# Impact of Strongly Correlated Conduction Electrons on the Functional Properties of the Transparent Conducting Oxide SrVO 3

Martando Rath, Moussa Mezhoud, Mudit Upadhyay, Xavier Portier, Sylvain Duprey, Philippe Marie, Julien Cardin, Christophe Labbé, Fabrice Gourbilleau, Simon Hurand, et al.

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## ► To cite this version:

Martando Rath, Moussa Mezhoud, Mudit Upadhyay, Xavier Portier, Sylvain Duprey, et al., Impact of Strongly Correlated Conduction Electrons on the Functional Properties of the Transparent Conducting Oxide SrVO 3. Advanced Materials Interfaces, 2026, pp.e01094. 10.1002/admi.202501094. hal-05520236

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Submitted on 20 Feb 2026

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RESEARCH ARTICLE

OPEN ACCESS

# Impact of Strongly Correlated Conduction Electrons on the Functional Properties of the Transparent Conducting Oxide SrVO₃

Martando Rath¹,² ID

Moussa Mezhoud¹

Mudit Upadhyay³

Xavier Portier²

Sylvain Duprey²

Philippe Marie²

Julien Cardin²

Christophe Labbé²

Fabrice Gourbilleau²

Simon Hurand⁴

Aïmane Cheikh¹

Adrian David¹ ID

Arnaud Fouchet¹

Wilfrid Prellier¹ ID

Ulrike Lüders¹ ID

¹CRISMAT, Université de Caen Normandie, CNRS UMR 6508, ENSICAEN, Caen, France | ²CIMAP, UMR 6252, UNICAEN, ENSICAEN, CNRS-Université de Caen, Caen, France | ³GREMAN, UMR 7347, CNRS-Université François Rabelais de Tours, Tours, France | ⁴Institut Pprime, UPR 3346, CNRS-Université de Poitiers-ENSMA, SP2MI, Chasseneuil-du-Poitou, France

**Correspondence:** Martando Rath (martando.rath@spin.cnr.it) | Wilfrid Prellier (wilfrid.prellier@ensicaen.fr) | Ulrike Lüders (ulrike.luders@ensicaen.fr)

**Received:** 7 January 2026 | **Accepted:** 5 February 2026

**Keywords:** optical conductivity | strongly correlated materials | transparent conducting oxides

## ABSTRACT

Electronic correlations and chemical disorder play a pivotal role in the emerging class of perovskite transparent conducting oxides (TCOs). These materials exhibit charge carrier densities typical of metals, yet achieve optical transparency through the enhanced effective mass, which shifts the plasma frequency beyond the visible range. In this study, we systematically tune the electronic and structural disorder of epitaxial SrVO₃ thin films on SrTiO₃ (001) substrates by annealing treatments. The controlled variations modulate the electronic correlations, leading to gradual changes in optical transparency and electrical resistivity. By analyzing these modifications of the functional properties, we elucidate the role of electronic correlations and disorder in SrVO₃, and propose the design strategies for optimizing the TCO performance of the perovskite oxides.

## 1 Introduction

The discovery of the transparent conducting state of the perovskite oxides SrVO₃ and CaVO₃ has underlined the importance of strongly correlated electronic systems for the development of technologically interesting properties [1–4]. In these systems, the 3d¹ electronic configuration of the vanadium ion is stabilized by the delocalization of the d-electron, leading to a charge carrier density in the range of 10²² cm⁻³ [5–10]. The large effective mass ($m^*/m_0$= 4, from optical experiments) of the conduction electron, typical of this kind of perovskite oxides, results in a reduction in the plasma frequency and therefore in the optical transparency under visible light [7, 11–14]. The strong role of electronic correlations was highlighted in previous studies, as was also the case of SrNbO₃ and SrMoO₃ [1, 3, 9–11, 15–23].

Their excellent TCO properties with resistivities that go down to the µΩcm range and high optical transparencies, leading to Figure of Merits (FoMs) comparable to the most widely used Indium-tin-oxide (ITO) [9, 16, 24–27].

The aim of this study is to highlight the influence of the characteristics of the strongly correlated conduction electrons on the functional properties of the TCO SrVO₃, i.e., its electrical resistivity and the optical transparency. The conduction electrons play a key role in the TCO, as they influence both the optical and the electrical properties. For the electrical resistivity $\rho$, the mobility and the number of charge carriers are the main governing factors, while for the plasma frequency $\omega_p$, the ratio between the charge carrier density and the effective mass is essential [28]. The charge carrier density is primarily determined by the oxidation

This is an open access article under the terms of the Creative Commons Attribution License, which permits use, distribution and reproduction in any medium, provided the original work is properly cited.

© 2026 The Author(s). *Advanced Materials Interfaces* published by Wiley-VCH GmbH
state of the transition metal ion in the perovskite oxide, i.e., the
oxygen stoichiometry of the material [29]. Both the mobility and
the electronic mass are influenced by the disorder in the system.
In strongly correlated systems, they are also affected by the
correlation strength [12, 30]. This correlation can be expressed in
terms of the Coulomb interaction *U*, although this quantity may
represent other structural effects in the lattice [31, 32]. Anyway,
for a metallic system, the increase of *U* leads to changes in the
band structure around the Fermi energy *E*<sub>F</sub>, with a transfer of
states from the partially occupied quasiparticle conduction band
located at *E*<sub>F</sub> to the so-called Lower and Upper Hubbard Bands
(LHB and UHB, respectively) [30, 33, 34]. The LHB is located
below *E*<sub>F</sub> and is fully occupied, while the UHB is located above
*E*<sub>F</sub> and is therefore empty. For large *U* values, the quasiparticle
band disappears, leading to the well-known Mott insulator state.
The perovskite oxide TCOs show metallic conduction in a Fermi
liquid state and have therefore an intermediate band structure,
where LHB and UHB are formed, but the quasiparticle band at
*E*<sub>F</sub> still exists [1, 30, 35, 36].

