# Enhanced performance of reversible solid oxide cells using high-surface-area nanostructured thin-film oxygen electrodes

Adeel Riaz, Kosova Kreka, Alexander Stangl, Fjorelo Buzi, Silvère Panisset, Federico Baiutti, Laetitia Rapenne, Carmen Jiménez, David Jauffres, Michel Mermoux, et al.

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## To cite this version:

Adeel Riaz, Kosova Kreka, Alexander Stangl, Fjorelo Buzi, Silvère Panisset, et al. Enhanced performance of reversible solid oxide cells using high-surface-area nanostructured thin-film oxygen electrodes. Journal of Power Sources, 2025, 650, pp.237492. 10.1016/j.jpowsour.2025.237492. hal-05408807

HAL Id: hal-05408807

https://hal.science/hal-05408807v1

Submitted on 10 Dec 2025

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# Enhanced performance of reversible solid oxide cells using high-surface-area nanostructured thin-film oxygen electrodes

Adeel Riaz<sup>a,●</sup>, Kosova Kreka<sup>b</sup>, Alexander Stangl<sup>a</sup>, Fjorelo Buzi<sup>b,●</sup>, Silvère Panisset<sup>a,c,●</sup>, Federico Baiutti<sup>b</sup>, Laetitia Rapenne<sup>a</sup>, Carmen Jiménez<sup>a</sup>, David Jauffres<sup>c</sup>, Michel Mermoux<sup>d</sup>, Albert Tarancón<sup>b,e,●</sup>, Mónica Burriel<sup>a,* ●</sup>

## HIGHLIGHTS

* Highly porous nanostructured La₂NiO₄+δ oxygen electrode thin films optimized.

* Exceptionally low ASRₚol of 0.07 Ω cm² at 600 °C.

* Outstanding electrochemical performance in commercial half cells with La₂NiO₄+δ thin film electrodes.

* Critical raw materials reduction by ~10–15 times (g cm⁻²) in the oxygen electrode.

## ARTICLE INFO

**Keywords:**

Solid oxide cells

Thin films

Oxygen electrode

Fuel cells

Electrolyzers

Critical raw materials

Lanthanum nickelate

## ABSTRACT

Reversible solid oxide cells (rSOCs) are highly efficient electrochemical energy conversion devices which provide a promising pathway to green energy challenges. An essential factor for achieving high performance and enabling their commercialization is the selection of the oxygen electrode material, as it significantly impacts polarization resistance at intermediate-to-low temperatures (< 700 °C). La₂NiO₄+δ (L2NO4) stands out as a promising material due to its mixed ionic and electronic conductivity, high oxygen exchange activity and low activation energy. In this work, L2NO4 nanostructured thin films were first optimized in symmetrical cells reaching polarization resistance values as low as 0.07 Ω cm² at 600 °C. The films were then integrated into state-of-the-art commercial button cells and tested in both fuel cell (SOFC) and electrolysis (SOEC) modes revealing extremely high performance, with a power density of 1.2 and 1.3 W cm⁻² at 0.7 V at 670 and 725 °C, respectively, in SOFC mode, and a current density of 1.2 A cm⁻² at 1.3 V and 670 °C in SOEC mode. The superior efficiency compared to commercial cells emphasizes the nano-engineering strategy for enhancing performance while minimizing the use of critical raw materials. This is achieved by employing approximately 10–15 times less oxygen electrode material (g cm⁻²), a crucial factor for the sustainable development of these devices.

## 1. Introduction

The hydrogen economy holds immense importance as a transformative solution to replace fossil fuels and mitigate the environmental impacts of conventional energy sources. Hydrogen's role as an energy vector is pivotal in reducing fossil fuel dependency as it provides a clean and versatile alternative energy that can be harnessed, stored and used

across various sectors including transportation, industry, and power generation. By adopting a closed-loop system, where hydrogen is produced through renewable energy-powered electrolysis and consumed in fuel cells, a carbon-free energy cycle can be achieved. Among different types of fuel and electrolysis cells, Solid Oxide Cells (SOCs) stand out as highly promising for diverse applications due to their significantly higher efficiency - up to 20–30 % greater than other types of fuel and

<footnote><sup>a</sup> Univ. Grenoble Alpes, CNRS, Grenoble INP, LMGP, Grenoble, 38000, France</footnote>

<footnote><sup>b</sup> Catalonia Institute for Energy Research (IREC), Jardins de les Dones de Negre 1, Sant Adrià del Besòs, 08930, Barcelona, Spain</footnote>

<footnote><sup>c</sup> Univ. Grenoble Alpes, CNRS, Grenoble INP, SIMaP, 38000, Grenoble, France</footnote>

<footnote><sup>d</sup> Univ. Grenoble Alpes, Univ. Savoie Mont Blanc, CNRS, Grenoble INP, LEMMI, Grenoble, 38000, France</footnote>

<footnote><sup>e</sup> Catalan Institution for Research and Advanced Studies (ICREA), Barcelona, Spain</footnote>

<footnote>* Corresponding author.</footnote>

<footnote>E-mail address: monica.burriel@grenoble-inp.fr (M. Burriel).</footnote>
electrolysis cells - and their ability to operate reversibly in both modes with the same cell structure.

