# In situ observation of phase transitions in La2NiO4+δ bulk and thin film samples via Raman spectroscopy

Adeel Riaz, Alexander Stangl, Mónica Burriel, Michel Mermoux

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Adeel Riaz, Alexander Stangl, Mónica Burriel, Michel Mermoux. In situ observation of phase transitions in La2NiO4+δ bulk and thin film samples via Raman spectroscopy. Acta Materialia, 2025, 301, pp.121570. 10.1016/j.actamat.2025.121570. hal-05408775

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Submitted on 10 Dec 2025

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# *In situ* observation of phase transitions in La₂NiO₄+δ bulk and thin film samples via Raman spectroscopy

Adeel Riaz<sup>a,b,*</sup>, Alexander Stangl<sup>a</sup>, Mónica Burriel<sup>a,b</sup>, Michel Mermoux<sup>b,*</sup>

## ARTICLE INFO

**Keywords:**

Lanthanum nickelate

Ruddlesden-Popper phases

*In situ* Raman spectroscopy

Ceramics

## ABSTRACT

La₂NiO₄+δ has attracted increasing interest in recent years, both as oxygen electrode in solid oxide fuel cells and electrolysers due to its high electrochemical activity at intermediate-to-low temperatures, and as key component of memristive devices for neuromorphic computing, owing to its variable oxygen stoichiometry. The integration of La₂NiO₄+δ into devices operating at different temperatures and oxygen partial pressures requires knowledge of the effects of hyper-stoichiometry (δ) on its crystalline structure. La₂NiO₄+δ is known to accommodate oxygen at interstitial sites allowing for large δ values, up to ~ 0.16. In addition, the O-doping - temperature phase diagram is known to be complex, exhibiting several phase transitions with increasing δ. Herein, we use Raman spectroscopy to monitor the effects of O-doping in the phase diagram and the various structures it contains. Throughout this work, we studied this material in its usual ceramic form, as well as in the form of thin films. Results are discussed in terms of phase transitions, chemical expansion, and some of the possible consequences of the low mean grain size inherent to such thin films.

## 1. Introduction

In recent years, one of the main challenges in the field of reversible Solid Oxide Cell (rSOC) materials has been lowering the operating temperature (T), typically to well below 800 °C [1–5]. For such a purpose, potential materials for oxygen electrode must fulfill several requirements, namely: (i) high electronic and ionic conductivities, (ii) high catalytic activity for oxygen reduction reaction (ORR) and oxygen evolution reaction (OER) as well as (iii) high stability under working conditions. In this context, rare-earth nickelates of formula Ln₂NiO₄+δ (Ln = La, Nd, Pr), have been widely studied as potential mixed ionic-electronic conducting (MIEC) oxygen electrode materials [6–9]. Throughout this paper, we will focus on lanthanum nickelate, La₂NiO₄+δ (L2NO4). In particular, it has recently been shown that tailored, architecturally designed in the form of thin films, L2NO4 oxygen electrodes can effectively improve the charge transfer kinetics by increasing the specific surface area [10]. This specific thin film architecture may be considered a first step to enable operating temperatures around 500 °C.

Rare-earth nickelates, and more generally their isostructural cuprate counterparts, were largely studied in the 1990's for their low temperature electronic transport and magnetic properties, and most importantly

for their superconductivity at high temperatures [11–16]. These oxides belong to the series of the Ruddlesden-Popper (R-P) family (Lnₙ+₁NiO₃n+₁), being the n = 1 member of the family [17]. Interestingly, rare-earth nickelates are known to accommodate interstitial oxygen, charge compensated by electron holes localized on the transition metal cation, leading to the observed mixed ionic and electronic conductivity. The available range of interstitial oxygen content or oxygen hyper-stoichiometry (δ) is known to depend on the size of the Ln cation radii and to be influenced by the synthesis and/or annealing conditions [18]. Concerning L2NO4, reported δ values usually lie in the 0 - 0.16 range, but it was shown that specific electrochemical processes can increase interstitial oxygen content up to δ ≈ 0.3 [19].

The structural and vibrational properties of L2NO4 as a function of δ can be studied using Raman Spectroscopy (RS). It is an inelastic light scattering technique, mostly described as a vibrational spectroscopy method. However, this description can be somewhat restrictive, as in this excitation frequency range, incident photons essentially probe electrons, while vibration modes just modulate the sample's electronic susceptibility and, hence, its optical properties. Nevertheless, we will adhere to this conventional description in the following discussion. Clearly, when analyzing (single) crystals, X-ray diffraction (XRD) and RS