Such a study of the Hubbard model has been carried out on SrVO3 thin films under different strains, showing the influence of the strain state on the electronic correlations, but fails to address the corresponding effects on the material's functional properties [35, 37]. In our previous studies, we reported the ageing effects in both epitaxial and polycrystalline SVO films [16, 38]. However, the underlying mechanisms and especially their relation to the electronic correlations remain unexplored and can only be accessed through optical conductivity measurements. In this study, we will show that the electronic correlations in SrVO3 epitaxial thin films can also be influenced by ambient air annealing. Heating for 24h at different temperatures ranging from 150 to 250°C gradually disrupts their electronic and crystalline orders in the films, allowing us to observe their effect on the macroscopic properties under the same strain. Optical conductivity spectra are used to analyze the characteristics of the band structure in the strongly correlated state, from which the main influences on the TCO properties are determined. We will show that a transition from a coherently conducting state to an incoherently conducting state has a relatively small influence on the macroscopic properties, while it is the chemical and structural disorders leading to the breakdown of the TCO phase. Therefore, the underlying mechanism of the transparency of the perovskite oxide TCOs is rather robust. In addition, we will be able to extract the main levers to optimize the optical properties of these materials while preserving a conductive state.

2 | Results and Discussion

## 2.1 | Tuning of the Functional Properties of Epitaxial SrVO$_3$ Thin Films by Various Annealing Temperatures

The SrVO₃ (SVO) films were subjected to different thermal
annealing treatments, carried out in the ambient atmosphere for
24h at various temperatures of 150, 200, and 250°C. Figure 1a
shows X-ray diffraction (XRD) θ-2θ scans of the zoomed part
around the (002) reflection of SVO thin films annealed at different
temperatures. The presence of Laue oscillations in the θ-2θ
scans reveals the high crystalline quality of the as-grown film.

<figure><img src="image_3.png" /><figcaption>FIGURE 1 | (a) θ-2θ XRD patterns of as-deposited and 150°C annealed SVO thin films on STO (001) showing the presence of Laue oscillations, that disappear after annealing at 200 and 250°C. (b) X-ray Reflectivity curves of different annealed SVO thin films. Red lines are fitting of experimental data using the GenX reflectivity tool. Each profile is shifted by a factor of 10 for improved clarity. Inset in Figure 1b shows the variation of thicknesses of the samples with the annealing temperatures.</figcaption></figure>

After annealing at 150°C, the intensity of the fringes is slightly
reduced and the peak of the main film shifts to higher angles,
but the structural changes remain small. The structural changes
are much more important for the films annealed at 200 and
250°C. In the 200°C annealed sample, the main SVO (002) peak
shows a strong reduction of intensity and the absence of Laue
oscillations, and for the 250°C annealed film, even the main SVO
peak disappears. Therefore, for these two samples, the annealing
has a strong effect on the crystalline order.

X-ray reflectivity (XRR) was also carried out on the films
(Figure 1b). Since the atomic electronic densities of the SVO film
and SrTiO₃ (STO) substrate are nearly identical, the XRR data
shows only low intensity thickness oscillations for the as-grown
and 150°C annealed films. The increase in the intensity of the
oscillations for the high temperature annealed samples indicates
a change in electronic density, suggesting a chemical alteration of
the films. The change in electronic densities of annealed samples
relative to the as-grown sample is confirmed by the scattering
length density (SLD) obtained from XRR fitting (see Supporting
Information [39] for details). The fitting analysis shows that the
density of the surface layer of air annealed SVO is smaller than
that of the crystalline one. A two-layer model was used to fit
the data with GenX, where a surface layer was added on the
SVO layer in order to capture any chemical changes related to
the annealing. However, the sample heated at 250°C was fitted
with a single layer model, which shows that the sample was
turned into a completely amorphous state after annealing. The
thickness of the films was found to increase from 35.2 nm for the
as-deposited film to 50.3 nm for the film annealed at 250°C, with
intermediate values of 37.4 and 39.1 nm for the samples annealed
at 150 and 200°C, respectively (inset in Figure 1b). To confirm
this thickness increase and identify its origin, cross-sectional
samples were prepared and imaged using transmission electron
microscopy (TEM).

A low magnification bright field TEM image of as-deposited SVO
film on STO substrate is shown in Figure 2a and indicates a
<figure><img src="image_4.png" /><figcaption>**FIGURE 2** | (a) Low magnification cross-sectional bright field TEM image of as-grown SVO thin film on STO substrate, (b) High resolution TEM image of the sample and insets show fast Fourier Transform (FFT) patterns of SVO grown on the STO substrate, (c) Low magnification STEM-HAADF image showing the similar Z contrast between SVO and STO, (d) High resolution STEM-HAADF micrograph showing perfect alignment of [001] SVO film on [001] STO. The dotted line indicates the interface position, and (e-h) STEM-EDX mappings of the elements O, Sr, Ti, and V, respectively.</figcaption></figure>

