A-site deficient chromite with in situ Ni exsolution as a fuel electrode for solid oxide cells (SOCs)


 A-site deficient chromite La0.65Sr0.3Cr0.85Ni0.15O3−δ (L65SCrN) decorated by in situ Ni exsolution was implemented as fuel electrode on 5 cm × 5 cm reversible electrolyte-supported solid oxide cells (rSOCs).


Introduction
Since the industrial revolution, carbon-rich fossil feedstocks have played an important role in our daily life in order to fulfil our needs for energy demand and for a broad range of household and commercial products. Nowadays, the chemical industry relies on crude oil, coal and natural gas to produce the key building blocks such as olefins and aromatics. Nevertheless, the improvement of the corresponding synthesis processes in terms of selectivity and energy consumption or the development of alternative routes have become a major priority for the modern industry due to limited recoverable natural oil reserves and growing environmental considerations regarding greenhouse gases emissions. 1 CO 2 is emitted in increasing amounts due to the growing need for power generation (coal-based plants) and industrial products, such as steel and chemicals, e.g. ethylene production by oxidative coupling of methane (OCM). 1 This greenhouse gas is also an essential feedstock for numerous chemical synthesis processes in combination with hydrogen. In some processes CO 2 is pre-reduced at high temperature with hydrogen through the reverse water gas shift (RWGS) reaction yielding CO -a more reactive molecule -as an essential building block for downstream chemical synthesis. Methanol, which is an important multipurpose intermediate commonly used for the production of various chemicals, is currently produced from syngas (H 2 + CO) which can also be generated via catalytic steam or autothermal reforming of methane. 1 In the light of syngas production and methanol synthesis, there are significant economic and environmental interests in valorizing renewable carbon sources. For this reason, since the last 10 years, the conversion of plant-derived materials (biomass) and CO 2 have attracted the attention for industry and academia with the aim to produce fuels and bulk chemicals by direct electrosynthesis routes with reduced CO 2 footprint. 2 Power-to-X concepts intend to convert excess renewable power into diverse fuels and chemicals that can be used for large capacity energy storage. 3,4 Among the various concepts, the technologies based on Solid Oxide Cell (SOC) operating at temperatures typically around 750 °C -850 °C enable conversion of electricity at high efficiency into valuable fuels (hydrogen or hydrocarbons) by means of high temperature electrolysis (HTE) without the need of precious catalysts. Interestingly, due to fast kinetics SOC enable the simultaneous electrolysis of H 2 O -CO 2 at high temperature into syngas, which can be further used for large-scale production of methanol and other green fuels and chemicals through the Fischer-Tropsch (F-T) synthesis. 1,5,6 Moreover, SOC offer the unique advantage to enable reversible operation. i.e. either energy storage or electricity production. Throughout the energy storage mode, electrical energy from renewable sources is converted to valuable fuels (hydrogen or hydrocarbons) by means of HTE, while on discharge mode these fuels could be used for power production through fuel cell operation. 4 A reversible Solid Oxide Cell (rSOC) system could effectively ensure large storage capacity and grid balancing.
State-of-the-art SOC rely on Ni-based cermet components, owing to the excellent electrical conductivity and high catalytic activity of Ni towards H 2 O -CO 2 splitting and hydrogen dissociation reaction at high temperatures. Ni-Zr 0.85 Y 0.15 O 2-δ (Ni -YSZ) cermet fuel electrodes -typically used in the so-called Anode-Supported Cells (ASC) -have been largely investigated in either operating modes. When operated in fuel cell mode, the electrodes are susceptible to poisoning with different fuel gas impurities such as sulfur species, that have deleterious effects on performance, especially in reformate gases, and long-term stability. 7,8 Moreover, they suffer from irreversible degradation when exposed to re-oxidation reactions. 9 These cermet electrodes are prone to Ni agglomeration leading to loss of electrical percolation and diminution of the triple-phase-boundary (TPB) length. 10 When operated in electrolysis, they suffer from irreversible microstructural alterations, especially at high temperatures, high current density and high pH 2 O. 11 In co-electrolysis operation, carbon formation has been observed at the electrode-electrolyte interface with a reactant conversion of ~67 % at 875°C, causing microstructural alterations accompanied with a deactivation of the active sites. 12 By contrast, Ni -Ce 1-x Gd x O 2-δ (Ni -CGO) based fuel electrodes -typically used in the Electrolyte-Supported Cells (ESC) -have been also investigated as fuel electrode materials because of their catalytic properties and CGO phase enhanced tolerance against carbon formation. 13 Nevertheless, due to the large content of metallic Ni, such Ni -CGO cermet electrodes are also vulnerable to dimensional alterations caused by grain coarsening upon redox cycling i.e. repeated alternate of oxidizing and reducing atmospheres, which adversely affects the apparent electronic conductivity leading to an increase of the ohmic resistance (R ohm ) and the gas transport properties of the electrode. 14 The large flexibility in use of SOC based electrochemical reactors, requires robust and durable fuel electrodes with high performance in either operating mode: fuel cell and electrolysis operation. This implies performance and durability in a broad range of partial pressures (pH 2 , pH 2 O, pO 2 , pCO and pCO 2 ) and a given dimensional stability nearly independent from the atmosphere.
Perovskite-based oxides (ABO 3 ) have been proposed as alternative materials to the Ni cermets as fuel electrodes for SOC because of their outstanding stability in both reducing and oxidizing atmospheres and their flexibility in terms of composition that enable a wide variety of doping elements on their Aand Bsites to tune their electrocatalytic properties. As a fuel electrode, high catalytic activity can be achieved when the A-site is a lanthanide and/or alkaline-earth cation and the B-site a transition metal cation, such as Mn, Co, Fe, Ni, Cr and Ti. 15 Strontium titanates have been widely studied and have shown remarkable performance as fuel electrodes in steam electrolysis at laboratory scale, where the perovskite's surface has been decorated with catalytically active Ni and Fe nanoparticles. 16 It has been reported that catalytically active nanoparticles surface decoration can be achieved by redox exsolution methods, where a catalytically active metal (i.e. Ni or Fe) is incorporated into the crystal lattice of the perovskite backbone in oxidizing conditions and is released (exsolved) on the surface as metal nanoparticles, either by exposure to a reducing atmosphere or by applying a large cathodic overpotential. 17,18 It is generally admitted that exsolution is favoured upon A-site deficiency: when the oxygen vacancy concentration is high enough to partially destabilize the perovskite lattice due to the high deficiency on Aand O-sites, metal particles from the B-site exsolve while charge balance of the lattice is maintained. 19 A recent study by Neagu et al. about Ni exsolution on lanthanum-calcium doped titanates and lanthanum-cerium doped titanates by in situ observation with environmental transmission microscopy (ETEM), showed that the exsolution phenomena and thus the shape of the resulting nanoparticle are significantly affected by the temperature and the oxygen partial pressure (pO 2 ), 20 being important operating parameters for the rSOC reactors.