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

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

<footnote>* Corresponding authors.</footnote>

<footnote>E-mail addresses: adeel.riaz@grenoble-inp.fr (A. Riaz), michel.mermoux@grenoble-inp.fr (M. Mermoux).</footnote>
complement each other, although not all space groups (SG) can be unambiguously identified by both methods [20], as each technique follows its own “selection rules”. One of the key differences between the two methods is that XRD probes mainly cationic sub-lattices, while RS is usually sensitive to oxygen-dominated vibrational modes. This is especially true in the high frequency region of the vibrational spectra of oxide materials, where oxygen makes a dominant contribution to the vibrational density of state (VDOS). The effects of disorder in the anionic sub-lattice can, therefore, be assessed, if necessary, although the result will depend on the material and the extent of the deviations from stoichiometry. Both methods are also sensitive to strain/stress effects, through the observation of Raman line (or XRD peaks) shifts and/or splitting, depending on the crystal symmetry and on the nature/-symmetry of the strain, isotropic or anisotropic. According to the quasi-harmonic approximation, the strain dependence of the phonon wavenumbers may be understood introducing the so-called Grüneisen parameters [21]. In its microscopic formulation, the Grüneisen parameter $\gamma$ is a measure of how a specific phonon frequency $\omega$ is altered under a small change in the geometry of the crystallographic unit cell. Although this approach may prove insufficient, particularly due to the isotropic assumption it is based on, which regards $\gamma$ as a scalar quantity, line shifts are still anticipated when the unit cell undergoes contraction or expansion. A simple rule applies: “the closer the atoms, the larger the spring constant values”. In most instances, this rule holds true and, at least qualitatively, explains the observed line shifts with reliable orders of magnitude. What is more, RS can be sensitive to disorder in the anionic sub-lattice, through the detection of forbidden modes that violate selection rules, or even broad, continuous signals that may reflect VDOS-like features. Finally, oxygen doping can also modify possible resonant scattering conditions, whether from disorder (simple smoothing of resonance conditions), or from chemical expansion considerations (variation in lattice parameters). For such situations, the analysis should be conducted for different excitation wavelengths. In particular, deviations from stoichiometry can lead to dramatic changes in the shape of the Raman spectra of specific oxides, as those crystallizing in the fluorite structure, see for example ref [22] that concerns cerium oxide. In these specific cases, it is not uncommon to observe forbidden, infrared symmetry-allowed longitudinal optical (LO) modes in the Raman spectra. In most cases, a Fröhlich electron-phonon [23] interaction is invoked to analyze/understand the presence of these modes in the Raman spectra.

As previously mentioned, in terms of applications, L2NO4 in thin-film form has attracted increasing interest in recent years for its potential applications as an oxygen electrode in reversible solid oxide cells (rSOCs), owing to its high electrochemical activity at intermediate to high temperatures [6,10,24–27] as well as in memristive devices for neuromorphic computing [28,29]. From a more academic perspective, L2NO4 seems quite amenable to study for different reasons. L2NO4 can accommodate oxygen at interstitial sites giving large over-stoichiometry $\delta$ values, up to ~0.16 or even higher. The O-doping phase diagram (T – $\delta$) of L2NO4 is known to be complex (and most likely sample-dependent, as indicated by the existing literature). Raman data on L2NO4 are limited [15,30,31], and there is no clear consensus among the few available datasets, Ambiguities remain in the interpretation of the peaks/modes observed, and, to date, no Raman results have been reported for thin films.

In this study we first investigate the two end-member phases of bulk L2NO4 samples, in line with previous literature, and address ambiguities in the interpretation of their Raman spectra. In addition, we analyze L2NO4 in thin film form (< 500 nm). This is of significance as the grain size in thin films is much smaller than in bulk samples leading to an increase in the density of defects, which can affect the phase diagram [32,33]. Phase transitions and an increase in grain boundaries can significantly impact oxygen ionic conductivity in various materials, thereby affecting the performance of devices used in applications such as low-temperature reversible solid oxide cells (rSOCs) and memristors

[34–36]. For example, Rodriguez *et al* [34] observed a surprisingly high oxygen transport properties in LaMnO₃ thin films due to a high concentration of defects and an unexpected phase transition at 500 °C. The increase in defect density can also lead to additional peaks or modes which need to be correctly identified to understand the changes during *in situ* or *operando* measurements. It is possible to observe phase transitions in bulk L2NO4 using XRD, however, with thin films it is very difficult due to the very high intensity of the substrate and the preferred growth orientation of thin films leading to several peaks not being observed which are typically used to identify these phase changes. RS can be used to identify phase transitions as it is a surface sensitive technique and it is possible to completely block the signal from the substrate. The rationale behind this study was to use *in situ* RS, in complement to XRD, to explore the high temperature structural behavior of L2NO4 thin films, using bulk pellets as reference, and to probe the O-doping phase diagram.

## 1.1. La₂NiO₄+δ structure and phase diagram

The layered structure of L2NO4, which belongs to the K₂NiF₄-type structure, is built of perovskite-type LaNiO₃ and rock-salt-type LaO layers alternating along the crystallographic c-axis. Because of the small size of the La ion, the La–O bonds are in extension while the Ni–O bonds are compressed. These distortions relax in different ways by inserting oxygen in the La–O layer. As a result, the O-doping phase diagram (T – $\delta$) is rather complex, and presents a succession of ordered domains and phase separation regions. The O-doping is not isovalent, meaning that holes are created for charge compensation, accompanied by a formal change of the Ni oxidation state. The hole concentration is usually approximated as 2$\delta$, which is likely a rough estimate, particularly in the case of thin films. The T – $\delta$ phase diagram of bulk L2NO4 has been studied over the years by several groups with different techniques, in particular X-ray and neutron diffraction [19,37–43]. Based on the literature, not all authors provided consistent boundaries or the same sequence of phases. Some of the key features are represented by the schematic phase diagram proposed by Odier *et al* [43], which is reproduced in Fig. 1, in particular the sequence of monophasic and biphasic domains. Synthesis methods have been suggested as the key factor influencing the determination of boundaries between the different phases formed [37,38,40]. Crystallite size variations may represent another parameter influencing phase transitions. To the best of our knowledge, this factor has not been previously considered.