sharp interface between the bottom of the film and the substrate. Figure 2b displays a typical TEM image of the as-grown sample. The thickness of the SVO film was found to be 37 nm, which is in good agreement with the XRR result. The insets in Figure 2b are the fast Fourier Transform (FFT) patterns of the SVO film and the STO substrate, showing very similar diffraction patterns. This indicates that the film is epitaxially grown on the substrate, with the following orientation relationship: (010) SVO// (010) STO and [001] SVO// [001] STO. Figure 2d shows a high resolution scanning transmission electron microscope energy high-angle annular dark field (STEM-HAADF) image following the [001] direction of both the film and the substrate, and corresponding to a magnified region of the film/substrate interface. In this figure, it can be noticed that the interface between SVO and STO is barely visible as both materials have similar densities (close mean Z values). The dotted line indicates the boundary between the thin film and substrate. An orderly arrayed atomic pattern between the film and the substrate illustrates that the as-grown film is phase pure, crystalline, and stoichiometric without structural defects in STO (Figure 2d). Chemical composition analysis by scanning transmission electron microscope energy dispersive X-ray spectroscopy (STEM-EDX) of the main elements (O K, Sr L, Ti K, and V K) is reported in Figure 2e-h. This analysis demonstrates the homogeneous spatial distribution of the main elements (O K, Sr L, Ti K, and V K) in the film. The absence of diffusion of Ti into the SVO and V into the STO substrate, i.e., compositional uniformity at the interface is evidenced by a sharp boundary between the sample and the substrate (Figure 2d).

As the film annealed at 150°C has not shown a strong change in the structure, its TEM images are not displayed here. On the other

<figure><img src="image_5.png" /><figcaption>**FIGURE 3** | (a) High resolution TEM image of 200°C annealed sample and Fast Fourier Transform (FFT) patterns of SVO grown on the STO substrate, as shown in the insets, clearing few nm amorphous layer (a-SVO) top of crystalline SVO thin film (b) High resolution STEM-HAADF micrograph showing perfect alignment of [001] SVO film on [001] STO. The dotted line indicates the interface position.</figcaption></figure>

hand, the sample annealed at 200°C shows starting crystalline degradation, as observed in XRD. In the low magnification TEM image, an amorphous layer can be observed at the top of the SVO layer (Figure 3a). The STEM-EDX analysis shows an enhanced Sr signal compared to V in the amorphous layer, underlining the different chemical composition of the amorphous layer, forming an Sr-enriched phase typically related to an over-oxidation [40], with respect to the crystalline layer, keeping the Sr/V ratio of 1 (see Supporting Information [39] for details). The driving forces,
<figure><img src="image_6.png" /><figcaption>**FIGURE 4** | (a) Cross-sectional bright field TEM image of SVO thin film on STO substrate, annealed at 250°C in air for 24h, (b) High resolution TEM image of the sample and insets are the FFT patterns of amorphous SVO (a-SVO) film after annealing at 250°C grown on the single crystal STO substrate, (c) Low magnification STEM-HAADF image showing the similar Z contrast between SVO and STO (d) High resolution HAADF-STEM micrograph of the interfacial region showing the amorphous nature of the SVO film except for few unit cells of crystalline SVO (c-SVO) on the top of STO, and (e–h) STEM-EDX elemental mapping of O, Sr, Ti, and V, respectively with inserted V line profile.</figcaption></figure>

such as elastic strain, electrostatic effects, etc., underlying Sr surface segregation in SVO thin films, are not yet fully elucidated. Moreover, this challenge is more complicated by concurrent V oxidation, which makes it difficult to disentangle the individual mechanisms governing Sr segregation [41]. The crystalline part, with a thickness of 18 nm of the SVO at the interface with the STO still exhibits high quality, with well-defined diffraction spots (Figure 3b).

The local structure and compositional variation of the SVO sample after annealing at 250°C are shown in Figure 4. Figure 4a depicts a cross-sectional low magnification bright field TEM image of the annealed sample with a thickness of around 48 nm, as estimated by XRR, which confirms the thickness's increase upon annealing at 250°C. The HR-TEM image in Figure 4b demonstrates the fully amorphous nature of the SVO sample, which is further confirmed by its FFT pattern (the absence of diffraction spots). However, atomically smooth terraces, i.e., few planes of crystalline SVO, close to the STO surface, are still observed by a close inspection of the enlarged interfacial region shown by the high resolution STEM-HAADF micrograph image in Figure 4d. Compositional analysis and mapping by STEM-EDX of the main elements (O, Sr, Ti, and V) are shown in Figure 4e–h. The slight increase of V concentration at the topmost surface and the interface between SVO and STO is detected, as shown by the line profile, inserted in Figure 4h.

Therefore, with the different annealing temperatures, it is possible to induce different structural disorders in the samples. While

the 150°C annealed film is comparable in its crystalline quality to the as-grown, in the 200°C annealed sample, both crystalline and amorphous parts are present. The crystalline part at the interface with STO is continuous, allowing for electrical conduction [7, 16]. The 250°C annealed film is fully amorphous, except for some probably disconnected crystalline parts directly at the interface with STO.