Lanthanum chromites present an alternative towards strontium titanates as another perovskite family that can also host B cations to be exsolved in situ on their surface to enhance the electrocatalytic activity. (La,Sr)(Cr,M)O 3 perovskites (M= Mn, Fe, Co and Ni) have been recently investigated for H 2 O electrolysis, CO 2 electrolysis and H 2 O -CO 2 co-electrolysis: mostly in stoichiometric formulations 21 and a few with A-site deficiency. 17,22 However, the Ni exsolution phenomena on lanthanum chromites upon temperature and atmosphere variation remain unclear, and the performance of such perovskite electrodes still needs to be improved in order to achieve comparable results with the typical Ni-cermet fuel electrodes.
Given the operating conditions of rSOC reactors with focus on Solid Oxide Electrolysis Cells (SOEC) applications, the lack of Ni exsolution research on chromites arouses the interest to investigate the performance of Ni-decorated chromites as fuel electrodes for SOC, raising as well the importance to evaluate their durability and performance in either mode on rSOC reactors.
In this paper we focus on the exploration of the A-site deficient chromite La 0. 65  O (98.5% Alfa Aesar) were dissolved in deionized water and mixed with glycine (J.T.Baker™). The glycine molar ratio for the total content of metal cations was 2:1. Next, these solutions were stirred and heated on a hot plate until a dark greencolored gel was formed. Previous thermogravimetric measurements in synthetic air performed on these gels indicated that the solvent evaporation takes place at ~91 °C followed by an exothermic selfcombustion reaction at ~220 °C. 24 Therefore, in this study, the gels were heated up to ~220 °C where self-combustion occurred. Finally, the resulting ceramic precursors were calcined in air at a rate of 3 °C /min up to 1400 °C for one hour since it was the minimal firing temperature at which a perovskite phase could be achieved, 24 which is consistent with previous studies on lanthanum chromites. 25

Characterization of LSCrN powders
Crystalline structure was investigated by X -ray diffraction (XRD) with a RIGAKU diffractometer operating at 40 kV and 30 mA with Cu-Kα 1,2 radiation source and Bragg-Brentano configuration in the range of 2θ from 20° -80° with a scanning rate of 0.4°/min. Crystalline phases were identified with the ICDD database. A different scanning rate of 0.1°/min was used for the as-prepared and reduced L65SCrN samples where phases were identified and quantified by Rietveld analyses using the FullProf.2k program suite.
Morphology and microstructure were observed with a scanning electron microscope Zeiss ULTRA PLUS SEM (Carl Zeiss AG, Germany) in combination with energy-dispersive X-ray spectroscopy (EDX) for elemental analysis, where a Bruker XFlash 5010 detector was operated at 125 eV with the Quantax 400 Software. The spatial resolution was ~ 100 nm and elements with atomic numbers higher than 4 (Boron) could be detected.
Surface chemistry investigations with X-ray photoemission spectroscopy (XPS) were carried out using in a system with a base pressure of 2x10 -10 mbar, with a hemispherical analyzer (ESCALAB250, ThermoFisher Scientific) and a monochromated Al Kα source with an X-ray energy of 1486.74 eV (XM1000, ScientaOmicron). The peak shape analysis was carried out with Unifit 2013, applying convoluted Gaussian/Lorentzian profiles and a Shirley background function. 26,27,28 The surface stoichiometry of the occurring atoms/signals was calculated using the numerically fitted peak areas, photoionization cross sections by Yeh and Lindau 29 and instrumental transmission functions given by the manufacturer.
The reducibility of the as-prepared L65SCrN and L70SCrN powders was characterized by means of thermogravimetric analysis (TGA) in reducing atmosphere (5% H 2 -Ar) with the analyzer Netzsch Jupiter 449C at a heating rate of 3 °C /min from 25 °C to 1200 °C. Such reducibility analysis by TGA was accompanied with a temperatureprogrammed reduction (TPR) performed on the flow-through quartz reactor TPDRO 1100 (Thermo Scientific, Italy), in which L65SCrN and L70SCrN powder specimens were introduced using quartz glass wool as a support. A thermocouple (type K) was placed in a thin quartz glass tube next to the specimen to monitor the temperature. The oven temperature was monitored and controlled by another thermocouple. The specimens were pre-treated in the reactor tube by flowing Ar gas at a flow rate of 20 mL/min and by heating at a rate of 10 °C/min from 30 °C to 150 °C with a holding time of 60 min. After cooling the specimen back to 30 °C, the TPR process was initiated with a reducing gas mixture of 5 % H 2 -Ar at a constant flow rate of 20 mL/min and a heating rate of 5 °C/min from 30 °C to 1000 °C with a holding time of 60 min at 1000 °C. At these conditions the sample temperature was ~ 1000°C while the oven temperature was 1100 °C. The exhaust gas from the reactor was analyzed by the Thermal Conductivity Detector (TCD).

Cell manufacturing
The L65SCrN electrocatalyst was implemented as fuel electrode into an electrolyte-supported cell (ESC) by screen printing, using a commercial square substrate (50 mm x 50 mm and 90 μm of thickness) of 3 mol% Y 2 O 3 -doped ZrO 2 electrolytes double-side coated with ca. 5µm of Ce 0.8 Gd 0.2 O 2-δ (CGO20-3YSZ-CGO20) from Kerafol GmbH, Germany. The fuel electrode ink was prepared by dispersing the L65SCrN powder in a solution (94 wt% α-Terpineol and 6 wt% ethyl cellulose) with a powder to solution ratio of 2:1, followed by mixing with the 3-roll milling machine EXAKT 80E EL. The prepared ink was printed on the electrolyte using the screen printer Aurel model 900 (Aurel automation s.p.a, Italy). The half-cell was fired at 1200 °C for 1 hour in air with a heating rate of 3 °C/min. Afterwards the air electrode was printed on the other half of the cell with a commercial ink of La 0.58 Sr 0.4 Fe 0.8 Co 0.2 O 3-δ (LSCF). The printed area for both electrodes was 16 cm 2 (40 mm x 40 mm). Platinum paste was brushed on the sintered fuel electrode surface for current collection. Finally, the cell was fired at a rate of 3 °C/min to 1050 °C in air and held for one hour.