To describe this O-doping phase diagram, it is common to first consider the ideal phase with highest symmetry. This ideal L2NO4 structure, commonly referred to as the High Temperature Tetragonal (HTT) phase, adopts a tetragonal structure with space-group (SG) I4/mmm or $D'_{4h}$. For the exact stoichiometric composition ($\delta = 0$) this phase is observed at elevated temperatures, typically above ~500 °C. This ideal tetragonal structure corresponds to an almost perfect matching of the bond lengths of the perovskite and rock-salt layers. However, as a function of temperature and $\delta$, the crystal cell often presents a lower symmetry than the ideal K₂NiF₄ tetragonal parent structure. In most studies, focusing on temperatures above ambient and for increasing $\delta$ values, L2NO4 undergoes sequential phase transitions from orthorhombic Bmab (Low Temperature Orthorhombic, LTO) [40], to pseudo-tetragonal Pccn [38] or P4$_2$/ncm (Low Temperature Tetragonal, LTO) [40], then to the HTT I4/mmm or $D'_{4h}$, and occasionally to orthorhombic Fmmm (High-Stoichiometry Orthorhombic, HSO) [41] SG, depending on $\delta$ and T, as shown in Fig. 1. Skinner *et al* [41] reported the HSO phase at 150 °C with $\delta = 0.17$, while Aguadero *et al* [19] reported it for $\delta = 0.3$. Biphasic regions have been evidenced as well [37, 43,44]. These biphasic regions can be understood as the result of miscibility gaps which follow from the non-uniform mixing of interstitial oxygen defects into oxygen-poor and oxygen-rich domains. It is believed that strong interstitial oxygen correlations give rise to these
<figure><img src="image_3.png" /><figcaption>Fig. 1. Schematic O-doping phase diagram (T – δ) for L2NO4, based on the phase diagram from [42]; LTO: low temperature orthorhombic, HTT: high temperature tetragonal, HSO, high stoichiometry orthorhombic.</figcaption></figure>

miscibility gaps [37]. In these regions, the oxygen-poor domains take the maximum δ value of the bordering pure low- δ phase and the oxygen-rich domains the minimum δ of the pure high- δ phase. As shown in Fig. 1, an excess oxygen concentration smaller than δ ≈ 0.01 is sufficient to induce phase separation into oxygen-poor and oxygen-rich phases. It should be noted that some of these transitions are indeed observed near room temperature (RT). Finally, the existence of a monoclinic (space group C2) superstructure at δ = 0.25 has also been reported at RT. This structure is four times larger than the average orthorhombic unit cell, resulting in the formation of a solid-solution with a composition of La8Ni4O17 [45]. It will not be considered further in the following discussion as the δ values achieved in this study are below this extreme value.

From an experimental perspective, XRD analysis provides characteristic reflections that are useful for distinguishing between the different phases and in determining phase fractions in biphasic materials [39,44]. Due to the orthorhombic distortion, the LTO phase is characterized by the splitting of certain reflections with h ≠ k, such as (200) / (020). In addition, weak superlattice reflections such as (212) or (032) are expected, an indication of coherent octahedral tilts. In the case of the LTT phase, the splitting of reflections with h ≠ k is absent, but the superlattice reflections (212) and (032) persist. This implies that mixed LTO/LTT phases can be identified by the coexistence of split and non-split reflections, along with the presence of two sets of the superlattice reflections for (212) or (032), due to the difference in lattice constants between the LTO and LTT phases. It is commonly accepted to give and compare lattice parameters on the basis of the $\sqrt{2a} \times \sqrt{2a} \times c$ parameters of the parent HTT cell. Finally, the room temperature evolution of the lattice parameters as a function of δ is known from different works [37,39,44], which gave rather consistent values. In the composition range that correspond to the broad HTT phase field, one can estimate the chemically-induced strain then as follows: ≈ -5 x 10⁻³ (a-axis) and ≈ +8 x 10⁻³ (c-axis), with the consequence that the unit cell volume does not vary dramatically in this composition range, at least close to RT.

The Raman analysis of L2NO4 requires knowledge of the expected normal modes. The k = 0 normal-mode enumeration of the different expected phases is available in a number of studies [46–49]. As previously mentioned, the highest symmetry HTT phase adopts a tetragonal structure with space group I4/mmm or $D_{4h}^{17}$ (also non-conventional F4/mmm symbol is sometimes used for a face centered cell of the parent I4/mmm structure). The atomic displacement patterns

corresponding to the different Raman modes of HTT tetragonal and Bmab orthorhombic L2NO4 are illustrated in Figure S1 and Figure S2, respectively, adapted from Ref [15,50]. The HTT tetragonal primitive cell includes one formula unit and, as a result, there are 21 optical modes of 2A₁g + 2Eg + 3A₂u + B₂u + 4E_u symmetry, from which the 4 “gerade” modes are Raman-active. According to all available computations, these Raman-active modes consist of the out-of-plane (c-axis, A₁g) motion and the in-plane (a and b-axis, Eg) motion of O and La ions [15,50]. Close to the stoichiometric composition (δ ≈ 0), the RT (LTO) structure is the orthorhombic Bmab (or $D_{2h}^{18}$) space group with two formula units in the primitive unit cell. Here, the phase transition from the ideal tetragonal structure to the orthorhombic one is induced by the rotation of the rigid NiO₆ octahedra about the original [110] tetragonal axis [40]. The sense of rotation and displacement for neighboring Ni sites needs to be opposite in sign, which results in a doubled unit cell whose a and b axes are approximately 2 times larger than and 45 to the original a axis. As a result, 42 vibrational modes are expected (5Ag, + 3B₁g + 4B₂g + 6B₃g + 4A_u, + 7B₁u, + 8B₂u, + 5B₃u), in which the 18 gerade modes are Raman active. As mentioned in [40] the LTO phase is also described with a slightly distorted orthorhombic cell, LTO, leading to the Pccn ($D_{2h}^{10}$) space group. In such a case, the irreducible representation becomes 8Ag + 8B₁g + 10B₂g + 10B₃g + 11A_u + 11B₁u + 13B₂u + 13B₃u. The 36 gerade modes are Raman active. Slightly increasing δ, and still close to RT, a second phase transition is observed leading to another tetragonal structure (LTT) (P42/ncm, $D_{4h}^{16}$). The structural difference between the LTO and LTT phases arise from the axis of rotation of the oxygen octahedral: in the LTT structure the NiO₆ octahedra rotate by equal amounts about the [110] and [1-10] axes of the simple tetragonal cell, doubling the in-plane unit cell as in the LTO case [40]. Again, the cell is doubled below the phase transition compared to the LTO one, which means that 81 modes (5A₁g + 3A₂g, + 3B₁g + 5B₂g, + 10Eg, + 3A₁u + 8A₂u + 8B₁u + 3B₂u + 13E_u) are expected. In this case, the A₂g, A₁u, B₁u and B₂u are silent so that 23 Raman-active modes are expected. Finally, the oxygen rich (δ ≥ 0.15) phase is sometimes indexed in the orthorhombic system (Fmmm, $D_{2h}^{23}$) or even F222, $D_{2}^{7}$), due to a slight orthorhombicity in the [110] direction of the I4/mmm structure [41]. In both cases, 21 modes are expected. For the Fmmm structure, the symmetry of the normal modes is the following: 2Ag + 2B₂g + 2B₃g + A_u + 4B₁u + 5B₂u + 5B₃u, among them the gerade modes are Raman active. Intuitively, the Fmmm structure is expected to be only a small distortion of the I4/mmm phase. In this case, a splitting of the Eg modes of the
parent phase into modes of B₂g and B₃g symmetry is expected. For the less symmetrical F222 phase, the irreducible representations are 3A + 4B₁ + 7B₂ + 7B₃. Among them, 3A + 3B₁ + 6B₂ + 6B₃ are Raman (and IR) active. In the latter case, the inversion center in the unit cell is lost, meaning that the zone center vibrational modes are both Raman- and infrared- active.