SOCs are highly efficient energy conversion solid state electrochemical devices with an oxide electrolyte, and ceramic oxide-based or cermet (ceramic-metal) electrodes. They can operate in two modes depending on the direction of flow of electrons and ions. The two modes are the fuel cell mode (SOFC), in which clean electrical power is produced from fuels such as hydrogen, with only water as the by-product, and the electrolysis mode (SOEC), which uses electrical power to produce fuels such as green hydrogen from water [1,2]. The reversible solid oxide cell (rSOC) system has recently garnered significant interest for its capability to alternate between SOFC and SOEC modes. In electrolysis mode, it consumes electrical power to produce oxygen and hydrogen fuels, while in fuel cell mode, it generates electrical power as needed. This makes it a potential candidate for building a future low-carbon energy supply system [3,4].

An important aspect for high performance and commercialization is the choice and amount of critical raw materials (CRMs) used in SOCs. Typical oxygen electrode compositions include La₁ₓSrₓCoO₃-δ (LSC), La₁ₓSrₓFeO₃-δ (LSF), La₁ₓSrₓCo₁_yFe_yO₃-δ (LSCF), which consist of CRMs such as La, Sr, Co, Ce, and Gd. These electrode materials need to be replaced, reduced and/or recycled [5] for industrial scale implementation to decrease the impact of CRMs in commercial SOCs. One of the strategies for reducing the use of CRMs is to use thin film oxygen electrodes [6]. There are several ways to nanoengineer these electrodes which include methods such as grain boundary engineering [7–9], strain engineering in thin films [10–13], multi-layered composite films [14, 15], vertically aligned nanostructures (VANs) [16–18] and nanostructures with high porosity [19–21]. All these cells show good performance; however, the scalability of these complex nanostructures remains a challenge.

The performance of rSOCs is usually limited by the electrochemical activity and stability of the oxygen electrode materials [3,22–27], thereby necessitating an in-depth understanding of the limiting factors in the oxygen reduction reaction (ORR) and the oxygen evolution reaction (OER). In SOECs, another factor leading to performance degradation is the delamination of the oxygen electrode due to the high oxygen partial pressures (pO₂) at the electrode/electrolyte interface [28]. Moreover, the high equivalent pO₂ at the electrode determines a depletion of the oxygen vacancy concentration and a potential drop in the oxygen evolution reaction kinetics for materials where ionic transport is dominated by oxygen vacancies [29]. Therefore, for rSOCs, the choice of oxygen electrode materials, their microstructure and long duration cyclic stability in both SOFC and SOEC modes need to be optimized.

Among the different oxygen electrode materials, La₂NiO₄+δ (L2NO4), a Ruddlesden-Popper phase material with alternating layers of perovskite and rock-salt structure, is one of the best performing electrodes at intermediate to low operation temperatures. Although bulk L2NO4 has been tested as oxygen electrode in SOCs [30,31], up to now perovskites have been preferred in commercial cells due to the instability of L2NO4 at high fabrication and operating temperatures (>700 °C) [32]. However, by using thin film deposition techniques that require relatively low deposition temperatures, and operating at low temperatures, these problems can be largely mitigated, making L2NO4 a highly promising alternative for oxygen electrodes. Thin film deposition by metal organic chemical vapor deposition drastically reduces the fabrication temperature of the cells to 600 °C, which enables the use of high performance materials such as L2NO4 and avoids the formation of insulating secondary phases. Additionally, L2NO4 is able to accommodate oxygen as interstitials (over-stoichiometry), which can prevent an increase in the pO₂ at the interface and avoid delamination of the oxygen electrode layer in SOEC mode [33,34] making it an ideal candidate for rSOCS.

Thin films of L2NO4 (20–500 nm) deposited at relatively low temperatures (< 650 °C) have shown high performance on symmetrical cells

[20,35–37] but have not yet been tested in operation conditions in SOCs. In a previous study [20], we demonstrated the improved electrode performance of nano-engineered thin films of L2NO4 grown on model single crystal substrates. Such improvement was attributed to a porous nano-columnar structure which increased the specific surface area, exposing surfaces with faster reaction kinetics and defects such as kinks and edges resulting in an enhancement of both the intrinsic and apparent activity of the electrode. In addition, the analysis of Electrochemical Impedance Spectroscopy (EIS) results combined with Finite Element Modeling (FEM) showed that an optimal film thickness exists, which corresponds to the transition from surface limited regime for thin films to bulk diffusion limited regime for thick films [37]. However, the performance of these electrodes within full cell and under polarization had not yet been assessed. This last study also demonstrated that it was possible to optimize the nanostructure of the L2NO4 electrodes by varying the deposition conditions leading to improved performance.