Electrochemical characterization
The electrochemical performance of the L65SCrN fuel electrode in ESC architecture was studied on the test bench described elsewhere. 24 The fuel electrode was contacted with a platinum mesh and the oxygen electrode with a gold mesh. A gold frame was used as sealant between fuel and air side. For commissioning, the cells were heated (3 °C/min) to 900 °C with N 2 (1 SLPM) and air (1 SLPM) for sealing purposes and subsequently reduced with H 2 (1 SLPM) for 1 hour on the fuel side. Afterwards, the operating temperature was adjusted to 860 °C. Electrochemical experiments were carried out in different fuel gas mixtures shown in Table 1. The equilibrium gas phase compositions and the theoretical OCV according to the Nernst voltage were calculated with the software CANTERA. 30 Electrochemical Impedance Spectroscopy (EIS) was performed in galvanostatic mode with the workstation Zahner PP-240 at a frequency range from 50 mHz to 100 kHz. The amplitude of the current stimulus was 500 mA. Distribution of relaxation times (DRT) calculations were carried out with the impedance analysis and modelling software ec-idea 31 and the equivalent circuit model-fit of the impedance data with the commercially available program ZView®. 32 In fuel cell (FC) mode the polarization curves (i-V) were measured from OCV to 0.8 A•cm -2 at a rate of 0.012A •s -1 and in electrolysis (EC) and co-electrolysis (co-EC) modes from OCV to -1.0 A•cm -2 at a rate of -0.012A•s -1 . For all operating modes the total fuel gas flow was maintained at 1 SLPM, except for co-EC which was kept at 0.8 SLPM.

Results and discussion
A comparative assessment towards LSCrN reducibility and Ni exsolution was performed between stoichiometric L70SCrN and Asite deficient L65SCrN perovskite powder samples. Furthermore, the temperature effect on the Ni exsolution was investigated on the L65SCrN and the electrochemical performance was evaluated in a full cell assembly with ESC architecture, in which the L65SCrN perovskite was implemented as fuel electrode.

Nickel exsolution assessment on LSCrN powders
Phase identifications of both ceramic powders L70SCrN and L65SCrN as-prepared and reduced (with 5% H 2 -Ar at 3 °C /min from 25 °C to 1200 °C) were performed by XRD in order to verify that the perovskite phase was stable after the reduction treatment. X-ray diffractograms of the L70SCrN before and after reduction are shown in Fig. 1. No NiO secondary phase or other impurities could be identified. Interestingly, no metallic Ni could be detected on the reduced L70SCrN sample as one may have expected upon reduction. However, a few nanoparticles could be observed on the SEM image on the reduced L70SCrN powder in the electronic supplementary information (ESI) in Fig. S1. This suggests that the total amount of metallic Ni in the reduced L70SCrN (originating from the reduction of a possible NiO secondary phase as well as from the likely to occur exsolution of metallic Ni upon reduction of the host perovskite) remained below the detection level of the XRD analysis, i.e. a phase content of less than 1 wt %.
On the contrary, secondary phases could be identified for L65SCrN on both as-prepared and reduced samples ( Fig. 2a and Fig. 2b). Therefore, Rietveld analyses were performed with the aim to quantify those secondary phases. An orthorhombic lattice (space group 62, Laue class mmm) was calculated for L65SCrN as-prepared sample with parameters to be a = 5.496 Å, b = 5.450 Å, c = 7.737 Å and V = 231.762 Å 3 . A secondary phase was identified as Nickel oxide (NiO) with cubic lattice (space group 225, Laue class m-3m) and lattice parameters a = 4.176 Å and V = 72.851 Å 3 . Although the NiO content was 4.22 mol %, it was not considered to be a detrimental impurity since it would be reduced operando into metallic Ni, being also catalytically active and electronically conductive. For the reduced L65SCrN sample, an expanded orthorhombic lattice (also space group 62, Laue class mmm) was calculated (a = 5.465 Å, b = 7.758 Å, c = 5.506 Å and V = 233.409 Å 3 ). A secondary phase of metallic Ni was calculated to be 4.83 mol % with cubic lattice (space group 225, Laue class m-3m) with a = 3.524 Å and V = 43.761 Å 3 , from which the metallic Ni characteristic peak (111) was identified at 44.4°. This corroborates that metallic Ni can be achieved upon exposure of L65SCrN to a reducing atmosphere at high temperatures. It is pertinent to note that the amount of metallic Ni after reduction (4.83 mol %) is greater than the content of NiO (4.22 mol %) on the as-prepared sample, indicating that the metallic Ni has been effectively exsolved from the L65SCrN matrix in these reducing conditions. The characteristic peak of the perovskite at 47.2° shifted slightly to lower diffraction angles (46.9°), which corresponds to an expansion of the perovskite orthorhombic lattice (from 231.762 Å 3 to 233.409 Å 3 ). This expansion could be due to the loss of Ni 2+ cations from the lattice that are reduced to Ni 0 and exsolved at the surface in correlation with the consumption of vacancies on the A-site, but also due to the oxygen loss. 33,34 A change in the Cr 4+ / Cr 3+ ratio in the host perovskite matrix could also contribute to this expansion, taking into account that the ionic radius in an octahedral environment for Cr 4+ is 0.55 Å and for Cr 3+ is 0.615 Å. 35 To better assess this, the L65SCrN as-prepared and reduced powders were also investigated by means of XPS.  At first, the XPS studies could not conclusively confirm the presence of a metallic Nickel phase in the reduced L65SCrN sample. As the typically used signal of the Ni2p electrons is overlapping with the very distinct 3d signals of Lanthanum of unusual quadruplet shape, the deconvolution of the traces of metallic Nickel was not possible this way. Similar to Nenning and Fleig, 25 the Ni3p region was used ( Fig.  3a), which is also partially overlapping with the Cr3s signal, but the chemical structure of the surface Chromium could be evaluated by means of the Cr2p signal ( Fig. 3b) and added as boundary to the numerical model for the Nickel 3p region, with its low cross section. While it was possible to identify two occurring Ni species in both the as-prepared and reduced L65SCrN samples, which are attributed to oxidic Ni 2+ (~67.2 eV) and surface Ni(OH) 2 (~69.6 eV), 36 the clear evidence for a Ni 0 species is lost in the signal noise. However, the signal deconvolution, which was performed using the aforementioned Chromium signature and a fixed Lorentzian peak width for the as-prepared sample, converged to a result with Gaussian line widths of 2.2 eV (Ni 2+ ) and 2.6 eV (Ni(OH) 2 ), which is corresponding to the overall findings of broader signals for the latter. For the reduced sample, it converged to a Gaussian linewidths of 3.0 eV (Ni 2+ ) and 2.2 eV (Ni(OH) 2 ) -which has to be considered a false solution, due to the resulting linewidths. The deconvolution with three components, Ni 2+ , Ni(OH) 2 , and Ni 0 , however could not be calculated successfully. The accuracy of the peak fit, which is relying on the data quality, may be argued, but the indirect indicator supports the finding from the XRD evaluation.