In most cases, the knowledge the density of vibrational states (VDOS) is of fundamental importance, often a prerequisite for the study of disordered materials. The measured and calculated phonon dispersion curves of tetragonal L2NO4 are reported in [46,47,49]. From the reported cutoff frequencies, no first-order eigenmodes are expected above ~700 cm⁻¹. In the case of the HTT structure, this frequency corresponds to a E₀ symmetry zone center mode, actually observed at ~660 cm⁻¹.

Regarding L2NO4, the primary experimental Raman data are described in [15,30,31,51], while additional data for the isostructural cuprate material can be found in [11,16,31,52–54]. A comprehensive summary of these datasets is provided in Table S1. Notably, the lattice dynamics of the La₂CuO₄ orthorhombic phase at the center of the Brillouin zone has been extensively characterized in [50], which details both the vibrational frequencies of the normal modes and their corresponding eigenvectors.

However, research over the past decades has been predominantly focused on La₂CuO₄-derived superconducting oxides, with most investigations targeting the structural and vibrational properties of these materials at temperatures significantly below room temperature (RT). Furthermore, as previously noted, discrepancies persist in the reported literature. These inconsistencies likely arise not only from experimental challenges but also from variations in sample nature and quality, contributing to a degree of scatter in the obtained results.

## 2. Experimental methodology

La₂NiO ₄+δ bulk samples were prepared using a modified Pechini method. The details of the preparation methods are described in [41]. La₂NiO ₄+δ thin films were deposited by pulsed-injection metal organic chemical vapor deposition (PI-MOCVD) on Si substrates covered with a 100 nm Pt layer. The platinum layer ensures that no modes from the substrate are detected as the metallic layer reflects all the incident photons. Dense thin films of 40, 200 and 500 nm L2NO4 were deposited at 750 °C. 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 3.15 and a solution concentration of 0.02 mol·l⁻¹. The solution was injected into the reactor with a frequency of 4 Hz, opening time of 2 ms. 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.

X-Ray Diffraction (XRD) in Bragg Brentano (θ-2θ) configuration were obtained using a Cu-Kα (λ = 1.54 Å) source on a Bruker D8 Advance diffractometer. The Raman spectra were recorded on two different Renishaw inVia Raman systems, mostly with an excitation wavelength at 532 nm. It was equipped with an air-cooled CCD detector, different specific gratings (1200, 1800 and 2400 grooves/mm), and dielectric rejection filters (cut-off frequency at about 80 cm⁻¹). Frequency calibration was performed using a silicon standard, whose frequency was set at 520.5 cm⁻¹. This configuration ensured a spectral resolution, more exactly an accuracy on the frequencies determination, much better than 1 cm⁻¹. Excitation wavelengths 488, 633 and 785 nm were periodically used to discriminate between Raman and photoluminescence (PL) origins, and to probe possible resonant conditions. Within the exception of samples close to stoichiometry, the line shape of the signals in the main wavenumber range investigated (≈80–1320 cm⁻¹) did not depend on the excitation wavelength, which in particular ensured that of the recorded signal were of Raman scattering origin. The in situ Raman measurements were conducted in a Linkam THMS600 temperature cell connected to different gas supplies, including a 10% H₂/Ar mixture,

synthetic air and oxygen. Measurements were conducted in a backscattering configuration; a long working distance 50x optic was used to focus the incident light at the sample and collect the scattered radiation. As polycrystalline samples were investigated, the polarization of the scattered light was not analyzed.

## 3. Results

Initial measurements on L2NO4 thin films revealed that RS typically provided identifiable signatures with, however, some data variability, including broad signals and the presence of weak intensity peaks. Furthermore, similar observations were made when analyzing ceramic-type L2NO4 samples prepared via conventional and well-established synthesis methods. These findings highlighted the necessity of obtaining additional Raman data on this specific compound, particularly in its ceramic/polycrystalline form, to ensure the reliability of the study. Consequently, we aimed to collect this information, which will be compared to existing studies and progressively presented in the following sections. In the subsequent step, we will focus on L2NO4 in its thin-film form.