In this study, the thickness of the porous nanocolumnar L2NO4 thin films was first optimized on symmetrical cells and the performance of the thin films was subsequently tested on state-of-the-art commercial fuel electrode-supported cells. Electrochemical performance was measured in both SOFC and SOEC modes at different temperatures. The results revealed very high performance of the oxygen electrode which can enable the cells to operate at temperatures as low as 600 °C.

## 2. Experimental methodology

### 2.1. Thin film deposition

Symmetrical L2NO4 thin films were grown by pulsed-injection metal organic chemical vapor deposition (PI-MOCVD) on commercial double side polished Yttria-Stabilized Zirconia (YSZ) single crystal substrates (Crystec) with (100) orientation. The precursor solution was prepared by dissolving commercial (Strem chemicals) powders La(tmhd)₃ and Ni (tmhd)₂ (tmhd = 2,2,6,6-tetramethylheptane-3,5-dionate) in m-xylene solution (Alfa Aesar) with a La/Ni ratio of 4.0 and a solution concentration of 0.02 mol l⁻¹. The solution was injected into the reactor using a frequency of 4 Hz, opening time of 2 ms and a deposition (substrate) temperature of 600 °C. The carrier gas concentration was fixed at 34 % Ar/66 % O₂ with a total pressure of 5 Torr. The thickness of the films was controlled by the number of pulses injected. To evaluate the performance in fuel cell and electrolysis modes, commercial state-of-the-art Ni-YSZ anode-supported button cells (20.5 mm diameter; ≈250 µm-thick Ni-YSZ anode and ≈9 µm-thick YSZ electrolyte, SolydEra S.p.A.) were employed. The gadolinium-doped ceria (CGO, Ce₀.₈₀Gd₀.₂₀O₁.₉₅) barrier layer (BL) deposition conditions are comprehensively described elsewhere [38]. The oxygen electrode active area on the fuel electrode-supported cell was 1.45 cm². For current collection a gold layer was brush painted on the oxygen electrode side and Ni paste was brush painted on the fuel electrode side and each side was heated to evaporate the organic components and fabricate porous films. Fine gold and nickel mesh were used for better current collector on the oxygen electrode and fuel electrode, respectively.

### 2.2. Structural characterization

The structural characterization and phase analysis of the L2NO4 films deposited was carried out by X-Ray Diffraction (XRD) in Bragg Brentano (θ-2θ) configuration using a Cu-Kα (λ = 1.54 Å) source on a Bruker D8 Advance diffractometer. The surface morphology and cross-section imaging were done using an FEG Gemini 500 scanning electron microscope (SEM) in secondary (SE) and backscattered electron (BSE) modes with an acceleration voltage of 3 kV and a working distance of 4 mm. Ni-YSZ/YSZ/CGO/L2NO4 TEM specimens were prepared in cross-sections using the semi-automated polishing tripod technique with a MultiPrep™ system (Allied High Tech Products, Inc.). A precision ion polishing system (PIPS) II from GATAN was used for the final polishing.
TEM and high resolution TEM (HRTEM) images were recorded with a JEOL JEM 2010 LaB6 microscope operating at 200 kV with a 0.19 nm point-to-point resolution. STEM EDX analysis and HAADF STEM images were recorded with a JEOL 2100F FEG microscope.

## 2.3. Functional characterization

The functional characterization of the symmetrical cells was carried out using electrochemical impedance spectroscopy (EIS) using a home-built setup with a tubular furnace. A gold current collection layer was painted and a fine gold mesh was used. EIS measurements were carried out using a Biologic SP-200e potentiostat between 1 MHz and 100 mHz using a sinusoidal voltage perturbation of 50 mHz in a symmetric dry air atmosphere at different temperatures. The area specific resistance (ASR) of the electrode was calculated by subtracting the electrolyte contribution and dividing by 2 due to the symmetrical configuration.

Full cells were measured using a NorECS ProboStat™ placed inside a tubular furnace. The temperature of the furnace was calibrated using an external thermocouple. The cells were sealed using a ceramic paste (Ceramabond-Aremco) and heated at 1 °C/min up to 750 °C. The sealing process requires a plateau of 2 h at 93 °C and another 2 h at 260 °C. EIS measurements were carried out using a Parstat potentiostat and the j-V curves were measured using a Keithley 2700 and a power supply. SOFC measurements were carried out under dry synthetic air on the oxygen electrode side (flow of 140 ml min⁻¹·cm⁻²) and dry hydrogen at the fuel electrode (flow of 56 ml min⁻¹·cm⁻²). SOEC measurements were performed using a 50/50H₂/steam composition at the fuel electrode (flow of 22 ml min⁻¹·cm⁻² for both) and dry synthetic air at the oxygen electrode (flow of 112 ml min⁻¹·cm⁻²).

