</sub> after rinsing [Ω·sq<sup>-1</sup>] [31]</td><td>-</td><td>0.043 ± 0.002</td><td>*</td></tr></tbody></table>

Ni was calculated for 10<sup>-2</sup> M sulfate concentration). Corrosivity of the sulfate salt for nickel should be enhanced when ammonia sulfate is used. This can be due, from one side, to the pH decrease by the hydrolysis of ammonia (reaction 3), which according to the Pourbaix diagram of Ni in the presence of the sulfate (see for example [14]) favors transformation of nickel hydroxide into soluble sulfate salts and, from other side, to the destabilization of the passive film by the formation of a soluble complex [Ni(NH<sub>3</sub>)<sub>6</sub>]<sup>2+</sup> (formation constant of 6 × 10<sup>8</sup> at 25°C)

$$\text{Ni(OH)}_2 + 6[\text{NH}_4]^+ + 2\text{H}_2\text{O} \rightleftharpoons [\text{Ni(NH}_3)_6]^{2+} + 4[\text{H}_3\text{O}]^+ \quad (6)$$

Difference in the chemistry of the surface film formed during aging with and without sulfate ions is schematically presented in Fig. 13 e and f. This difference can explain strong difference in the electrical conductivity after aging with and without the pollutant. Ni(SO<sub>4</sub>)·nH<sub>2</sub>O compounds are much less conductive (10<sup>-7</sup>·10<sup>-8</sup> S·m<sup>-1</sup> [79]) than initial Ni contact (Ni: 10<sup>7</sup> [80], NiO: 10<sup>-2</sup>·10<sup>-8</sup> [81,82], α-Ni(OH)<sub>2</sub>: 10<sup>-13</sup>·10<sup>-14</sup> [59]) and are certainly responsible for the significant conductivity loss (Table 4). The high solubility of Ni sulfates (406–750 g·L<sup>-1</sup> [83]) may explain the pitting corrosion of nickel, visible in Fig. 11c. The fact that conductivity was not restored by rinsing after the aging with the pollutant (Table 4) could be explained by the Al exposure to the atmosphere due to the Ni thinning or by a limited elimination of resistive soluble species trapped in the pits.

To sum up, for all the three types of the contacts (front contact, TCO and back contact), both degradation intensity and mechanism seem to be strongly impacted by the presence of (NH<sub>4</sub>)<sub>2</sub>SO<sub>4</sub>, if compared to only temperature and humidity aging. This observation suggests that, from the chemical point of view, the loss of CIGS cell electrical efficiency after aging in the presence of (NH<sub>4</sub>)<sub>2</sub>SO<sub>4</sub> can be due to the enhanced contacts corrosion, leading to the increase of the series resistance and therefore the fill factor. Among the three studied types of contacts, the front contacts are the first to meet the pollutant when encapsulation fails, thus their corrosion could be the most critical for the cell performance.

<caption>**Table 5**</caption>

Observed migration in different samples and its possible impact on the cell properties.

<table><thead><tr><th>Elements</th><th>Aging condition</th><th>Na</th><th>Mo</th></tr></thead><tbody><tr><td rowspan="4">SLG/Mo/ CIGS (Fig. 5)</td><td>Ref</td><td>Accumulation on the CIGS surface and Mo/CIGS interface</td><td>No detected migration</td></tr><tr><td>Blank</td><td>High migration to the CIGS surface in dots (Na<sub>3</sub>Mo<sub>4</sub>O)</td><td>Mild migration through CIGS in dots</td></tr><tr><td>H<sub>2</sub>O</td><td>leading to shunt paths and R<sub>sh</sub> loss</td><td></td></tr><tr><td>(NH<sub>4</sub>)<sub>2</sub>SO<sub>4</sub></td><td>Low concentration in CIGS</td><td>No detected migration</td></tr><tr><td rowspan="4">Cell (Fig. 8)</td><td>Ref</td><td>No Na<sup>+</sup> on the surface and low concentration in CIGS (Na<sup>+</sup> washed away by the drop)</td><td>No Mo detected either in CIGS nor on the surface</td></tr><tr><td>Blank</td><td>Accumulation at Mo/CIGS interface</td><td>No detected migration</td></tr><tr><td>(NH<sub>4</sub>)<sub>2</sub>SO<sub>4</sub></td><td>Migration to p-n junction leading to recombination at p-n junction and V<sub>oc</sub> loss</td><td>No detected migration</td></tr><tr><td></td><td>Migration to AZO or surface</td><td>No detected migration</td></tr></tbody></table>

<footnote>* Too resistive to be measured</footnote>
## 4.2. Elements migration during aging and its consequences for the cell performance

The observed element migrations and possible impact of each element redistribution on the cell performance are summarized in Table 5.