The primary objective of this study was to identify and characterize the distinct Raman signatures associated with the end-member structural phases of L2NO4 at both room and elevated temperatures, with a particular focus on the thin-film form. Furthermore, by comparing the Raman spectra of bulk samples and thin films, we aim to distinguish the intrinsic signals of L2NO4 from those associated with the higher defect density typically present in thin films. To this end, we first used a bulk sample prepared using the usual citrate route, then sintered to obtain a ceramic pellet. The bulk L2NO4 sample was successively heated for 9 hours at 575 °C in 5% H₂/Ar and for 1 hour at 200 °C in dry air, respectively, to obtain δ values close to 0 (reduced sample) and close to δ ~0.13 (oxidized sample). The corresponding RT diffraction patterns after the high temperature treatments are shown in Fig. 2a-b. The XRD pattern of the reduced sample was effectively refined (Rietveld method) with an orthorhombic Bmab model with no additional peaks observed (see Figure S3). Splitting of diffraction peaks such as (200) and (313) and appearance of Braggs peaks with h+k = 2n+1 (where n is an integer) could be effectively observed, in complete agreement with the Bmab SG model, see for example [38,39], which confirms the LTO phase assignment. The refined cell parameters a = 5.5390 Å, b = 5.4669 Å and c = 12.5439 Å are in close agreement with published data for to δ ≈ 0 [37,39]. As expected, the SG of the oxidized sample was found to be tetragonal I4/mmm with refined cell parameters of a = 5.4674 Å (reduced cell parameter √2a) and c = 12.6691 Å, corresponding to δ ≈ 0.13, according to [37–39].

Fig. 2c-d shows the corresponding Raman spectra of the reduced and oxidized samples. As previously mentioned, as these L2NO4 samples are in the form of polycrystalline ceramics, selection rules cannot be considered. Nevertheless, it is possible to assign some of the peaks, as the frequencies and the corresponding ionic motions and Raman tensors are described in a number of previous studies [15,30]. As shown in Fig. 2c and expected from the mode enumeration, many peaks were observed for the reduced sample (LTO phase), at about 105, 150, 165, 230, 260, 307, 382, 440 and 658 cm⁻¹. According to [50], other Raman-active modes are expected below the cutoff frequencies of the instruments used in the work. The lines observed at 224 and 675 cm⁻¹ appear as shoulders, but the resolution of the measurements is sufficient to evidence them. These lines were consistently observed for several oxidation/reduction cycles. As for the spectra obtained for the reduced phase, only the relative intensities fluctuate from one spectrum to the other.

Despite the limited data available in the literature for L2NO4, this spectrum looks similar to that given in references [30,55,56] with the exception of the doublet observed at around 658 cm⁻¹, which was not consistently observed throughout the references quoted above. This spectrum (Fig. 2c) is also similar to those known for the isostructural cuprate materials [11,16,31,52–54], again with the exception of the
<figure><img src="image_4.png" /><figcaption>**Fig. 2.** (a) XRD pattern of the reduced (stoichiometric, $\delta = 0$) and oxidized sample ($\delta = 0.13$), (b) magnified spectra showing the splitting of peaks in the LTO phase. The red markers correspond to the tetragonal I4/mmm phase (ICDD: 04-015-2147) and the black markers correspond to the orthorhombic Bmab phase (ICDD: 04-010-9707); Room temperature Raman spectra of the (c) reduced orthorhombic structure, and (d) oxidized tetragonal L2NO4 structure.</figcaption></figure>

high frequency 658 cm⁻¹ line, for the same reasons. From the polarized measurements conducted on single crystals [31], Raman modes at about 105, 225, 260, 440, and ≈ 500 cm⁻¹ are expected to be of $A_g$ symmetry. Without formal proof from our data, the other modes are expected to have $B_g$-type symmetry. According to the literature survey, no Raman-active mode is expected around 660 cm⁻¹. On the other hand, as already mentioned, this frequency corresponds quite closely to that of an active infrared mode of the HTT structure of $E_u$ symmetry with an observed low LO-TO (longitudinal optical and transverse optical) splitting, see for example [57].

For the oxidized sample (Fig. 2d), two main lines are clearly observed, peaking at about 224 and 437 cm⁻¹, in agreement with most reports [15,30] and computation studies [50]. In the oxidized high $\delta$ stoichiometry state (HTT phase), 4 Raman-active modes are expected. The fully symmetric $A_{1g}$ modes are expected to have higher intensities than the in-plane $E_g$ modes. The symmetric $A_{1g}$ vibration of the apical oxygen atoms along the tetragonal axis is observed at about 440 cm⁻¹. In spite of some controversy, the second main mode at about 224 cm⁻¹ can be identified as the $A_{1g}$ vibration of the La atoms along the tetragonal c axis. We were able to further confirm this attribution by analyzing samples enriched in ¹⁸O. The *in situ* Raman spectra of the ¹⁸O exchanged and completely back-exchanged (BE) L2NO4 thin film are shown in Figure S4. The peak of the $A_g$ mode, corresponding to the out-of-plane movement of the oxygen ion, shifts with changes in the ¹⁸O concentration, while the peak around 225 cm⁻¹ remains constant, which confirms that this mode is indeed not related to the movement of the oxygen ions but most probably the out-of-plane movement of the La ion.