# 3. Results and discussion

## 3.1. Optimization of L2NO4 thin film electrode

L2NO4 nanocolumnar films with thicknesses of 500 and 1000 nm were deposited by Pulsed Injection Metal Organic Chemical Vapor Deposition (PI-MOCVD) at 600 °C on YSZ single crystal substrates in a symmetric configuration. MOCVD is a widely used deposition technique in the semiconductor, optoelectronics, and energy industries for

producing high-quality, uniform thin films with precise control over composition and thickness. This fabrication strategy can therefore be easily scaled up to a commercial level. A first optimization step consisted of a decrease of the deposition temperature from 650 °C [20,37] to 600 °C in order to create a nanocolumnar dendritic morphology, i.e., a more porous film with a larger active surface area. Fig. 1 (a and b) show the surface morphology of L2NO4 films deposited on YSZ single crystal, revealing a highly porous microstructure for both films. The 500 nm film surface shows triangular shaped grains (Fig. 1 (a)) while the 1000 nm film shows large grains which consist of smaller grains merged together (Fig. 1 (b)). The cross-section images in Fig. 1 (c & d) show a nano-columnar structure of the film with open porosity. The 500 nm thick film grows with a specific nanostructure which mostly consists of V-shaped columns. These columns evolve into a dendritic morphology in the 1000 nm thick film consisting of a dense core column with smaller structures branching out, largely increasing the specific surface area of the film.

Fig. 2 shows the XRD pattern of the L2NO4 thin films deposited on YSZ single crystals. The films grow polycrystalline due to the high lattice mismatch between the *a* cell parameter of L2NO4 (a/√2 = 5.47 Å) and the cubic unit cell of the YSZ substrate (5.14 Å). L2NO4 grows with a mean tetragonal structure (ICDD: 01-076-6116) mostly observed for the highest hyper-stoichiometries [39]. Some impurity peaks that can be attributed to lanthanum carbonate (La₂CO₃) are also observed in both films, which originates from an excess of La in the precursor solution, which may react with decomposition by-products of the organometallic precursors during the deposition. However, this impurity is not expected to significantly impact electrode performance, although its removal could further enhance it.

To determine the optimal film thickness for electrochemical performance the area specific polarization resistance (ASR~p~) of the symmetrical cells was evaluated by Electrochemical Impedance Spectroscopy (EIS). The EIS Nyquist plots for the two samples are shown in Fig. 3 (a) and were fitted with an equivalent circuit corresponding to the elementary reactions taking place at different regions in the electrodes. The Nyquist plots for the two samples are shown in Fig. 3 (a). The impedance spectra of the two symmetrical cells are distinctly different with the 1000 nm L2NO4 cell showing two semi-circle arcs at high frequency (HF) and low frequency (LF) regions, while the cell with 500

<figure><img src="image_4.png" /><figcaption>**Fig. 1.** SEM images showing the surface morphology and cross-section nanostructure of (a, c) 500 nm thick, and (b, d) 1000 nm thick L2NO4 films deposited on YSZ substrates.</figcaption></figure>
<figure><img src="image_5.png" /><figcaption>Fig. 2. XRD patterns of the L2NO4 films deposited on YSZ single crystal substrates. YSZ single crystal substrate peaks are marked with green squares according to ICDD: 00-030-1468. L2NO4 peaks are labelled with red circles according to ICDD: 01-076-6116 and La₂CO₅ peaks are marked with yellow triangles according to ICDD: 00-037-0804. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)</figcaption></figure>

nm L2NO4 has an extended arc which can be further characterized with an additional feature in the medium frequency (MF) region. These impedance spectra could not be fitted with a Gerischer element which is typically used for porous mixed ionic-electronic conductors [40,41] indicating that the electrode is not limited by the ionic oxygen diffusion in the bulk of the electrode, but rather by the surface exchange reaction. For such surface limited electrodes, the impedance spectra can be fitted with a series of Randles' circuits - a resistor and a constant phase element connected in parallel (R-CPE). Each R-CPE corresponds to a physicochemical process taking place at the electrode. To avoid over-fitting and to determine the number of R-CPE required for fitting the data accurately the EIS data was further analyzed by distribution of relaxation times (DRT). DRT does not require prior knowledge of the number of processes involved as it deconvolutes the impedance spectra by the characteristic relaxation frequency (or time constant) of each contribution. Here, we do not use DRT for quantitative analysis, but rather to verify the number processes taking place in the electrode, and hence, the number of R-CPE that are used for fitting the EIS data. The DRT of the two symmetrical cells are shown in Fig. 3 (b). There are 4 peaks for the 500 nm film. The peak at high frequency can be attributed to a DRT artefact [42], which is supported by the fact that 3 R-CPE elements are sufficient to model the impedance spectra. The artefact peak could arise from parasitic effects of non-ideal inductance and a low coefficient of the CPE at high frequency [43]. Apart from this artefact peak there are 3 peaks corresponding to the processes taking place in the electrode labelled P1 (at HF), P2 (MF) and P3 (LF). For the 1000 nm L2NO4 there are only 2 peaks (P1 at HF and P2 at LF). Accordingly, the Nyquist plots from EIS are fitted with 3 and 2 R-CPE in series for the 500 and 1000 nm L2NO4 cells, respectively. The total ASRₚ extracted from EIS at different temperatures for the two samples is shown in the Arrhenius plot in Fig. 3 (c). The ASR values show an enhancement in the electrode performance going from 500 nm to 1000 nm with a decrease in polarization resistance from 0.29 Ω cm² to 0.07 Ω cm² at 600 °C.