In order to determine the changes in the band structure related to these crystalline modifications, we have carried out, the optical conductivity, $\sigma_1(\omega)=\omega\epsilon_0\epsilon_2(\omega)$ analysis, where $\omega$ is the angular photon frequency, $\epsilon_0$ is the vacuum permittivity, and $\epsilon_2(\omega)$ is the imaginary part of complex dielectric function [28]. The results are shown in Figure 5a. For the as-grown SVO sample, the main characteristic peaks typical for bulk SVO are observed [14, 42]. At low energy, the signal increase is due to the presence of the Drude peak. This peak is related to excitations in the quasiparticle conduction band at $E_F$. At higher energy, two peaks labelled A and B are observed at energies around 1.79 and 3.43 eV, respectively. They were assigned to the specific band structure for strongly correlated systems [4, 14]. In these systems, the quasiparticle band at $E_F$ splits gradually with the electron-electron repulsion $U$ into two bands, the occupied Lower Hubbard Band (LHB) lying below $E_F$ and the unoccupied Upper Hubbard Band (UHB) [30, 43, 44]. The as-grown SVO film shows the characteristic peaks related to these bands: peak A corresponds to the Mott-Hubbard d-d transition, i.e., from the LHB to the quasiparticle band, and from the quasiparticle band to the UHB and peak B is a mixture of the d-d transition with the 2p-3d transition, but a higher spectral
<figure><img src="image_7.png" /><figcaption>**FIGURE 5** | (a) The optical conductivity of SVO film of as-grown and annealed at different temperatures, such as 150, 200, and 250°C. The metallic spectra present the Drude tail at low energy, the peak (A) related to the transfer between the quasiparticle bands and the Hubbard bands, the peak (B) related to the transfer between the Hubbard bands and the charge-transfer transitions (C). The inset shows the enlarged view of the Drude and mid-gap regions of the films. (b) Evolution of spectral weight (W) of three energy regions: W₁ (0.6–1.5 eV), W₂ (1.5–2.5 eV), and W₃ (2.5–3.8 eV) across the four films.</figcaption></figure>

weight of the latter. Therefore, the energy position of the peak A corresponds to *U*/2. The observed position of this peak in the as-grown film is at 1.79, which is very similar to the one determined in refs. [14, 37, 45] for bulk SVO. Thus, the typical band structure for strongly correlated metallic systems is well present in the as-grown film. The high energy region in the optical conductivity spectra is ascribed to the charge transfer (CT) peaks (C) where the signatures are governed by the interband *pd* transitions between the O 2*p* and the V 3*d* bands.

When regarding the evolution in the samples after annealing, again, a gradual change of the electronic structure is observed. After annealing at 150°C (red line in Figure 5a), the position of peak A remains similar, but there is a clear transfer of spectral weight (W) from peaks A to B (inset in Figure 5a). Therefore, annealing leads to a decreased density of the quasiparticle states and an increased density of the LHB, indicating an increase in incoherent transport at the expense of coherent transport. After annealing at 200°C, peak A has an even smaller spectral weight, and peak B decreases in energy. In the model of Makino et al., this evolution corresponds to an increasing loss of coherent transport [14, 46]. The observed shift of the B peak is interesting to analyze. Related to the O 2*p*-V 3*d* transition, this shift indicates a difference in the energy separation of the oxygen bands compared to the V bands. Such a change is complicated to explain in terms of structural analysis, where it was shown that the crystalline part retains a perovskite structure without any fundamental changes compared to the as-grown film. Combined with the absence of the A peak, indicating a very small spectral weight of the quasiparticle band, this shift can be interpreted as a change in the positioning of the UHB with respect to the O bands, and therefore a decrease in *U*. However, more studies are necessary to determine the exact origin of this energy shift. However, the 250°C annealed film no longer shows the characteristic peaks of the SVO, nor does it show the tail of the Drude peak. The only peak still observed is situated in the interband transition region. It is noteworthy that the amorphous surface layer of the annealed samples is insulating and the oxidation state of V is 5+ with a 3*d*⁰ electronic

configuration [3, 7, 16]. This implies that the quasiparticle band at the Fermi energy is absent and any optical transitions related to the *d*-bands should occur at higher energies in the amorphous layer. Hence, the C region of the 250°C aged film is related to the optical transitions in the amorphous layer.

Thus, the different annealing temperatures allow us to vary the electronic band structure of the SVO films and, moreover, the strength and the quality of the electronic correlations. The as-grown SVO film shows the signatures of bulk SVO and will serve, therefore as a reference to interpret the functional properties of the films in the framework of a strongly correlated, semi-coherent quasiparticle system. The film annealed at 150°C allows to study the influence of less coherent transport on the functional properties. The structural, chemical and electronic characteristics are well comparable to the as-grown film, but the spectral weight (as shown in Figure 5b) shifts observed in the optical conductivity indicate an increase in incoherency of the conduction electrons. For an annealing at 200°C, the crystalline structure of the film has a lower quality. Here, the comparison of the results of the dc electrical measurements and the optical results concerning the characteristics of the charge carriers allows discrimination between the conducting crystalline phase, which will be probed by the DC measurements, while the optical measurements will probe the entire film, including the amorphous parts. The sample annealed at 250°C will again serve as a reference, as the film is (nearly) entirely amorphous.