It has to be noted that the high surface sensitivity of this method might misguide the interpretation if not considered. The information depth, i.e. 3·, where  is the escape depth of the relevant electrons, is below 10 nm, 37 which means that only the topmost surface of the crystallites is visible to this method. This surface sensitivity may also explain the presence of Ni(OH) 2 , that can easily form on top of Nickel exsolved surface particles. 38 The Nickel ratio was calculated with the Ni3p signal for the asprepared and reduced samples. The concentration in Nickel at the surface was 11.2 at % for the as-prepared sample and 3.8 at % for the reduced sample (Table 2). These results, which may seem surprising at first, could be explained by an agglomeration of surface Nickel in nanoparticles upon reduction. Since their typical particle size (Fig. 3d) is larger than the information depth that is less than 10 nm, they may not be fully probed by XPS. This is an important finding because this suggests that the arrangement of the Nickel at the surface of the L65SCrN sample changed from a distribution, one can expect homogeneous at the surface of the as-prepared perovskite, to a more heterogeneous distribution where the Nickel is agglomerated and likely concentrated into nanoparticles during reduction. This indicates that the top surface of the perovskite after reduction tends to be depleted in Nickel, yielding overall a reduction of the Nickel concentration at the surface of the material analyzed by XPS.
The ratio between the surface states of Chromium was determined by a peak fit of the Cr3p region. The separated components were identified, according to systematic studies by Biesinger et al., 39 to Cr 3+ (oxide) at ~575.7 eV, Cr 3+ (hydroxide) at ~577.6 eV, and Cr 6+ at 579.5 eV. A Cr 4+ state, as a possible cause for the observed change in lattice parameters has been discussed in earlier works on LSCr perovskites 40 , but cannot be distinguished with this method due to the very close binding energies of Cr 3+ and the anomalous Cr 4+ . 41,42 Considering the inherent surface sensitivity of photoemission spectroscopy, the bulk properties of the investigated perovskite crystallites are not accessible anyway, and the surface states of Chromium are not relevant to the lattice parameters. However, the surface stoichiometry (Table 2) of these Chromium species can give indirect insight. The reduced L65SCrN sample shows an abundance of surface Cr 3+ (both oxidic and hydroxide), which would be consistent with the mentioned lattice expansion observed by XRD (Fig. 2). In order to better understand and highlight the role of the A-site deficiency on the Ni exsolution, TGA and TPR were performed in reducing atmosphere (5% H 2 -Ar) on the as-prepared L70SCrN and L65SCrN powders. Since no other volatile species or compounds are expected to be formed during such thermal treatments, the net weight loss measurement by TGA is attributed to the net loss of oxygen, assuming that the oxygen from the perovskite lattice (and from the NiO secondary phase) reacted with the hydrogen to form H 2 O. 34 The TGA and TPR profiles are shown on Fig. 4a and Fig. 4b, respectively. In Fig. 4a, the maximum onset for weight loss, corresponding to the maximum value of the derivate (DTG Max ), occurs at 442 °C for L65SCrN and at 472 °C for L70SCrN.  quantity of perovskite moles is conserved, it is possible to assume the following relation: Where and are the molecular weights of the 3 3as-prepared and reduced perovskite samples respectively. This expression can also be written in terms of the atomic weight of oxygen as: Knowing that the weight loss measured by TGA is given by: eqn (2) For the case of L65SCrN and L70SCrN, i.e. with Ni cations on the Bsite that can be reduced and forced out of the lattice, the reduction and exsolution process of the Nickel will result in a consumption of the oxygen vacancies, so that it is not possible to relate directly net oxygen mass loss with oxygen deficiency. Nonetheless, though the initial oxygen stoichiometry was not determined in each of the compounds, a net specific oxygen consumption δ O can be calculated by analogy with eqn (4). Therefore, by considering the net weight losses of 1.94 wt% and 2.68 wt%, the net oxygen consumption for each perovskite was calculated to be δ O (L70SCrN and L65SCrN, respectively. The significant difference in the net specific oxygen consumption upon reduction between L65SCrN and L70SCrN suggests that the A-site deficiency enhances the reducibility of the LSCrN, which seems to favor the exsolution of metallic Ni. To better understand the influence of the temperature and time on the formation of metallic Nickel, the Ni nanoparticles' morphology was investigated in two isothermal reducing treatments: below and above 600 °C. These treatments were performed on the L65SCrN since it showed superior reducibility and are explained and detailed in the following section.