The microstructural features of the 1000 nm L2NO4 film could play a role in the higher performance of the electrode with smaller nano-columnar grains and dendrite-like features, increasing the number of active sites for oxygen redox reaction to take place. Furthermore, we previously observed that apart from an increase in the surface area, the intrinsic activity of the electrode is also enhanced due to an increase in the defects (kinks, edges etc.) on the lateral surface of the nano-columns

[20,44]. Going to a dendritic structure, as seen in the 1000 nm film, further increases the number of these defects leading to a sharp increase in the electrochemical activity of the electrode. To further understand the mechanisms contributing to the reaction kinetics, the pO₂ dependence of the best performing electrode (1000 nm L2NO4) was measured at 600 °C. The DRT of this film at different pO₂ is shown in Fig. 4 (a). Within the whole analyzed pO₂ range, the cell has only two distinguishable features in the EIS spectra, whereas the characteristic time scales vary only slightly with decreasing pO₂. As shown in Fig. 4 (b), the specific capacitance of the high frequency feature (P1) is in the range of 10⁻³ F cm⁻². These values are in agreement with reported values in L2NO4 dense thin films and porous bulk electrodes for charge transfer and oxygen incorporation [35,45]. The specific capacitance of the LF arc (P2) increases with pO₂ and is in the range of 0.05–2 F cm⁻². This high capacitance cannot be associated with a double layer capacitor. The activation energy for this process is 1.16 eV, as shown in Fig. 4 (c). Therefore, this process is not related to gas diffusion, as in such case it should be almost independent of the temperature. This process could be associated with molecular oxygen adsorption at the surface of the electrode [35,43]. This is in agreement with the ASR.

Commonly, the pO₂ dependence of the ASR is used to obtain better insight into the underlying elementary reaction [35,45,46]:

$$ \frac{1}{R_p} \propto p_{O_2}^n \qquad 1 $$

The value of the slope, n, of the ASR as a function of pO₂ can give insight into the oxygen species involved and the possible rate-limiting processes. The observed n value contains direct and indirect contributions. In a simplified manner three regimes are distinguished, whereas for (1) and (2) the direct contribution governs the pO₂ dependence:

(1) n ~ 1: the reaction mechanism occurs once per O₂ molecule, and thus can include molecular gas diffusion, gas adsorption and/or molecular dissociation on the surface of the electrode.

(2) n = 1/2: the corresponding reaction repeats twice per O₂, hence, involving only atomic species.

(3) n < 0.5: same as (2) but indirect pO₂ dependencies of involved defect species cannot be neglected either due to strong variations with pO₂ of vacancies/interstitials, electron/holes and/or the involvement of several different defects.

The n values for the two processes (P1 and P2) are given in Fig. 4 (d). The ASRₚ₁ increases with decreasing pO₂ with the value of n being 0.27 for P1. This value indicates that this process could be related to the charge-transfer reaction at the surface and incorporation of oxygen into the crystal lattice of the electrode, with several defect types being involved. P2 has an n value of 1.03, close to 1, which means that molecular oxygen would be involved in this step. The pO₂ dependence of the total ASR shows that the RDS changes depending on the pO₂ regime. At low pO₂ the reaction is most likely limited by the gas adsorption and dissociation process, while at high pO₂ it would be limited by the charge transfer and incorporation processes.

## 3.2. Fuel electrode-supported rSOC

In this study a commercial fuel electrode-supported cell was used with Ni-YSZ as fuel (hydrogen) electrode and YSZ as electrolyte. During cell fabrication and operation, high temperatures (> 700 °C) could lead to interdiffusion and reactivity of the oxygen electrode with the electrolyte. This can provoke the formation of secondary phases at the electrode/electrolyte interface which impede the ionic reaction and diffusion pathways and significantly decrease the performance of the cell. A Barrier Layer (BL) is typically employed to prevent the reaction between the YSZ electrolyte and lanthanum from the oxygen electrode (e.g., lanthanum-based cobaltites, manganites, and nickelates). This reaction can result in the formation of insulating interlayers such as
<figure><img src="image_6.png" /><figcaption>Fig. 3. (a) Nyquist plots from EIS at 600 °C in dry air fitted with an equivalent circuit, (b) DRT deconvolution of the EIS spectra at 600 °C, and (c) Arrhenius plot showing the ASR~p~ of the 500 nm and 1000 nm L2NO4 symmetrical cells.</figcaption></figure>

La₂Zr₂O₇ [47–49]. Morales *et al.* [50] showed a 70 % increase in performance using a 1.5 µm thick gadolinium-doped-ceria (CGO, Ce₀.8₀Gd₀.₂₀O₁.₉₅) barrier layer deposited by Pulesd Laser Deposition (PLD) between state-of-the-art oxygen electrode and electrolyte materials (i.e., La₁-xSrₓCo₁-yFeᵧO₃-δ and YSZ). The cell showed long term stability with no insulating interlayers forming after 14000 hours (h) of operation in SOFC mode [51]. Bernadet *et al.* [52] further optimized the thickness of the CGO BL to 200 nm with no significant degradation in performance or stability after 3500 h of operation in SOFC mode. Here, this same type of CGO BL with an optimized thickness of 150 nm was deposited on these cells by PLD.