From optical measurements, additional information about the materials and charge carriers can be determined (as summarized in Table 1). Figure 6 illustrates the optical band gaps of as-grown and annealed SVO samples at different temperatures. To determine these values, UV-visible-NIR reflection spectra were recorded as a function of wavelength at room temperature. The optical reflection spectra were then transformed using the Kubelka–Munk function, $F(R) = \frac{K}{S} = \frac{(1-R)^2}{2R}$, where *R* is the reflectance of the materials, *K* is the absorption coefficient, and *S* is the back-scattering coefficient. The Tauc plot, which
<table><caption>**TABLE 1** | Summary of calculations of transport parameters such as effective mass from the Ellipsometry data (together with dc transport data) and optical parameters like zero crossing ($\epsilon_1=0$) or ENZ, $LF_{max}$, scattering parameters from Ellipsometry data, direct optical bandgap ($E_g$) from the reflection data of as-grown SVO and samples with different annealing temperatures.</caption><thead><tr><th>Aging temperature [°C]</th><th>ENZ [eV]</th><th>LF<sub>max</sub> [eV]</th><th>τ [fs]</th><th>m*/m<sub>0</sub></th><th>τ<sub>ee</sub> [fs]</th><th>E<sub>g</sub> [eV]</th></tr></thead><tbody><tr><td>As-grown</td><td>1.3</td><td>1.33</td><td>2.34</td><td>4.72</td><td>6.11</td><td>2.52</td></tr><tr><td>150</td><td>1.22</td><td>1.27</td><td>1.90</td><td>4.67</td><td>3.66</td><td>2.31</td></tr><tr><td>200</td><td>1.17</td><td>1.21</td><td>2.13</td><td>4.39</td><td>1.49</td><td>2.22</td></tr></tbody></table>

<figure><img src="image_8.png" /><figcaption>**FIGURE 6** | (a) Indirect and (b) Direct optical band gaps derived from Tauc plots using UV-vis-NIR reflectance data. The band gaps were determined through linear extrapolation of the onset absorption coefficient data, as indicated by the long dashed lines. Samples were as-grown and annealed at various temperatures on an STO (001) substrate. Inset in Figure 6b: Displays the variation in the direct optical bandgap across all samples. Notably, the optical bandgap of the sample annealed at 250°C in air matches that of the bare STO substrate.</figcaption></figure>

involves linear extrapolation of the onset absorption coefficient data ($\alpha\hbar\nu)^n$ versus photon energy ($\hbar\nu$), provides insight into the optical band gaps of the materials. Here, $F(R)$ is proportional to the extinction coefficient $\alpha$, and $n$ is the electronic coefficient, which is 1/2 for indirect and 2 for direct optical band gaps. Upon analysis, the bandgap evolution for both indirect and direct types results in nearly identical values (Figure 6a Indirect and Figure 6b Direct optical bandgap). The measured direct band gap of the as-grown SVO thin film is 2.5 eV, aligning with previously reported results. Additionally, a decreasing trend in the bandgap was observed with increasing annealing temperatures, as shown in the inset of Figure 6b. However, after annealing at 250°C for 24 hours, the crystalline sample became amorphous, its shape and bandgap resembling that of the bare STO substrate [47]. The notion of the optical bandgap has to be considered carefully in the context of the metallic perovskite TCOs. In semi-conducting TCOs, this value corresponds to the gap in the band structure around $E_F$, and the impurity bands leading to the electrical conduction are not considered. In the case of metallic TCOs, the bands at $E_F$ have a higher density of states, so that, in principle, the initial definition of a bandgap does not hold. On the other hand, the optical bandgap describes the transparency range of the material, which is pertinent also in the case of metallic TCOs. Therefore, we have decided to keep here the notion of the optical bandgap, as an expression of the transparency range to be expected. An interesting observation is that the onset of the optical conduction in the as-grown SVO film is determined by the peak. Thus, the optical bandgap depends mainly on how the oxygen bands are positioned relative to the quasiparticle peak, as electronic correlations have little effect on it. This may be

an interesting aspect of the theory-based development of new strongly correlated TCOs.

One of the most important parameters for TCOs is the reduced plasma frequency $\omega_p$, defining the low energy boundary of the transparency range. However, this energy is largely governed by the charge carriers, i.e., the quasi-particle bands at $E_F$, and therefore less influenced by the electronic correlations. We have determined its value first from the real part of the complex dielectric function $\epsilon(\omega) = \epsilon_1(\omega) + i\epsilon_2(\omega)$ and where $\omega$ is the photon angular frequency, shown in Figure 7a, by its zero crossing $\epsilon_1(\omega_p^*) = 0$, also called epsilon near zero (ENZ) and second, from the energy Loss Function $LF = -\text{Im}[\epsilon^{-1}(\omega)] = \epsilon_2(\omega)/(\epsilon_1^2(\omega)+\epsilon_2^2(\omega))$, calculated in the generalized Drude model, where strong plasmon peaks are visible at $\omega_p^*$ i.e., the maximum of the loss function spectra ($LF_{max}$) (inset of Figure 7b). Both values are given in Table 1. For the as-grown SVO film, $\omega_p^*$ takes a value of about 1.3 eV, comparable to what is found in literature [11]. Regarding the samples annealed at 150 and 200°C, a slight decrease of $\omega_p^*$ is observed, probably reflecting a slight bandwidth change of the quasiparticle band.

In a simple Drude model, the plasma frequency is given by $\hbar\omega_p = \hbar e \sqrt{\frac{N}{m^*\epsilon_0}} = \hbar\omega_p^* \sqrt{\epsilon_r}$, with $m^*$ the electronic effective mass. $\epsilon_0$ and $\epsilon_r$ are the material's vacuum permittivity and relative permittivity, respectively. $\hbar$, $e$, and $N$ stand for the reduced Planck constant, the electronic charge, and the carrier concentration of the material, respectively. Hence, by taking for $N$ the values measured under DC by Hall effect, the extracted values of the effective mass of
<figure><img src="image_9.png" /><figcaption>**FIGURE 7** | (a) Real part of the complex dielectric function $\varepsilon_1(\omega)$ of SVO thin films with different annealing temperatures (as-grown, 150, 200, and 250°C), annealed STO (100) substrate. Horizontal dashed line indicate zero crossings of $\varepsilon_1(\omega)$ of as-grown, 150, and 200°C heated samples. b) Imaginary part of complex dielectric function $\varepsilon_2(\omega)$ of various SVO thin films and inset shows the variation of screen plasma frequency ($\omega_p^*$), i.e., maximum of Loss function spectrum, with the annealing temperatures of the samples.</figcaption></figure>

the as-grown and annealed samples at different temperatures (150 and 200°C) are 4.72, 4.67, and 4.39 $m_0$, respectively, while assuming $\varepsilon_r = 4$ following in ref. [11]. Thus, the effective mass enhances with the observed loss of coherent transport, i.e., the stronger population of the LHB and UHB. Therefore, the Coulomb repulsion is the main influence on the effective mass, as expected for strongly correlated systems.