Influence of the temperature on the Ni exsolution on L65SCrN
L65SCrN powder samples were exposed to pure hydrogen at different temperatures. For comparison, powder samples were annealed in hydrogen at either 500 °C or 900 °C for an annealing time of 3 hours. After reduction the presence of metallic Ni concomitant with the perovskite phase was confirmed in the XRD patterns (Fig. 5). However, the presence of impurity traces such as SrO 2 and La 2 O 3 could also be noted for the reduced samples and traces of Sr 3 (CrO 4 ) 2 for the as-prepared sample. SEM imaging of the two powder samples, i.e. L65SCrN reduced at 500 °C for 3 hours and L65SCrN reduced at 900 °C for 3 hours revealed different morphologies (Fig.  6). On the L65SCrN sample reduced at 500 °C for 3 hours one can observe spherical well-dispersed nanoparticles on the perovskite surface of a diameter of ~ 8 nm (Fig. 6a, surface type 1) to ~ 30 nm (Fig. 6a, surface type 2). The presence of the nanoparticles is correlated with a local enrichment in Nickel which suggests that those nanoparticles are very likely made of metallic Nickel. These observations are detailed in the ESI, in figures S2 and S3. The corresponding EDS mapping of Ni and Cr are shown in figures S4 and S5 respectively. Interestingly, it could be suggested that the exsolution of Ni nanoparticles at the surface of the perovskite grains may depend on the crystallographic orientation: surfaces denoted by type 1 appear to have qualitatively higher Ni nanoparticle density than type 2 surfaces (Fig. 6a). Though it was not possible to determine the specific crystallographic orientations of those surfaces, this strongly suggests that exsolution of nanoparticles is influenced by the surface characteristics of the perovskite grains. This is in agreement with the observations made on titanates by Neagu et al. as they found that during exsolution the particles remained socketed in the [110] crystallographic orientation with respect to the perovskite lattice, which is one of the key structural features that grants to exsolved nanoparticles their stability. 20 Such orientation relationship is in accordance with previous reports, whereby the diffusion direction for B-site cations in perovskite lattices is along the [110] orientation. 19,20,45 On the L65SCrN sample reduced at 900 °C for 3 hours coarser Ni nanoparticles of irregular shape of ~30 nm up to 100 nm could be observed (Fig. 6b). Such nanoparticles growth is consistent with the observations made on the corresponding XRD pattern (red pattern on Fig. 5) that reveals a more intense Ni characteristic peak. A closer inspection of the different surfaces could not reveal a variation in nanoparticle density, which appears qualitatively lower than the one on the sample reduced at 500 °C. This lower particle density at the surface of the perovskite and their coarser particle size suggest a growth mechanism of the Ni nanoparticles that takes place upon temperature increase.   48 They observed that the Ni nanoparticles had coarsened significantly from a particle size of ~ 10-15 nm (after 3 hours) to an average hemisphere diameter of 50-60 nm after 311 hours of reduction. In contrast, for the Ru-doped chromite after 311 hours of reduction at 800 °C there was no significant change on the Ru nanoparticles since their size did not exceed 10 nm. From these observations and by analogy to thinfilm nucleation they concluded that particle coarsening may be due to a fast surface diffusion, which would allow nuclei to be fed by adatoms yielding larger and more widely spaced nuclei. 48 Therefore, they suggested that the faster Ni particles coarsening was likely explained by larger Ni surface diffusivities in comparison to Ru on the chromite, although quantitative data on chromites are not available. 48 This would be consistent with the above-mentioned observations made on Nickel concentration determined by XPS at the surface of the as-prepared and reduced L65SCrN samples. Such high surface diffusivity was highlighted by Sakai et al., who investigated the chromium diffusion in lanthanum chromites between ~700 °C -1400 °C by 50 Cr tracer diffusion and secondary ion mass spectrometry (SIMS). 49 They estimated that independently of the temperature the grain boundary diffusion coefficient was 10 5 times larger than the bulk diffusion coefficient. 49 Interestingly, this behavior has also been observed in other perovskites families, such as strontium titanates SrTi 0.75 Co 0.25 O 3-δ , where exsolved Co particles diffuse onto the existing Co nanoparticles rather than nucleating in new locations at the grain boundaries, due to the increment of Co diffusivity at high temperatures (above 700 °C). In this case the distances between the particles previously nucleated are assumed to be shortened. 46 Another interesting observation was made by Kousi et al. on La 0.7 Ce 0.1 Co 0.3 Ni 0.1 Ti 0.6 O 3-δ (LCCNT) where they identified different sizes on the exsolved Ni-Co nanoparticles in the bulk as compared to the ones exsolved on the surface: the bulk particles were smaller (~10 nm) than the ones exsolved at the surface ~40 nm, noting as well that on the bulk the nanoparticle population was significantly higher. These conjectures were made based on a SEM cross-section evaluation, where it is possible to identify the surface and the bulk. 50 Moreover, it is possible to correlate the difference on the observed Ni exsolution morphologies at 500 °C and 900 °C upon reduction ( Fig.  6) with the formation of oxygen vacancies, since the exsolution is favored when these vacancies reach a high concentration. 19 The formation of oxygen vacancies upon reduction may take place either in a surface site or in a bulk site. 51 It has been claimed that there is a strong correlation between the bulk and the surface kinetics, which indicates that not only the oxygen vacancy concentration but also the pO 2 plays a significant role in the surface exchange processes in mixed ionic and electronic conductors (MIEC) such as chromites. 52 Another approach was made by Gao et al. who investigated Ni exsolution phenomena on a Sc-based A-site deficient perovskite La 0.4 Sr 0.4 Sc 0.9 Ni 0.1 O 3-δ (LSSN) at different temperatures and annealing times. They proposed that Ni exsolution could be considered as a chemically driven heterogeneous phase transformation, being a consequence of four physical processes: diffusion, reduction, nucleation and growth. They found that the nucleation is affected by: mechanical stresses, related strains on the perovskite lattice, metallic Ni wetting angles, and A-site and oxygen vacancies. These factors significantly determine where the nucleation would take place. Moreover, parameters such as the atmosphere (e.g. pO 2 , pH 2 ), annealing time and temperature may affect the particle growth. 53 In the case of A-site dopants diffusion, a study made on manganitebased perovskite oxides, LnMnO 3 , demonstrated that the surface oxygen vacancy attracts the dopant (that partially substituted the host on the A-site) driving it to the surface on the host sublattice. 33 It has been assessed by DFT (density functional theory) that the elastic and electrostatic interactions of the dopant with the surrounding lattice are driving forces for dopant segregation on perovskite compounds. The factors that may affect these driving forces are the dopant size, the lattice parameter, and the distribution of charged vacancies. 33 However, the diffusion phenomenon from the segregating cations should be carefully studied since the surface composition depends both on thermodynamics and kinetics, 33 which could be also the case for the B-site dopants diffusion in other perovskite systems.