On top of this BL a L2NO4 nano-columnar thin film with the optimized thickness of 1000 nm was deposited by PI-MOCVD. A schematic representation of the cell is shown in Fig. 5 (a). Fig. 5 (b) shows a porous surface microstructure by SEM for the L2NO4 film on CGO similar to the one directly on the YSZ single crystal substrates (see Fig. 1 (b)). The film shows large, free-standing grains (up to 200 nm) formed by a conglomeration of many small grains (10–20 nm), leading to a very large surface area. A cross-section SEM overview of the cell with the dense YSZ electrolyte and the porous Ni-YSZ support is shown in Fig. 5

(c). Despite the roughness of the electrolyte, the L2NO4 film shows a very high conformity and homogeneity in thickness. An in-depth microstructural analysis of the as-deposited film was carried out by cross-section TEM and SEM as shown in Fig. 5 (d) and Figure S1 (a-b), respectively. The cross-section SEM images indicate that the L2NO4 film grown on the CGO layer retains the highly porous nanostructure. The TEM image in Fig. 5 (d) shows a dense and homogenous layer of CGO on top of the YSZ electrolyte. The L2NO4/CGO and CGO/YSZ interfaces show a good adhesion between the different layers. The L2NO4 film is formed by a 10–20 nm dense bottom layer grown on the CGO BL, on top of which the dendritic nanostructure grows.

Fig. 6 shows the STEM EDX mapping of the different layers. There is a clear distinction in the distribution of the elements of each layer with no cationic inter-diffusion observed between the CGO and YSZ (cf. Ce and Zr signals) nor between CGO and L2NO4 (cf. La and Ce signals) during the growth of the L2NO4 thin film at 600 °C. The low MOCVD deposition temperature used presents a clear advantage compared to typical synthesis methods which require high sintering temperatures (> 1000 °C) and often leads to the formation of secondary phases [32,53].

The electrochemical performance of the cell in fuel cell mode was
<figure><img src="image_7.png" /><figcaption>Fig. 4. Analysis of the EIS spectra measured for the 1000 nm L2NO4 symmetrical cell at 600 °C (a) DRT deconvolution of the EIS spectra, and (b) specific capacitance of the two processes at HF (P1) and LF (P2), (c) Arrhenius plot of the ASR in dry air and (d) ASR as a function of $pO_2$ of different processes.</figcaption></figure>

assessed in the temperature range of 600–750 °C. The current density-voltage (j-V) and current density-power density (j-P) curves are shown in Fig. 7 (a). All measurements were carried out using a dry synthetic air flow on the oxygen electrode side and a dry hydrogen flow on the fuel electrode side. An open circuit voltage (OCV) in the range of 1.11–1.13 V was obtained for both cells, which indicates an adequate sealing and gas separation between the fuel electrode and oxygen electrode chambers. The cell shows high performance with the current density and power density reaching maximum values of 1.2 A cm⁻² and 1.2 W cm⁻², respectively, at 0.7 V and 670 °C, and of 1.3 W cm⁻² at 725 °C. The bending of the curves at high current densities could be due to a low availability of H₂ at the fuel electrode side. In electrolysis mode, the measurements were carried out under dry synthetic air flow on the oxygen electrode side and a mixture of 50/50 % H₂/H₂O and Ar on the fuel electrode side. An OCV in the range of 0.9–0.95 V was obtained. The cell shows high performance with the current density reaching –1.20 A cm⁻² at 670 °C and 1.3 V, as shown in Fig. 7 (b). This cell was tested in both SOFC and SOEC modes and remained at high temperatures (600–725 °C) for approximately 200 h. As shown in Fig. S2, the performance measured after 200 h showed a decrease in power density from 0.94 W cm⁻² to 0.84 W cm⁻² at 0.7 V and 615 °C. This degradation

in performance is believed to be mainly related to the microstructural densification of L2NO4 as it was deposited at 600 °C and later operated at much higher temperatures. Figure S1 (c-d) shows the cross-section SEM images of this cell after the measurements. Compared to the pristine cell the nanostructure is relatively denser although some porosity is still present. Additionally, the brush-painted gold current collector undergoes spheroidization over time, leading to a decrease in the measured effective area, which may create a false impression of decrease in activity. Fig. S3 shows the cross-section STEM-EDX images of this cell with no evidence of interlayer cationic diffusion or formation of an insulating phase, indicating the chemical stability of these thin films.