In the same model, the electron scattering time ($\tau_{ee}$) that correlates the electrical conductivity and plasma frequency ($\omega_p$) is given by $\tau_{ee} = \sigma_0/\varepsilon_0\omega_p^2$, where $\sigma_0$ is the zero-frequency conductivity obtained from the DC measurement and $\omega_p$ is the plasma frequency obtained from the optical spectra. The calculated $\tau_{ee}$ of the as-grown and the samples annealed at different temperatures (150 and 200°C) are 6.11, 3.66, and 1.49 fs, respectively. The global scattering time ($\tau_g(\omega)$) can be deduced from the offset between the plasmon peak position in the loss function $\omega_p^*$ and the value of the ENZ. From the offset or red shift of the plasmon peak, the global scattering rate is given by
$$ \frac{\hbar}{\tau} = \sqrt{E(LF_{max})^2 - (\hbar\omega_p(\varepsilon_1 = 0))^2} $$
According to Matthiessen's rule, the global scattering rate ($1/\tau_g(\omega)$) is the sum of the scattering rates of the different scattering phenomena, including the electron scattering [48]. This global scattering time was calculated to be 2.34, 1.90, and 2.13 fs for the as-grown, 150, and 200°C annealed samples, respectively. Therefore, the global scattering time ($\tau_g(\omega)$) in the conducting part of the sample does not change, but $\tau_{ee}$ shortens with the annealing treatments. This is in accordance with the observed spectral weight transfer from the quasiparticle band at $E_F$ to the LHB and UHB, translating the loss of charge transport coherence and therefore a stronger electron-electron scattering.

## 2.2 | Determination of the Functional TCO Properties

The electrical resistivity of the films is shown in Figure 8a, except for that of the 250°C sample, where the resistivity was too high

to be measurable in the PPMS system. All conducting samples show the typical T² temperature dependence of the Fermi liquid (see Supporting Information [39] for details). The main change observed is an enhancement of the residual resistivity in the films. This finding is most interesting in the case of the 200°C annealed film. The TEM analysis has shown a two-layer film, with about half of the film (≈ 18 nm) being crystalline. As the amorphous phase is insulating, as observed before and in the case of the 250°C annealed film of this study, the observed electrical resistivity can be ascribed only to the crystalline part [7, 16, 49].

When correlated to the observed changes in the electronic band structure, the evolution of the resistivity is consistent. All three films retain their metallic character (at least for part of the film) with signatures of strong electronic correlations, which is consistent with the observation of the peaks related to the LHB/UHB in the optical conductivity. Therefore, even for a large change of the structure as observed in the 200°C annealed film, the remaining crystalline part of the SVO retains its metallic character. However, the strong increase of the residual resistivity of this film appears to point to a strong increase in disorder even in the metallic, i.e., crystalline part of this film, although the local diffraction patterns in the TEM study (Figure 3a) stay well defined.

Figure 8b displays the results of the Hall effect measurements, showing the charge density and charge mobility of different annealed samples (see Supporting Information [39] for details). The as-grown film shows the typical values of the charge carrier density in the $10^{22} \text{ cm}^{-3}$ range. The observed value is with $2.5 \times 10^{22} \text{ cm}^{-3}$, which is slightly enhanced compared to the theoretical value of $1.77 \times 10^{22} \text{ cm}^{-3}$ (counting one 3d electron per V in the film volume). Regarding mobility, the observed values are in the range of the highest values observed in the literature until now [7, 10, 16]. For the sample annealed at 150°C, both the charge density and mobility decrease, although only in a limited way. It can therefore be noted that the gradual loss in charge carrier coherence does influence the electrical properties of the films,
<figure><img src="image_10.png" /><figcaption>**FIGURE 8** | (a) Temperature dependence of electrical resistivity of the epitaxial SVO for different aging temperatures. (b) Charge mobility (left axis) and the charge density (right axis) extracted from Hall effect measurements.</figcaption></figure>

<figure><img src="image_11.png" /><figcaption>**FIGURE 9** | UV-visible optical transmission of epitaxial SVO thin films in the as-grown state and after 24h of aging at different temperatures. The FoMs, which correlate with the optical transmission factor and electrical resistivity of the various annealed samples, are shown in the inset.</figcaption></figure>

where a high degree of coherence is necessary to obtain excellent resistivity values.

For the sample annealed at 200°C, both the charge density and the mobility decrease. Here, only the thickness of the crystalline, i.e., conducting part of the film has been taken into account to calculate the transport characteristics. The charge density decreases only slightly, staying in the range of the theoretical charge density of SVO. The primary cause of the observed resistivity increase appears to be the significant decrease in charge mobility, which drops by a factor of 4 for the sample annealed at 200°C.