Regarding the L65SCrN in the present study, it is unclear how the Ni enrichment at the surface originating from the bulk may occur. If we consider the assumption by Gao et al. that the mass transport process during the Ni exsolution is critical for the particle growth, it is likely that such growth results from the Ni 2+ ions diffusion followed by the reduction reaction to metallic Ni, which may be limited by two possible models: mechanical energy effects (as a function of the strain activation energy) or limited Ni supply. 53 They found that the Ni particles preferably nucleate at the surface rather than in the bulk due to tendency to decrease the strain activation energy. 53 By DFT calculations they concluded that these models are likely to represent the actual Ni exsolution mechanism on LSSN. This contrasts with the observations reported on lanthanum titanates based fuel electrodes for which the particle -substrate interaction prevails and thus stabilizes the Ni exsolved nanoparticles on the surface in the temperature range between ~ 650 °C -900 °C in reducing atmospheres. 20 3 . They found that the location of the particlesocket did not change during the timescale of the experiment, indicating that the nanoparticle was locked/socketed in place once it was formed (exsolved). 20 More precisely, they found during the experiment that additional particles formed within nanoscale proximity of the first ones but neither of them moved nor drifted under the environmental transmission microscope (ETEM) electron beam. This revealed that particle-support interactions are strongly dominant over particle-particle interactions for titanates. 20 Contrary to the case of lanthanum chromites, particle-particle interaction prevails at high temperatures.
Although lanthanum titanates do not seem to exhibit the same behavior as strontium titanates, lanthanum scandates or lanthanum chromites towards Ni particle growth, recent studies highlighted that for perovskite oxide-based electrocatalysts, there is a strong influence of the gas atmosphere, i.e. pO 2 in correlation with the temperature, on the shape of the exsolved nanoparticles, 20,25 which is susceptible to affect their electrocatalytic performance.
In this study, the reduction of L65SCrN at low temperatures (~500 °C) yielded the finest and well-dispersed Ni exsolved nanoparticles. Such operating conditions are still far below from the typical operating temperatures (T > 800 °C) of ESC in SOC applications. Since it is shown that Ni particle-particle interactions prevail at high temperature in chromites, yielding particle coarsening, it is an important aspect to consider for the implementation of the L65SCrN as fuel electrode into a SOC stack because it may affect the morphological stability of the reactive surfaces. Possibly, the Ni exsolved particle size could be optimized with a rigorous investigation on the Ni exsolution phenomena in these A-site deficient chromites. Synthesis and processing parameters may also play a significant role: porous structures may be one key factor for Ni exsolution. 53 For a better understanding, additional investigation of the Ni concentration profiles across the perovskite grains upon reduction as well as the Ni particle size evolution as a function of temperature, annealing time, gas atmosphere, grain size and porosity would be necessary, e.g. by TEM and TOF-SIMS, accompanied with DFT modeling. 33,46,53 Therefore, it is questionable how stable a L65SCrN electrocatalyst may perform at the stack level since pO 2 gradients are usually observed along the gas channels: for instance, high pH 2 at inlet and high pH 2 O at outlet are characteristic of FC operation, while on EC mode high pH 2 O at inlet and high pH 2 at outlet are typical (assuming high fuel and steam utilizations for both modes). Moreover, reversible operation (FC-EC) would expose the electrocatalyst alternatively to pO 2 gradients.
In the following section we focus on the electrochemical performance of cells with L65SCrN fuel electrode upon variation of the operating conditions.

SOC electrochemical performance with L65SCrN electrocatalyst as fuel electrode
The electrochemical performance of a full cell with ESC architecture and a L65SCrN fuel electrode has been evaluated in FC, FC-EC, EC and co-EC modes as described in Table 1.

Performance evaluation in FC, FC-EC and EC operation
The measured open circuit voltage (OCV) was 1.25 V at 900 °C with pure H 2 /air, demonstrating appropriate gas tightness of the sealing. For the three different operating modes the measured OCV was slightly above (~ 20 mV) the theoretical Nernst potential E 54 : Such difference was assigned to a deviation of the inlet gas composition due to inaccuracy of the steam supply mass flow control.
Polarization curves at 860 °C in fuel cell (FC), reversible (FC-EC) and electrolysis (EC) operation are shown in Fig. 7a. In FC mode, i.e. with a 90 % H 2 -10 % H 2 O fuel gas mixture the oscillations observed find their origin in pH 2 O fluctuations from the steam supply due to marginal operation. In reversible mode FC-EC, i.e. with a 50 % H 2 -50 % H 2 O fuel gas mixture, the I-V characteristic evolves continuously from either side of the OCV which reflects the reversible functionality of the full cell and thus the L65SCrN electrocatalyst in either mode.