The interface between CIGS and Mo in GD-OES profiles is much less well defined inside the dots for SLG/Mo/CIGS configuration (Fig. 5.2b). It is possibly due to a small migration of negatively charged soluble molybdates, at the origin of the formation of sodium molybdates observed inside the dark dots. This effect was not observed for complete cells thus one should be careful when looking only at a reduced model, the deposition of additional layers can slow down, stop or hinder the migration of the elements.

Sodium is often described as an element easily migrating in the form of $Na^+$ ions, playing a detrimental role in the cell electrical stability, as demonstrated in other studies [43,45,84,85]. In SLG/Mo models (Fig. 5b), $Na^+$ accumulates just below the oxide layer, thus the thickness of this layer limits $Na^+$ diffusion to the CIGS and the power conversion efficiency of the cells [44]. Humidity seems to be the driving force of $Na^+$ migration through the CIGS to the top surface. In dry heat test, no spots were formed on the samples (Fig. 5.5a). In Blank conditions, a thin water layer was induced by cyclic conditions, promoting the migration of $Na^+$ to the surface of CIGS (in SLG/Mo/CIGS, Fig. 5.2) or to the p-n junction (in full cells, Fig. 8b). Under the drop, which represents a thick layer of the electrolyte, $Na^+$ migrated to the CIGS surface, where it was dissolved in water and therefore “washed out” (Figs. 5.3 and 4) [43,86] or continued across the p-n junction to the AZO inside the cells (Fig. 8c).

Differences in the sodium migrations between the layers and interfaces of the CIGS cells aged by temperature and humidity cycling without and with the pollutant are schematically sum up in Fig. 14, which evidenced gradual sodium release from the CIGS cell in the presence of the pollutant. This sodium redistribution which can be responsible for the efficiency reduction by multiple mechanisms: creation of defects promoting recombination or formation of shunt path. The accumulation of $Na^+$ at p-n junction and CIGS layer in Blank sample can generate some deep defect states that increase recombination [87] or reduce the hole density by removing the $V_{Cu}$ acceptor defects [88] leading to the loss of the open circuit voltage ($V_{oc}$) of the cell reported in literature [31]. The sodium molybdates $Na_2MoO_4$, $Na_2Mo_2O_7$ and copper selenides $Cu_{(2-x)}Se$ detected inside the dark dots can also act as recombination center, impacting $V_{oc}$.

Additionally, two types of possible shunt paths were observed in the aged cells due to corrosion and elements migration. First, near the scribing P3 revealing Mo to atmosphere, a dark corrosion product identified as $Na_2MoO_4 \cdot 2 H_2O$ (Fig. 7a) expended with aging. This conductive product ($10^5 S \cdot m^{-1}$ [89]) can be a good candidate for shunt path. Second, the formation of $Na^+$ rich-spots and accumulation of $Na^+$ at the p-n junction can be also responsible for shunting [84], greatly impacting shunt resistance ($R_{sh}$) and $V_{oc}$ [45,84].

With the pollutant, $Na^+$ accumulated less in CIGS and at p-n junction. This is consistent with the hypothesis that the $V_{oc}$ and $R_{sh}$ degradation is less important for the aging in the presence of pollutant than in Blank conditions, proposed in the previous work [31].

The chemical modifications reported in this work confirm the loss of performance related to the CIGS layer and p-n junction in the case of Blank, as described in this section, and related to contact corrosion and optical loss reported in the presence of the pollutant, as described in Section [4.1].