In order to compare with literature data, we first probed the *in situ* LTT (P4$_2$/ncm) /LTO (Bmab) /HTT (I4/mmm) phase transition sequence as a function of temperature for $\delta$ close to 0. To this end, the

sample was examined in a reducing atmosphere (Ar/5% $H_2$), from liquid nitrogen temperature to about 550 °C. The evolution of the Raman spectra with increasing temperature is shown in Fig. 3. The LTT/LTO phase transition is expected to occur in the -200 to -190 °C range [56,58, 59], which according to Refs [55,56] is a first-order phase transition. The spectra obtained at the lowest temperatures (-196 and -150 °C) are consistent with the RT spectrum discussed above. The main differences, depending on the observed modes, are some line shifts and strong differences in line widths. Only weak differences are observed comparing spectra obtained at -196 °C and -150 °C, which mainly concern weak lines peaking below ca 300 cm⁻¹. This is in line with the only two references on this phase transition investigated using Raman spectroscopy. It seems that the temperature reached in the low-temperature cell was not low enough to fully observe the full phase transition [56]. Increasing the temperature, a strong broadening of all the lines is observed. Most of the modes that are characteristic of the LTO phase progressively shift and weaken, vanishing above 450 °C. In addition, there is a strong downshift of the low frequency mode observed at about 115 cm⁻¹ (104 cm⁻¹) at -196 °C and RT respectively. This mode is consistently assigned to the staggered rotational mode of $NiO_6$ octahedra, as described in Ref [48,50] and effectively shows softening toward the orthorhombic-tetragonal phase transition temperature at 242 °C. It becomes overdamped above 350 °C, triggering the LTO / HTT phase transition. The behavior of this mode, in agreement with a second-order phase transition, closely follows the literature data that concern both L2NO4 and the isostructural cuprate compound [58,60].

Of interest here are the frequencies of the lines peaking at about 223 and 437 cm⁻¹, that are common to the LTO and HHT phases. The linewidth of the line peaking at about 437 cm⁻¹ increases by a factor of about 3 in the temperature range investigated, while its frequency remains
<figure><img src="image_5.png" /><figcaption>**Fig. 3.** (a) *In situ* Raman spectra from -190 °C to 500 °C in dry air flow for bulk stoichiometry L2NO4 (δ = 0) (b) magnified ranges to show evolution of different modes in selected regions.</figcaption></figure>

nominally constant to 1 or 2 cm⁻¹ within the measurement accuracy. While the line broadening seems consistent with the increase in temperature, this is clearly not the case with the rather subtle downshift observed. In most cases, line shifts of around 10 cm⁻¹ are expected in the temperature range investigated here [61]. This is also in line with most of the published results, see in particular [53]. The behavior of signals observed in the 200 to 300 cm⁻¹ frequency range will also be discussed. At low temperatures, 3 lines are clearly observed, which seem to merge into a single broad line for the highest temperatures. The apparent frequency of this line is 240 cm⁻¹, i.e., around 10 cm⁻¹ above the frequency expected for the HTT phase at high temperature. Based on the observations in Fig. 3, the data suggest that the LTO-HTT phase transition remains incomplete at 500 °C.

In the second step, the reduced sample underwent gradual annealing in synthetic air to traverse the oxygen-doping phase diagram as slowly as possible. This process was conducted within a temperature range below 250 °C, where all previously described phases, along with their respective two-phase coexistence regions, are expected to be observed. For each temperature, successive spectra were recorded every 30 seconds until there were no further observable changes in the line shape. At these temperatures the oxygen incorporation kinetics is expected to be slow enough to be able to record subtle changes in the line-shape of the spectra vs exposition times. Below 200 °C, the spectra remained invariant, as shown in Fig. 4a. At 200 °C, we found a progressive disappearance of the characteristic signals of the LTO phase (around 105, 260 and 660 cm⁻¹) with time in favor of those of the HTT phase (only two peaks around 225 and 440 cm⁻¹), as indicated by the Raman spectra in Fig. 4b-c. This gradual phase transition was observed in situ after around 1500 s. Notably, this transition is evidenced by the

apparent average downshift of spectral lines from 240 cm⁻¹ to 225 cm⁻¹, and the broadening of the 440 cm⁻¹ peak. Once again, no clear or systematic shift was observed in the spectral line common to all phases, specifically the vibrational mode at around 440 cm⁻¹, as a function of oxygen enrichment. This implies that the so-called chemical expansion cannot be unambiguously detected in this specific system. This stands in contrast to observations in certain non-stoichiometric oxides, where such effects are more clearly measurable, and also deviates from the initial expectations of this study. Finally, note the sharp decrease in overall signal intensity during this transition, which accompanies the change in color of the sample from yellowish to black (see Fig. 4d-e).

Similar oxidation/reduction in situ protocols were repeated for L2NO4 dense 200 nm thin films deposited by MOCVD. Optimization of deposition conditions to obtain a La/Ni ratio as close as possible to 2 is described elsewhere [28]. The grain size in the thinnest films is around 20 nm as shown by Stangl et al [10]. This grain size is much smaller than bulk ceramics (hundreds on nanometers). For the deposited thin films, the diffraction patterns could be indexed according to the HTT tetragonal phase. We initially encountered significant variations in the Raman signals, particularly as a function of film thickness, c.f. Fig. 5. As demonstrated in the following analysis, we successfully identified the two distinct Raman modes characteristic of the HTT phase. However, these Raman modes were superimposed on a broad background signal, along with weaker spectral features, in particular in the low frequency range of the spectra. In general, the thinner the layer, the more intense this broad signal becomes, as illustrated in Fig. 5. This broad signal has an apparent maximum at around 575 cm⁻¹, and a cut-off frequency at around 750 cm⁻¹. Notably, this cut-off frequency aligns with known experimental and computed vibrational state density of states. It is likely that this signal partly reflects this oxygen partial VDOS, influenced by the small grain size of the crystals (10–20 nm) forming the layers and the high proportion of grain boundaries associated with them (see Figure S5 for SEM and XRD of the 500 nm thin film). Such a high defect density in those grain boundary regions may effectively induce a partial relaxation of the *k* ≈ 0 selection rule. This effect is not observed in the Raman spectrum of the bulk sample as the grain size is larger (hundreds of nanometers). From this point onward, films with a thickness of 200 nm or greater were used for subsequent studies as the spectra show a smaller number of peaks arising from defect associated with the small grain size of thin films.