In EC mode, with a 20 % H 2 -80 % H 2 O fuel gas mixture, the I-V characteristic shows a nearly linear evolution from OCV until an inflection point which corresponds to the thermoneutral voltage of steam electrolysis (~1.29 V). 3 Above this value, the slope of the polarization curve decreases yielding a curve flattening. This is explained by the exothermal nature of the steam electrolysis at higher cell voltage, causing a net heat production. Since the oven of the test bench was operated isothermally, this heating effect cannot be controlled at the cell level causing a net increment of the temperature and enhancing the electrode reactions kinetics. 3 EIS data recorded near OCV conditions are shown on the Nyquist Plot of Fig. 7b. The ohmic resistance R ohm is estimated at 0.47 Ω•cm 2 with 10 % H 2 O, 0.46 Ω•cm 2 with 50 % H 2 O, and 0.45 Ω•cm 2 with 80 % H 2 O, suggesting a sensitivity of this parameter to the pH 2 and thus the pO 2 in the feed gas. As p-type conductor, this increment of R ohm upon increase of pH 2 could be explained by a decrease of the conductivity in the L65SCrN, since the positively charged oxygen vacancies that are created upon reduction hinder the transportation of electrical holes, 55 and decrease the effect of the alkaline earth doping on the electrical conductivity of lanthanum chromites. 56 However, it is important to note that the ohmic resistance values are comparable with commercial ESC references, since they are slightly lower than for a state-of-the-art Ni-CGO fuel electrode tested also at 860 °C with an estimated value of 0.55 Ω•cm 2 . 3 At low frequencies, i.e. below 1 Hz, EIS data were scattered due to the small voltage variations induced by fluctuations in the steam supply, what made difficult to determine accurately the total area specific resistance from these spectra. Therefore, the total area specific resistance (ASR DC_Total ) was calculated as the slope from the polarization curves at ± 0.3 A•cm -2 (linear range) where the influence of gas conversion and concentration polarization are expected to be minimal. The ASR DC_Total at -0.3 A•cm -2 was calculated from the polarization curve (Fig. 8a) to be 0.676 Ω•cm 2 . This value lies in the same order of magnitude (being lower) as the one reported for an ESC reference with a Ni-CGO fuel electrode tested in co-electrolysis (25 % H 2 , 25 % H 2 O, 25 % CO 2 , 25 % CO) at 830 °C which showed an ASR DC_Total of 0.84 Ω•cm 2 at -0.3 A•cm -2 . 58 Even though these conditions are not directly comparable this first approach suggests promising performance of the L65SCrN electrocatalyst as fuel electrode. The polarization curve shows a linear evolution from OCV until the thermoneutral voltage of coelectrolysis that is determined in the tested conditions to ~1.32 V lying between the steam electrolysis value (1.29 V) and CO 2 electrolysis (1.46 V). 3,59 Above this point, the curve flattens due to the decrement on the ASR DC_Total, probably enhanced by the temperature increment at this exothermal regime. 3 From the EIS spectra recorded at OCV (Fig. 8c and 8d) a R ohm of 0.47 Ω•cm 2 was identified for the L65SCrN (blue pattern). Furthermore, the polarization resistance ASR AC_Pol was estimated to be 0.29 Ω•cm 2 , which can be directly compared with the estimation by Dueñas et al.
on an ESC with Ni-CGO fuel electrode under the same operating conditions (green pattern): 0.23 Ω•cm 2 . 3 Regarding the electrode polarization processes, they are dominated by one main contribution which suggests a convolution of the electrode losses with the gas losses (conversion and diffusion processes).
A calculation of the distribution of relaxation times (DRT) was performed with the aim to identify relevant processes within the L65SCrN fuel electrode-cell. Five contributions or processes could be identified (Fig. 8b), which allowed to propose the equivalent circuit model (ECM) shown on Fig. 8e, with good fitting of the EIS spectra on Fig. 8c and 8d (χ 2 = 1.55×10 -4 ). The proposed ECM comprised an ohmic resistance R o of 4.69 ×10 -1 Ω•cm 2 to model the ohmic losses and a series connection of five RQ-elements, where Q represents a constant-phase element (CPE).
The different five physical and electrochemical processes were identified as shown in Table 3, where the resulting peak frequency (f DRT_peak ) and the polarization resistance values (area under the peaks) from the DRT analysis were used to calculate the capacitance of an ideal RC element from the general eqn (6): 60 for n = 1. The resulting R and Q (which has the characteristic of a capacitor C for n=1) were used as initial parameters for the equivalent circuit model fitting in ZView for the five RQ-elements. The fitting parameters including R, Q and n (for n between 0 and 1) were then iterated until the new f ECM_peak matched with the f DRT_peak .
For the case of the LSCF charge transfer process (Peak#2) and the LSCF/CGO interfacial double-layer capacitance (Peak#1), characteristic frequencies were identified with comparable values as the ones reported by Yurkiv et al. 61 For the electrochemical process on the L65SCrN fuel electrode there is a lack of reported EIS data which makes difficult the precise identification of the processes.  Nonetheless, since no additional high frequency processes are visible in the EIS spectra it is suspected that the L65SCrN charge transfer process are correlated with large chemical capacitance linked to its oxygen non-stoichiometry which overlaps with the gas conversion process. For this impedance feature, the DRT analysis suggests two processes (Peak#3 and Peak#4) at a frequency between ~ 3 -25 Hz that, by analogy with the behavior of MIEC materials reported by Adler et al., 62 are likely to be connected to each other and be characteristic of the response of the L65SCrN. A last process (Peak#5) at low frequencies between 0.1 Hz and 1 Hz was attributed to the gas losses where there is an overlapping of the RWGS reaction (catalytic conversion) 63 , the diffusion in the electrodes and the gas conversion processes (within the gas channels). At frequencies above 10 3 Hz, it was not possible to fit the EIS spectrum due to an artifact of the measurement.
Overall this corresponds to a first approach to model and understand the electrochemical behavior of ESC with a L65SCrN fuel electrode in co-electrolysis. However, this would need to be confirmed and further investigated to understand in detail the different electrochemical processes that take place within the cell and does not preclude other approach to better reflect the behavior of the electrode materials. Especially, with the proposed model, if two connected processes are reasonable to consider for the impedance feature of the L65SCrN, a parametric study varying temperature, current density and gas compositions at fuel electrode with DRT studies will be needed in order to understand the response of the L65SCrN fuel electrode and make a clear process assignment. 64

Fig. 9
Cross-section view of ESC implementing the L65SCrN electrocatalyst as fuel electrode after rSOC operation at 860 °C.
After the above described electrochemical tests were performed, the tested ESC was observed by SEM as shown in Fig. 9, where no delamination or mismatch at none of the interfaces platinum current collector-L65SCrN, L65SCrN-CGO20, CGO20-3YSZ and CGO20-LSCF could be observed on the polished cross section, which suggests good thermo-mechanical compatibility between the different cell components. Elemental analysis performed by EDS revealed the presence of silicon, likely in the form of SiO 2 species, within the CGO20 barrier layer ( Fig. S6 and able S1 in ESI). Therefore, long-term evaluation in co-electrolysis with ultra-pure feed water was performed with a new cell in order to exclude any exogenous sources of degradation and clearly assess the stability of the L65SCrN fuel electrode in operation in view of SOC applications.

Evaluation of the long-term stability in co-EC operation
Long-term steam co-electrolysis at 860 °C was carried out during 950 hours in galvanostatic condition at a fixed current density of -0.46 A•cm -2 at 860 °C with an initial voltage of ~1.3 V (Fig. 10a). EIS measurements at OCV (Fig. 10c and 10d) were performed in order to monitor the evolution upon operation. Compared to the previous cell (Fig 8), a higher R ohm was measured likely due to a contact issue between the electrodes and the current collectors for this long-term test. However, the polarization resistance (R pol ), for testing times above 677 hours, but possible already from the first 100 hours, coincides with the one reported on Fig. 8c, being ca. ~ 0.3 Ω•cm 2 for both cases. The same accounts with the imaginary part of the impedance Z Im from Fig. 8d and Fig. 10d, which are in the same order of magnitude, i.e. close to -0.1 Ω•cm 2 in both cases.