The performance in fuel cell mode of the cell with 500 nm L2NO4 is shown in Figure S4 (a). The power density is relatively lower compared to the cell with 1000 nm L2NO4 which is in agreement with the polarization resistance measurements of the symmetrical cells (see Fig. 3 (c)). A degradation test in SOFC galvanostatic mode was carried out at 650 °C for this cell at an operating current density of 0.5 A cm⁻² under dry hydrogen and synthetic air. Figure S4 (b) shows the change in voltage with time measured for approximately 500 h. It should be noted that the gap in data points is due to the limit of the data logger, however, the cell was continuously under operating conditions. Typically, the first
<figure><img src="image_8.png" /><figcaption>**Fig. 5.** (a) Schematic of the full rSOC configuration showing the Ni-YSZ fuel electrode support and functional layer, YSZ electrolyte, CGO BL and L2NO4 oxygen electrode; (b) Nano-columnar surface morphology of the L2NO4 film on the full rSOC cell observed by SEM; (c) Cross-section SEM image of the full rSOC cell, and (d) TEM image showing the L2NO4/CGO/YSZ interfaces.</figcaption></figure>

<figure><img src="image_9.png" /><figcaption>**Fig. 6.** STEM EDX elemental mapping of an as-deposited full cell with a 1000 nm L2NO4 oxygen electrode, 150 nm CGO BL and YSZ electrolyte.</figcaption></figure>
<figure><img src="image_10.png" /><figcaption>Fig. 7. Polarization curves of the fuel electrode-supported cell with 1000 nm L2NO4 in: (a) fuel cell mode, and (b) electrolysis mode between 600 and 725 °C.</figcaption></figure>

200–250 h are considered as the stabilization time where major changes
in microstructure and morphology, and interdiffusion of cationic species
can occur. After 250 h the cell showed notable stability with a small

degradation rate of 2.8 % $kh^{-1}$ (% of voltage loss per 1000 h). Current-Voltage (j-V) measurements were carried out periodically during the study. Figure S4 (c) shows the j-V and j-P curves measured at 650 °C

<figure><img src="image_11.png" /></figure>

<figure><img src="image_12.png" /><figcaption>Fig. 8. Comparison of the performance of the full cell with 1000 nm L2NO4 in this work in SOFC mode at 0.7 V and 700 °C: (a) per CRM mass, and (b) per active area of the oxygen electrode; and in SOEC mode at 1.3 V: (c) per CRM mass, and (d) per active area of the oxygen electrode with commercial and other high performance fuel electrode-supported cells [18,19,30,31,55,57,58].</figcaption></figure>

<figure><img src="image_13.png" /></figure>

<figure><img src="image_14.png" /></figure>
initially (black circles), after heating the sample to 750 °C (red circles), and then measuring at 650 °C for 150 h (blue circles) and for 500 h (green circles). Post-mortem cross-section SEM images, shown in Fig. S5, reveal densification of the nanostructure similar to the 1000 nm L2NO4. This change in morphology and the evolution of the current collector are believed to be related to the degradation in performance. However, this nanostructure is still very porous, hence the small decrease in performance after 500 h.

A comparison of the performance in SOFC mode of the 1000 nm L2NO4 cell with that of other fuel electrode-supported cells reported in literature is tabulated in Table S1 and plotted in Fig. 8 (a) and (b). Fig. 8 (a) compares the power density per CRM mass of the oxygen electrode and Fig. 8 (b) compares the absolute power density of the full cells at an operating voltage of 0.7 V and temperature of 700 °C (values interpo- lated for comparison). The commercial state-of-the-art cell from Sol- ydEra S.p.A. shows a power density of 830 mW cm⁻² and 180 mW mg⁻¹ of CRM [19]. A recent review (2023) by Morales-Zapata *et al.* [54] on the electrochemical performance of RP-phase oxygen electrodes revealed that Pr₄Ni₃O₁₀+δ [55] has the highest maximum power density reported till now for this family of materials. The reported cell with a 10 µm thick oxygen electrode gave a power density of 680 mW cm⁻² and 148 mW mg⁻¹ of CRM. The best performance of bulk full cells with only L2NO4 was reported by Lee *et al.* [31]. The cell was composed of a 20 µm L2NO4 layer and showed a maximum power density of 375 mW cm⁻² and 65 mW mg⁻¹ of CRM. Buzi *et al.* [56] reported a remarkable power density per CRM mass of 6400 mW mg⁻¹ of CRM using composite thin film of LSC-SDC. This performance is 1.5 times higher than the L2NO4 films used in this work, however, comparing the absolute power density, the performance of the L2NO4 thin films is almost 5 times higher.

Fig. 8 (c) and (d) show a comparison of the SOEC performance (in current per CRM mass) of the 1000 nm L2NO4 thin film electrode from this study with other high performance cells and absolute current den- sity, respectively, at an operating voltage of 1.3 V. Vibhu *et al.* [31] tested several RP-phase materials and a commercial cell with LSCF in SOEC mode. LSCF and L2NO4 showed a low performance with a current output of 130 mA mg⁻¹ and 940 mA cm⁻² while the best performing material was PNCO20 (Pr₂Ni₀.₈Co₀.₂O₄+δ) with 303 mA mg⁻¹ and 1320 mA cm⁻² under a mixture of 50/50 % H₂O/H₂ as the fuel and dry air as oxidant. Tan *et al.* [57] fabricated an impregnated LSCN-CGO (La₀.₈Sr₀.₂Co₀.₈Ni₀.₂O₃-δ - Ce₀.₉Gd₀.₁O₂-δ) oxygen electrode for rSOC. In SOEC mode, the current output reached 487 mA mg⁻¹ and 1600 mA cm⁻² under 90 vol% absolute humidity. The LSC-SDC thin film from Buzi *et al.* [38] again showed a remarkable current output per CRM mass of 12 A mg⁻¹ using composite thin film of LSC-SDC, however, the cur- rent per unit area was relatively lower than for the L2NO4 thin films. Shimada *et al.* [58] also reported a high performing composite electrode of SSC-SDC with 2.3 A cm⁻² but a much lower performance per CRM mass (750 mA mg⁻¹).