The optical transmission of the as-grown and annealed samples are shown in Figure 9. The former shows the typical spectrum of SVO, with a maximum transmission at about 500 nm and a strong decrease at lower wavelengths. In comparison, the samples annealed at 150 and 200°C show very similar optical transmission. Therefore, the observed changes in the band structure, which significantly affect the electrical properties, have only a limited impact on the optical transparency of the annealed films. In the

case of the 200°C sample, the optical properties are therefore governed by the crystalline part of the film, although an enhancement of the transparency is observed due to the amorphous part, acting here as an anti reflecting coating, since it is rather transparent and has a somewhat lower optical index, as can be seen in Figure 7 [25]. The optical results on the 250°C annealed sample highlight the fact, that in the amorphous phase, no distinctive signature is present in the visible range.

In many cases, the optical transparency $T_{opt}$ and the electrical conduction ($\sigma = 1/(R_s d)$) show some trade off in TCOs and figures of merits (FoM) are used to evaluate their performance. $R_s$ is the sheet resistance, and $d$ is the thickness of the film. In Table 2, the results of the electrical and optical properties of the samples are summarized.

The figures of merit (FoM) for TCOs was proposed by Haacke (FoM<sup>H</sup>) [50]. FoM<sup>H</sup> is more widely used in the literature FoM<sup>H</sup> = $T_{opt}^q/R_s = T^q\sigma d$ where $q$ is a parameter being application dependent. $q = 10$ was chosen here because a transmittance of 90% is sufficient for most purposes [50, 51].

From a separate analysis, we have calculated the FoM based on the optical transmission factor $\eta_T = \frac{\int_{\lambda_{min}}^{\lambda_{max}} T(\lambda) d\lambda}{\int_{\lambda_{min}}^{\lambda_{max}} T_{max}(\lambda) d\lambda} \times 100$, where $\lambda_{min}$ and $\lambda_{max}$ are the minimum and maximum wavelengths of interest, respectively, and the electrical resistivity ($\rho$) of the sample. The FoM is given by $FoM = \eta_T/\rho$ [52, 53]. Thus, the FoM is a measure of the total transmittance over a spectral range (400–800 nm), while FoM<sup>H</sup> is a single point measurement. The calculated FoMs decrease, indicating a reduced optical performance in the annealed samples to the as-grown one (inset of Figure 9). The most important influence is the increase of the resistivity in the 150 and 200°C annealed films, therefore, the same conclusions are drawn from the resistivity data.

### 3 | Conclusion

The annealing experiments on epitaxial SVO films grown on STO (001) substrates allow tuning of the degree of electronic and crystalline disorders. This approach enables a systematic investigation of how changes in the electronic band structure, related to the electronic correlations, affect the functional
<table><caption>**TABLE 2** | The variation of sheet resistance ($R_s$), electrical resistivity ($\rho$), thickness ($d$), optical transmission ($T$), and $FoM^H$ of as-grown and annealed SVO samples with different temperatures from 150 to 250°C. It is important to note that for the 250°C annealed sample, the values of $R_s$ and $\rho$ exceed the measurable range of the PPMS system. So it is mentioned out of range for $R_s$ and also for $FoM^H$. However, the electrical resistivity given in the table is calculated from the resistance value obtained after the completion of 24h aging cycle, i.e., from the evolution of resistivity as a function of time, which is recorded from the Keithley 2450 sourcemeter [38].</caption><thead><tr><th>Aging temperature [°C]</th><th>R<sub>s</sub> [Ω/sq.]</th><th>ρ [μΩ-cm]</th><th>d [nm]</th><th>T [%]</th><th>FoM<sup>H</sup> [Ω<sup>-1</sup>]</th></tr></thead><tbody><tr><td>As-grown</td><td>41.3</td><td>113</td><td>35.2</td><td>71.2</td><td>8×10<sup>-4</sup></td></tr><tr><td>150</td><td>76.3</td><td>206.7</td><td>37.4</td><td>72.4</td><td>5.1×10<sup>-4</sup></td></tr><tr><td>200</td><td>400.2</td><td>557.8</td><td>39.1</td><td>76.9</td><td>1.8×10<sup>-4</sup></td></tr><tr><td>250</td><td>Out of range</td><td>1.7×10<sup>9</sup></td><td>50.3</td><td>95.2</td><td>Out of range</td></tr></tbody></table>

properties of this new kind of perovskite TCOs. In addition, the robustness of the correlation effects with annealing treatments can be assessed. After annealing at 150°C, the spectral weight transfer to the Hubbard bands indicates a slight loss of charge carrier coherence, resulting in a decrease in the charge mobility and an increase in the residual resistivity. When annealed at 200°C, the remaining crystalline part of the film shows a much higher resistivity, although the order in the crystalline part is preserved. The breakdown of the quasiparticle band in this sample, observed by optical conductivity, indicates a strongly incoherent transport at the origin of the increased resistivity. However, the optical transmission increased slightly because of the insulating, amorphous and transparent part of the film.

An important role of the electronic correlations and coherent transport on the electrical properties has been identified, while the optical transparency is rather unattained by the changes in the band structure. The FoM of SVO is therefore governed by the electrical resistivity, which should concentrate the research efforts.