The first 100 hours of test are marked by a decreasing voltage of the cell, meaning an improvement of the cell performance (Fig 10 a and  10b). This improvement is expressed by a decrease of both R ohm and R pol over time, which could be due to the following factors: (i) the oxygen electrode needed ca. 100 hours to be correctly contacted, due possibly to the gold mesh current collector stabilization.
(ii) a removal of impurities that are present in trace amount in the CO 2 that may have been adsorbed at the surface of the perovskite and poisoned the electrode during the 24 hours of cell operation at OCV, before the start of the galvanostatic durability test.
(iii) an activation process of the L65SCrN fuel electrode to accommodate the defect chemistry with the testing conditions (~ 100 hours) during which one can speculate further Nickel exsolution to occur.  57 Despite the difference in operating conditions and the lower applied current density, such a low voltage degradation rate reported for an ESC with L65SCrN fuel electrode is promising. After cooling, SEM investigation of the non-polished surface of the L65SCrN fuel electrode after 950 hours of co-electrolysis operation was performed (Fig. 11). As expected, the contamination of the electrode by silicon was negligible, suggesting a minimal impact on the transport properties of the L65SCrN fuel electrode. (EDS analysis shown on ESI in Fig. S7 and table S2). Investigation of the surface of the L65SCrN perovskite grains revealed the presence of Ni exsolved nanoparticles well distributed at the surface of the perovskite grains over the whole area of the analyzed sample. This is illustrated in a representative manner in Figure 11 and Figure S8. The average size of those exsolved nanoparticles was estimated in the range of 10 to 25 nm on this sample, which is significantly lower than the one estimated on the powder sample at 900°C in pure hydrogen (Fig. 6b).
Considering the 950 hours operating time, this observation suggests that particle coarsening was limited and did not occur in the testing conditions. This is consistent with observations reporting that the size of the exsolved nanoparticle increases until a critical stable radius, as a function of the temperature and time. 46,53 Knowing that the oxygen partial pressure pO 2 was about ~ 10 -24 bar in the case of powder sample reduction at 900°C, 67 and only ~ 10 -21 bar in this case of long-term co-electrolysis, 30 it seems that the pO 2 plays the most important role on the morphology of the exsolved nanoparticles. The relative predominance of the parameters that affect the Ni exsolution morphology could be qualitatively ranked as following: pO 2 > temperature > time. Since electrolysis operation implies reducing conditions to the fuel electrode, it is nonetheless reasonable to expect that the applied current density and thus the overpotential at the L65SCrN fuel electrode is going to have an additional influence on the exsolved Nickel nanoparticles. Since Asite deficiency and exsolution are intimately correlated, it is difficult to decorrelate the impact of the two phenomena on the cell performance. Working with A-site deficient perovskite is of high interest with perspective of industrialization since it enables to control the impurities (as the B-site elements) in the produced materials. ESC are intrinsically characterized by a high operating temperature to counter-balance the ohmic losses induced by the thickness of the electrolyte. This feature limits the impact of the electrocatalysis on the overall cell performance. However, given the contingencies of the ESC cell configuration, we believe that the observed behavior and performance is to a large extend due to the exsolution of Nickel.

Conclusions
Lanthanum strontium chromite perovskites with Nickel partial substitution were investigated for the sake of developing an alternative electrocatalyst to traditional cermets as fuel electrode for SOC applications. The materials have been synthesized and their propensity to exsolve Ni nanoparticles under exposure to reducing atmosphere has been investigated ex-situ. Introduction of a deficiency up to 5% on the Asite of the perovskite was shown to be effective to enhance the exsolution capability of the synthesized materials (L65SCrN), compared to a full stoichiometric perovskite (L70SCrN). The density and particle shape of the exsolved nanoparticles on the surface of the perovskite were shown to be sensitive to the crystallographic orientations of the surfaces and the pO 2 . This behavior is consistent with the observations made on other families of perovskites with exsolution of nanoparticles such as titanates. However, in contrast to the titanates, the evolution in shape, size and coverage ratio of the Ni nanoparticles, upon temperature increase characterized by a particle coarsening suggest that particle -particle interactions prevails over particle -substrate interactions at the surface of lanthanum chromites. This significant morphological change of the nanoparticles upon operating conditions could affect their catalytic activities over time and thus impact the overall performance and durability of an electrode made of these materials.
Tested on cells and with optimal contacting solutions, L65SCrN electrodes demonstrated performance levels that are comparable with the ones of state-of-the-art cermet fuel electrodes: in fuel cell, electrolysis and co-electrolysis modes which are representative conditions for a reversible SOC system. Excellent voltage stability was reported in co-electrolysis operation over 950 hours with a voltage degradation of about 3.2 mV / 1000 hours. Qualitatively it is suggested that pO 2 is the main factor governing particle size followed by the temperature and then time. This suggests that the nanoparticles can be dimensionally stable when the system is operated isothermally, or when exsolution takes place at a temperature higher than the nominal operating temperature of the cell. Therefore, considering a SOC stack implementing cells with L65SCrN fuel electrodes, one can reasonably expect that the exsolution would take place during the commissioning of the stack, yielding coarsened Ni nanoparticles that are dimensionally stable during operation, fulfilling the durability requirements.
However, as a disadvantage the high temperature thermal treatment that is usually performed for stack commissioning would yield coarsened nanoparticles that may impede further electrode performance optimization. One aspect to consider and being advantageous would be to maintain the exsolved nanoparticles as fine as possible to optimally boost the performance of L65SCrN electrodes by tuning for instance the pO 2 .
Additional investigation of the exsolution phenomena by varying parameters such as temperature, time and pO 2 would be thus needed in order to better understand the mechanisms of exsolution. Another important aspect to evaluate is how the Ni nanoparticle size impacts the performance of the electrode and how to fine tune the exsolution parameters for maximizing electrocatalytic activity of this electrode. This would enable to improve the presented L65SCrN fuel electrode and to develop durable electrode morphology for rSOC applications.

Conflicts of interest
There are no conflicts to declare.