The *in situ* oxidation tests yielded results similar to those of the bulk material: while the LTO phase was observed, the Raman signals evolved rapidly over time - even at temperatures below 200 °C, as shown in Fig. 6a. The *in situ* reduction of the thin films produced interesting results. Surprisingly, the film got reduced in less than 300 s at 50 °C, as shown in Fig. 6b. The typical LTO modes appear, including the soft mode around 99 cm⁻¹ which is associated with HTT to LTO phase transition. The film was measured at different points, confirming a homogenous reduction across its entire surface. For comparison, the reference bulk sample reduced to the LTO phase at a much higher temperature (450 °C), whereas the thin film exhibited remarkably fast reduction kinetics at a very low temperature (50 °C). The enhanced surface reactivity of the thin film due to reduced diffusion paths and modified vacancy formation energies from the higher grain boundary densities might have contributed to this rapid reduction, however, the exact origin could not be identified. There was no further change in the line spectra after 1800 s, so the temperature was increased to 100 °C where the only change observed was the increase in the peak width and intensity of the mode at 660 cm⁻¹ indicating that the film is already in the most reduced state.

## 4. Discussion

Our findings underscore the challenges inherent in working with ultrathin layers, particularly when the material exhibits nanoscale grain sizes (approximately 10 nm) and a high density of grain boundaries. From the standpoint of the techniques employed in this study, the
<figure><img src="image_6.png" /><figcaption>**Fig. 4.** *In situ* evolution of the Raman spectra of stoichiometric bulk L2NO4 in dry air at (a) temperatures from 25 to 150 °C, (b-c) at 200 °C with time. The positions of the peaks corresponding to the L2NO4 modes from tetragonal (green dotted lines) and orthorhombic (orange dotted lines) phases are marked in the graph; (d-e) Optical surface images of the bulk pellet showing the color evolution from (d) reduced to (e) oxidized phase L2NO4 at 200 °C.</figcaption></figure>

<figure><img src="image_7.png" /><figcaption>**Fig. 5.** Raman spectra of L2NO4 thin films with thicknesses of 40, 200 and 500 nm.</figcaption></figure>

primary effect of using thin films instead of bulk samples was a pro-
nounced broadening of all spectral signals, accompanied by additional
features induced by the high density of crystalline defects - most notably,
the broad Raman band centered at approximately 550 cm⁻¹.

Our results prompt several questions. First, can we conclusively

attribute the observed Raman signatures to the three anticipated pha-
ses? The HTT phase Raman signature is effectively the expected one,
irrespective of material processing (thin film or bulk). The LTO phase
Raman signature was most likely obtained for bulk L2NO4, insofar the
XRD pattern revealed the expected lattice parameters and confirmed the
orthorhombic distortion. Moreover, the observed Raman signature
closely aligns with those reported in previous studies. We successfully
detected comparable Raman signals for materials in the form of thin
films; however, some degree of variability was observed in the mea-
surements. Regarding the LTT phase, the measurements were not
entirely conclusive, indicating potential uncertainties or limitations in
the data obtained. Based on the analysis of all the proposed phase dia-
grams, theoretically this specific phase should have been observed,
particularly during the annealing of the ceramic material in synthetic air
at a relatively low temperature (200 °C). However, the absence of this
phase in the results suggests a discrepancy with the expected behavior.
The T-δ diagram reproduced in Fig. 1, extracted from [42], is based on
data from different works, based on investigations on single crystals
and/or powders. If we consider the data from ref [37,42,44,62] to be the
most relevant on this issue, there are many differences between the
datasets. Firstly, as explicitly written in [42], some differences remain in
the phase determinations obtained from powders and crystals, even for
conditions considered to be close to equilibrium. This implicitly means
that results are sample-dependent, considering for example single
crystal vs powder, as already discussed, not to mention crystallite size
effects. Considering only the 25–200 °C temperature range, the limit
<figure><img src="image_8.png" /><figcaption>**Fig. 6.** *In situ* Raman spectra of a L2NO4 thin film annealed at different temperatures in: (a) oxidizing conditions; (b) reducing conditions.</figcaption></figure>

between which a mixture of phases and that where only the HTT phase is observed vary in the $\delta = 0.06 - 0.1$ range, depending on the study considered. There is therefore a considerable scatter between all the data sets, even for conditions considered close to thermodynamic equilibrium. Probably more importantly, as mentioned in [62], the phases detected are quite sensitive to the cooling rate down to RT. Analysis conditions are therefore of prime importance. However, thermodynamic equilibrium cannot be assumed within the framework of this study. Secondly, according to [42], the shaded areas in Fig. 1 denote $\delta$ ranges the X-ray data were analyzed according to two phases, choice which was however not clearly justified. Ultimately, they concluded only that, in these $\delta$ ranges, the reflections broadened sufficiently to be indexed as a mixture of two phases. Nonetheless, this broadening could also result from an ordering phenomenon—one that cannot be evaluated with the present data.

A striking finding was the high to very high intensity of the line (in fact a doublet) centered at about $660 \text{ cm}^{-1}$ systematically observed close to $\delta \approx 0$ for all the samples studied in both in the ceramic or thin film form. This peak has been observed previously in stoichiometric samples, however, the origins of this peak have been identified or discussed [30, 63]. According to all computations, Raman-active modes are not expected to appear within this frequency range. Different hypotheses are to be considered: i) parasitic phases: none of the signatures related to the parasitic phases likely to be observed, whether due to the synthesis process, sample storage, or even wrong analysis conditions can account for the observed signal; ii) aliovalent doping on the lanthanum sites. The literature survey clearly shows that such a substitution, when exceeding a few atomic percent, can induce significant modifications in the line-shape of the spectra. For instance, studies on the isostructural $\text{La}_2\text{CuO}_4$ compound, as shown in references [54,64], provide examples of these effects. In such a case, symmetry-breaking arguments are commonly invoked, allowing for the detection of symmetry-forbidden

Raman lines. Nevertheless, such an argument does not apply in our case, as samples from two distinct synthesis processes exhibited the same trends. Therefore, there is no rationale for expecting the doping effects to be similar across both synthesis methods; iii) specific resonant conditions. Why should we consider this hypothesis? Resonant scattering has indeed been observed for the isostructural $\text{La}_2\text{CuO}_4$ compound close to the exact stoichiometry, tuning the excitation wavelength in the near infrared range (1064 nm) [65]. For this specific situation, a strong resonant one-phonon scattering was observed for seven peak frequencies that correspond to longitudinal optic (LO) phonons with both $E_u$ and $A_{2u}$ symmetry. The results were interpreted on the basis of a Fröhlich-type electron-phonon interaction mechanism.