The primary interest of this study lies in establishing a relationship between the functional properties of TCOs and the well-known Hubbard model for strongly correlated systems. There is a large number of perovskites that have been described theoretically with the Hubbard model earlier, so we can now rely on these studies in order to identify interesting systems to be realized experimentally and to be characterized. From theoretical calculations of strongly correlated systems, the bandwidth also shows a strong influence on the transport properties [12]. It was not possible to investigate this parameter here in this study, as a priori, no bandwidth variations were observed. However, it would be important to investigate this parameter, where the solid solution of $Sr_{1-x}Ca_xVO_3$ was used in the past and a bandwidth variation is observed due to a structural transition from the SVO cubic structure to an orthorhombic one for $CaVO_3$ [14, 18, 46]. Therefore, the optical band gap is rather governed by the relative position of the oxygen bands with respect to the quasiparticle peak, which is only slightly influenced by the electronic correlation. From extensive theoretical studies, it is quite well known how to influence this parameter, especially by the chemical substitution of the perovskite transition metal site. However, the optical properties were seldom studied, as the TCO application was not the main interest at that moment.

## 4 | Experimental Section

A polycrystalline $Sr_2V_2O_7$ target was prepared by standard solid-state reaction, for the growth of perovskite $SrVO_3$ (SVO) thin film by Pulsed Laser Deposition (PLD). Prior to the growth of thin films, single crystal $SrTiO_3$ (001) substrates were cleaned by ultra-sonication with successive acetone and ethanol solvents. An ultraviolet KrF excimer laser with a wavelength of 248 nm was used during the PLD deposition, carried out under vacuum and at a substrate temperature of 500 °C. SVO films with a thickness of 30 nm were grown on single crystal $SrTiO_3$ (001) substrates as described in more details in refs. [7, 16, 38].

In this study, we conducted thermal annealing on the samples using a hot plate under an ambient atmosphere for 24h. To assess their electrical resistivities, a 4-point-probe was positioned on each sample, placed on the hot plate initially at room temperature. Subsequently, the hot plate's temperature was ramped up in the range of 150 to 250 °C. A constant ramping rate of 20°C/min was maintained during the heating. Once the target temperature was reached, we initiated the resistivity measurements using a Keithley 2450 sourcemeter. The evolution of the resistivity of the SVO films during the heat treatments is described in ref. [38]. After annealing of the samples at various temperatures, we have studied their structural and optical properties.

The crystal orientation and crystalline quality of the different annealed thin films were examined by X-ray diffraction (XRD) with a Bruker D8 Discover with monochromatic Cu Kα₁ radiation (λ = 1.5406 Å). X-ray reflectivity (XRR) measurements were carried out on the same machine. XRR data were fitted using GenX software [54].

The transmission electron microscopy (TEM) observations were performed using a double corrected ARM200F JEOL microscope operated at 200 kV equipped with an energy dispersive X-ray spectrometer (JEOL CENTURION). The spatial resolution using the high-angle annular dark field (HAADF) detector is about 0.8 Å. The TEM samples were prepared by the focused ion beam (FIB) method (Dual beam FEI HELIOS Nanolab 660 equipment). The films were protected by prior depositions of carbon (about 20 nm thick) and platinum (about 1μm thick) films before the grinding process, and the thinning procedure was done until electron transparency.
Optical properties were measured across the UV-visible-Near
Infrared range using a Lambda 1050 Perkin-Elmer spectropho-
tometer in reflection mode with the universal reflectance acces-
sory (URA) and in transmission mode.

Variable-angle spectroscopic ellipsometry (VASE) measurements were performed in reflection mode at room temperature from 0.6 to 6.5 eV (UV-vis-NIR), using a HORIBA Jobin Yvon instrument. The ellipsometric parameters (Δ, Ψ) of the samples are recorded at various incident angles 60°, 70°, and 75° from the sample normal. The optimized incidence angle of 70° was used consistently for all ellipsometry measurements. The complex optical indices (n+ik) (where n is normal refractive index and k is extinction coefficient) and dielectric functions, ε(ω) = ε₁(ω)+iε₂(ω) are extracted from the data using DeltaPsi2 software with the experimentally by XRR, determined film thickness measurement. The experimental ellipsometry spectra collected over the near IR to UV range are fitted with the configuration of air/film/substrate with no surface roughness in the model. A combination of Drude, new amorphous, and Lorentz oscillator functions is used for SVO thin films on the STO substrate. Here, the Drude oscillator is used to represent the free charge contribution in the intraband transition and the new amorphous oscillators, along with an extra Lorentz oscillator, are used to represent the interband transitions, i.e., O 2p → V 3d transition. A constant value ε∞ is added to the real part ε₁(ω) to fit the model with further interband transitions at even higher energy [3, 20, 55].

For transport measurements, electrical contacts were made with an ultrasound wire bonder, allowing for soldering through the surface insulating layer due to the destruction of the thin film under the contact. The crystalline SVO being the only conducting material in the sample, we can therefore conclude on the conduc- tion by the SVO layer. Low temperature resistivity and Hall effect measurements were carried out using the van der Pauw contact geometry in a commercial Physical Properties Measurement System (PPMS) from Quantum Design, under magnetic fields of ±9 T.

Acknowledgments

The authors would like to acknowledge financial support from the CNRS
prematuration project CoCOT, the Normandy region through the projects
RIN PLDsurf and Cibox, and the funding of the Ph.D. thesis of M.M.
This work is also supported by the French government, managed by the
Agence Nationale de la Recherche under the France 2030 programme,
under the reference ANR-23-EXES-0001, and by the Normandy region,
and the CNRS Federation IRMA (FR 3095).

**Conflicts of Interest**

The authors declare no conflict of interest

**Data Availability Statement**

The data that support the findings of this study are available from the
corresponding author upon reasonable request.

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## Supporting Information

Additional supporting information can be found online in the Supporting Information section.

**Supporting File:** admi70435-sup-0001-SuppMat.pdf.