Despite having a much thinner oxygen electrode, the performance of the cell (per active area and per CRM mass) reported in this study is much higher than the ones previously reported in the literature for RP-phase materials. These promising results show the potential of using submicron thin electrodes, which use far less amount of CRM (~10–15 times) in the oxygen electrode, and show similar or even higher performance than state-of-the-art thick electrodes [5] enabling operating temperatures down to 650 °C, 100–150 °C lower than the current operating temperatures of 750 °C in commercial cells, and thereby minimizing thermally activated degradation processes, such as cationic interdiffusion.

## 4. Conclusions

Nano-engineered La₂NiO₄+δ thin films were optimized and initially evaluated in symmetrical cell configuration as potential oxygen elec- trodes for reversible solid oxide cells. At 600 °C the 1000 nm L2NO4 films showed a very low polarization resistance of 0.07 Ω cm⁻²

compared to 0.29 Ω cm⁻² of the 500 nm films. The rate limiting pro- cesses were analyzed by measuring the pO₂ dependence of the electro- chemical properties of the electrodes. This pO₂ dependence is in agreement with the rate limiting process being the oxygen incorporation and charge-transfer reaction at the surface of the electrode. The elec- trochemical performance of the optimized L2NO4 thin film was then evaluated in SOFC and SOEC modes using state-of-the-art commercial fuel electrode-supported cells (CGO/Ni-YSZ/YSZ). The use of thin films allowed to reduce the sintering temperature of the L2NO4 oxygen electrode by at least 300 °C compared to conventional sintering, thus avoiding secondary phase formation. The electrochemical performance of the cell significantly outperformed commercial state-of-the-art cells with a power density of 1.2 W cm⁻² at 0.7 V in SOFC mode and a current density of 1.2 A cm⁻² at 1.3 V in SOEC mode at 670 °C. Thus the exceptional performance of the L2NO4 oxygen electrode allows the cell to operate at least 100 °C lower than standard operating temperatures. These results demonstrate that nano-engineering of thin-film electrodes can achieve performance exceeding that of bulk commercial electrodes while using approximately 10–15 times less critical raw materials in the oxygen electrode, which is essential for the sustainable development of these electrochemical devices.

### CRediT authorship contribution statement

Adeel Riaz: Writing – review & editing, Writing – original draft, Methodology, Investigation, Formal analysis. Kosova Kreka: Writing – review & editing, Methodology, Investigation, Formal analysis. Alexander Stangl: Writing – review & editing, Supervision, Formal analysis. Fjorelo Buzi: Writing – review & editing, Methodology, Investigation. Silvère Panisset: Writing – review & editing, Investigation, Formal analysis. Federico Baiutti: Writing – review & editing, Supervision, Methodology, Conceptualization. Laetitia Rapenne: Writing – review & editing, Methodology, Investigation. Carmen Jiménez: Writing – review & editing, Methodology, Investigation. David Jauffres: Writing – review & editing, Supervision. Michel Mer- moux: Writing – review & editing, Supervision. Albert Tarancón: Writing – review & editing, Supervision, Funding acquisition, Concep- tualization. Mónica Burriel: Writing – review & editing, Supervision, Methodology, Funding acquisition, Conceptualization.

### Declaration of competing interest

The authors declare no competing financial interest. ■

### Acknowledgements

This project has received funding from the European Union's Horizon 2020 research and innovation program under grant agreements No. 824072 (Harvestore project) and No 101017709 (EPISTORE project). The authors acknowledge the facilities, and the scientific and technical assistance of the CMTC characterization platform of Grenoble INP supported by the Centre of Excellence of Multifunctional Architecture Materials "CEMAM" n°ANR-10-LABX-44-01 funded by the "Investments for the Future" Program. The authors acknowledge Jan Pieter Oueltjens and Dario Montinaro (SolydEra S.p.A., Italy) for the half-cells supply employed in this work. A.R. acknowledges the IDEX-Université Grenoble Alpes mobility grant for the research stay at IREC under which part of this work was carried out.

## Appendix A. Supplementary data

Supplementary data to this article can be found online at https://doi.org/10.1016/j.jpowsour.2025.237492.
# Data availability

All sample datasets and materials related to this work are made available under CC BY 4.0 license in the zenodo repository: https://doi.org/10.5281/zenodo.12737119.

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