For all the samples examined in this study, particularly near $\delta = 0$, the spectra show a clear dependence on the excitation wavelength. This is illustrated in Figure S6, which pertains to the thin film of L2NO4. We focus on this sample for several reasons: i) the peak around $660 \text{ cm}^{-1}$ is not observable at 785 nm, but becomes detectable at approximately 633 nm, becoming clearly visible at wavelengths below 532 nm. This specific mode is effectively resonantly enhanced. The frequency of this mode corresponds closely to that of an $E_u$ mode in the HTT parent phase, making the hypothesis of a Fröhlich-type interaction plausible. In this situation, multiple harmonics of such a mode are often observed. The first harmonic of this mode has been clearly observed. Concerning L2NO4, only one IR-active mode should be concerned by this specific interaction. ii) Fig. 4 shows that this peak around $660 \text{ cm}^{-1}$ vanishes at the onset of oxidation. This may explain why this line has not been observed systematically, as it is only detectable very close to $\delta = 0$. Without precise control of the hyper-stoichiometry, it becomes challenging to observe this line consistently. iii) Apart from this specific mode, all other modes become clearly observable beyond 532 nm, which supports the same conclusion. However, it should be noted that the electron-phonon interaction to be considered is different for the Raman-allowed modes. Finally, the wavelength dependence of the Raman spectrum was also examined for the oxidized sample, that concerns the tetragonal phase. No specific wavelength dependence of the spectra was observed in this case as shown in Figure S6.

Another important finding is related to the mode observed around $225 \text{ cm}^{-1}$ which has commonly been linked to the *in-plane* movement of the apical oxygen ions [15,30]. Following $^{18}$O isotope exchange, the Raman spectra reveal no shift in this mode, suggesting it does not involve oxygen ions. This mode is attributed to the motion of La ions, which explains its insensitivity to the change in mass.

Regarding the line shape of the broad signal observed around $575 \text{ cm}^{-1}$ in thin films it can be correlated to the small average size of the crystallites that constitute the film. It is not a disallowed mode which appears due to breaking of symmetry from the small grain size. This line does not strongly depend on the excitation wavelength, exhibiting the same cut-off frequency at about $750 \text{ cm}^{-1}$ and must therefore be correlated with disorder considerations.

## 5. Conclusions

This work was devoted to a systematic study of the Raman spectra of $\text{La}_2\text{NiO}_{4+\delta}$ samples in the form of bulk pellets and thin films. To eliminate any signal arising from grain size effects and/or defects typically present in thin films, the reference spectra of the reduced orthorhombic phase (LTO) and the oxidized tetragonal phase (HTT) were obtained using a $\text{La}_2\text{NiO}_{4+\delta}$ sintered pellet free of impurities. The sharp peak around $660 \text{ cm}^{-1}$ in the reduced samples (both bulk and thin films) which had been previously correlated with forbidden infrared mode was concluded to be due a specific resonance effect. As for the thin films, a broad Raman peak observed around $575 \text{ cm}^{-1}$ in oxidized samples could be correlated to the small grain size (~20 nm). Moreover, phase transitions from LTO ($\delta \approx 0$) to HTT ($\delta > 0.1$) in both the pellet and thin films, as well as the reverse transition, were successfully observed *in situ* using Raman spectroscopy. *In situ* Raman studies also revealed a
surprisingly fast reduction of the La₂NiO₄+δ thin films at a very low temperature of 50 °C in a reducing atmosphere of 10% H₂/Ar. Finally, it was possible to correct the identification of the mode around 225 cm⁻¹ which had been correlated to the movement of oxygen ions. ¹⁸O isotope exchange revealed that this mode does not involve oxygen ions and is related to the movement of the La ions. The results in this work correct some of the ambiguities in the interpretation of peaks and modes observed in L2NO₄. Moreover, it opens a pathway for *in situ* analysis of L2NO₄ thin films under different temperature and oxygen partial pressure conditions.

### Supporting Information

All sample datasets and materials related to this work are made available under CC BY 4.0 license in the zenodo repository: 10.5281/ zenodo.15209655

## CRediT authorship contribution statement

**Adeel Riaz:** Writing – review & editing, Writing – original draft, Methodology, Investigation, Formal analysis, Data curation. **Alexander Stangl:** Writing – review & editing, Supervision, Formal analysis. **Mónica Burriel:** Writing – review & editing, Supervision, Methodology, Funding acquisition, Conceptualization. **Michel Mermoux:** Writing – review & editing, Writing – original draft, Supervision, Methodology, Investigation, Formal analysis, Data curation, Conceptualization.

## Declaration of competing interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

## 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). This research has benefited from characterization equipment of the Grenoble INP - CMTC platform supported by the Centre of Excellence of Multifunctional Architecture Materials "CEMAM" n ARN-10-LABX-44-01 funded by the "Investments for the Future" Program. The authors acknowledge Prof. Stephen J. Skinner (Imperial College London, U.K.) for preparing and providing the reference La₂NiO₄ pellets.

## Supplementary materials

Supplementary material associated with this article can be found, in the online version, at doi:10.1016/j.actamat.2025.121570.

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