## 5. Conclusion

For the first time in the photovoltaic research, atmospheric aerosols are considered as stress factors susceptible to impact the stability of PV cells. The results presented in this work are based on the corroboration between chemical and morphological characterizations and

<figure><img src="image_15.png" /><figcaption>Fig. 14. Schematic representation of sodium migration between layers and interfaces of CIGS cell during temperature and humidity aging without (a) and with (b) the presence of $(NH_4)_2SO_4$ drop.</figcaption></figure>

optoelectrical properties loss. To conclude, the presence of ammonium sulfate during temperature and humidity cycling significantly increases degradation and modifies chemical degradation roads in the Cu(In,Ga) $Se_2$ - based photovoltaic cells and relevant materials.

In Blank conditions (heat and moisture only), sodium migration leading to the optoelectrical degradation was the most important mechanism. Sodium accumulation at the p-n junction can be interpreted from the optoelectrical point of view as defects formation affecting $V_{oc}$. Sodium was also accumulated in dark spots on top of SLG/Mo/CIGS structures. These dark spots could act as potential shunt paths affecting $R_{sh}$. The fact that $Na^+$ in black dots was present in the form of soluble molybdates suggests that its migration can be governed by the migration of molybdate anions (a charge compensation effect).

The fact that dark spots formation on the surface was limited in the zones delimited by a liquid drop containing or not $(NH_4)_2SO_4$ can be considered as an indirect confirmation of sodium migration mechanism promoted by degradation of Mo/CIGS interface. Indeed, thick liquid layer on the top of CIGS is expected to limit $O_2$ diffusion to the Mo/CIGS interface decreasing both Mo corrosion and cathodic pH increase, necessary for the formation of soluble molybdates.

During aging with $(NH_4)_2SO_4$, the main degradation was due the enhanced corrosion of the contacts - metallic Mo and NiAlNi and transparent conductive Al-doped zinc oxide. Localized corrosion of these materials with formation of insulating corrosion products such as zinc hydroxysulfates on AZO, $Ni(SO_4) \cdot nH_2O$ on NiAlNi and $MoO_3 \cdot nH_2O$ on Mo led to significant conductivity loss. Contact corrosion was negligible and the surface chemistry of the contacts tested in Blank condition was different, confirming that the enhanced corrosion of the contacts in the presence of $(NH_4)_2SO_4$ is due to the modification of the contacts corrosion mechanisms by the pollutant. Localized dissolution of AZO and formation of insoluble zinc hydroxysulfates was also possible thanks to $(NH_4)_2SO_4$ which can also explain rapid degradation of optical properties.

The proposed methodology taking into account potential aerosol pollutants effect on the reliability of thin film photovoltaic cells can be considered as a systematic study method which can be extended to different thin film technologies and applied to verify the effect of specific atmospheric pollutants expected in the environments imposed by emerging photovoltaic applications. As an example, chlorides (e.g. $NaCl$, $CaCl_2$) need to be considered for offshore and nearshore floating PV and specific fertilizers, fungicides etc. for Agri-PV.
CRediT authorship contribution statement

**Debono Adèle:** Writing – review & editing, Writing – original draft, Visualization, Validation, Supervision, Methodology, Investigation, Formal analysis, Data curation. **Fikree Noor:** Writing – review & editing, Validation, Investigation. **Rebai Amelle:** Investigation. **Schneider Nathanaelle:** Writing – review & editing, Validation, Supervision. **Guillemoles Jean-François:** Writing – review & editing, Validation, Supervision. **Volovitch Polina:** Writing – review & editing, Writing – original draft, Validation, Supervision, Methodology, Formal analysis, Conceptualization.

Declaration of Competing Interest

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

The authors wish to thank Gustavo Lins de Oliveira for his work concerning the aging and characterization of NiAlNi electrical contacts and Jiarui Wang for his contribution to the aging and characterization of MoSe₂ samples.

The study was supported by the French Government Program of Investment for the Future "Investissements d'Avenir" (ANR-IEED-002-01) via a free access to the IPVF Platform. A.D. thanks the Doctoral School ED 388 of the PSL University for the PhD fellowship.

Appendix A. Supporting information

Supplementary data associated with this article can be found in the online version at doi:10.1016/j.corsci.2025.112829.

Data availability

Data will be made available on request.

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