# Growth and characterization of strontium vanadate films deposited by reactive magnetron sputtering

Axel Rouviller, Moussa Mezhoud, Christian Dufour, Clara Grygiel, Ulrike Lüders, Xavier Portier, Fabrice Gourbilleau

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Axel Rouviller, Moussa Mezhoud, Christian Dufour, Clara Grygiel, Ulrike Lüders, et al., Growth and characterization of strontium vanadate films deposited by reactive magnetron sputtering. APL Materials, 2025, 13, pp.061111. 10.1063/5.0253471. hal-05101868

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# Growth and Characterization of Strontium Vanadate Films Deposited by Reactive Magnetron Sputtering

Axel Rouviller,¹ Moussa Mezhoud,² Christian Dufour,¹ Clara Grygiel,¹ Ulrike Lüders,² Xavier Portier,¹ and Fabrice Gourbilleau¹

¹) CIMAP, Normandie Université, ENSICAEN, UNICAEN, CNRS, 14050 Caen Cedex 4,
France

²) CRISMAT, CNRS UMR6508, Normandie Université, ENSICAEN, UNICAEN, 14050 Caen Cedex 4,
France

(*Electronic mail: axel.rouviller@ensicaen.fr)

(Dated: 9 May 2025)

Indium-tin oxide (ITO) is the most commonly used transparent conducting oxide (TCO) but suffers from resource depletion. Among all the alternatives to replace it, $SrVO_3$ is a promising candidate which has been highly studied in recent years. In this study, structural, electrical and optical properties of 30 nm-thick films grown on low-cost substrates by reactive sputtering of a $SrVO_3$ target are presented. The structural, electrical and optical properties of the obtained films have been compared with those of our previous study in which thin films were grown using a $Sr_2V_2O_7$ target, illustrating that the use of an $SrVO_3$ target allows us to obtain layers with lower roughness and better opto-electronic properties. Moreover, it was observed that the layers are non-stoichiometric compared to $SrVO_3$ which can explain their semiconducting properties instead of the expected metallic behaviour.

## I. INTRODUCTION

$SrVO_3$ (SVO) is a strongly correlated metal which has received extensive research in recent years¹. Indeed, the results achieved up to now make it a promising transparent conducting oxide (TCO) for replacing ITO (Indium-Tin Oxide), which is the most used TCO and is subject to significant market price swings due to the low abundance of indium on earth¹. The electrical conductivity of this SVO material is guaranteed by its important charge carrier density, which is in the same range as that of typical metals ($10^{22} \text{ cm}^{-3}$) to which must be added an optical transparency in the visible range. This last point is due to the important effective mass $m^*$ of the SVO electrons because of the electronic correlations which cause a shift of the reflection edge of the light by the free carriers out of the visible range, contrary to typical metals. Thanks to its interesting electrical and optical properties, this material could be used in microelectronic devices such as photovoltaic panels or light-emitting diodes.

However, in order to have these desired physical properties, SVO needs to crystallize in its cubic form which unfortunately drastically reduces its development possibilities since the obtention of this crystallographic phase requires the use of expensive substrates such as $SrTiO_3$ (STO), $LaAlO_3$ (LAO) and $(LaAlO_3)_{0.3}(Sr_2TaAlO_6)_{0.7}$ (LSAT), which makes them incompatible with the microelectronic industry¹⁻¹². Concerning the growth process, only few articles present the achievement of cubic SVO thin films by sputtering. Among them, we can cite Jung et al. who have demonstrated the growth of crystallized films on LSAT substrates by a reactive sputtering of an $Sr_2V_2O_7$ target¹³. Marks et al. have shown that heating a layer in a reducing atmosphere after a growth by RF magnetron sputtering permits the obtention of crystallized $SrVO_3$ on an STO substrate¹⁴. Finally, after a growth on low-cost substrates, Sharma et al have identified

an orthorhombic phase in their $SrVO_3$ films, which are transparent in the visible range but highly resistive¹⁵. The study of the growth of $SrVO_3$ using magnetron sputtering and low cost substrates is therefore interesting and challenging to propose a low-cost process compatible with the microelectronics requirements. Unfortunately, the growth of cubic $SrVO_3$ on Si substrates or glass ones results in a lower crystalline quality in comparison to the specific substrates detailed above that permit an epitaxial growth. Nevertheless, recent studies have demonstrated that crystalline $SrVO_3$ layers may be produced on such less expensive substrates by the simple use of an anatase $TiO_2$ buffer layer¹⁶,¹⁷.

Following this idea, a recent study made by our group has shown the possibility to obtain transparent and conductive strontium vanadate layers by sputtering of a $Sr_2V_2O_7$ target on a silicon substrate on which an anatase $TiO_2$ buffer layer was deposited beforehand by means of atomic layer deposition (ALD)¹⁸. We thus showed that with an optimisation of the growth temperature ($T_G$ and the ratio between $H_2$ and Ar fluxes ($r_H$) we can obtain semiconducting films with resistivity values between 1 and $2 \times 10^{-3} \Omega \cdot \text{cm}$, which is lower than those of semiconductors classically used as TCOs such as undoped $SnO_2$¹⁹⁻²¹, undoped $ZnO$²²⁻²⁴ and Al doped ZnO without annealing²⁵⁻²⁷. However, in order to reduce the oxygen provided by the sputtering of the target, a significant amount of hydrogen in the plasma is necessary for growing a SVO film with required TCO properties. This leads to an important roughness of this layer driven by a competition between deposition and etching phenomenon due to the hydrogen species present in the plasma²⁸,²⁹. This is why, with the recent development of the perovskite sputtering target market, we have pursued our study using an $SrVO_3$ target with the aim to improve the quality and the properties of our layers obtained to date with a $Sr_2V_2O_7$ target.

In this work, we report experimental investigations on the
growth of strontium-vanadate films by reactively sputtering of an $SrVO_3$ target on Si substrates with a $TiO_2$ buffer layer. We investigated the influence of the substrate temperature $T_G$ and the hydrogen rate $r_H$ on the structural, electrical and optical properties of the obtained films. The results of this study have been compared with the results reported in our first study¹⁸.

## II. EXPERIMENTAL DETAILS

The strontium vanadate layers of this study were grown by reactively sputtering a 2" target of $SrVO_3$ by applying a radio frequency of 45 W. The Ar plasma was enriched with hydrogen and its pressure was set at 3 mTorr for all growths. The substrates used were, (100) oriented, p-doped Si single crystals with a resistivity in the 5-15 Ω.cm range. Beforehand to the growth of the SVO layer, 10 nm of anatase $TiO_2$ buffer layer has been grown by ALD. Its growth parameters are the same than those found by Jolivet et al.³⁰. Two main sputtering growth parameters have been studied to analyze their influence on the structural, electrical and optical properties of the strontium vanadate layers. These are (i) the substrate temperature $T_G$ which has been varied from 350 to 550°C with a step of 50°C for an hydrogen rate of 50% and (ii) the hydrogen rate $r_H$ with values of 0, 10, 30 and 50% while $T_G$ was fixed at either 450°C, 500°C and 550°C.

In order to analyze the influence of these growth parameters on the crystallization of the SVO films, X-ray diffraction (XRD) experiments were carried out using a D8 Bruker Discover diffractometer equipped with a Cu source and a monochromator ($\lambda_{k\alpha1} = 1.5406$ Å). Crystallographic analyses were also made by transmission electron microscopy (TEM) using a double corrected cold FEG JEOL ARM200 operated at 200 kV. TEM images were treated using the DigitalMicrograph software by Gatan. The EDX spectrometer was a CENTURIO equipment from JEOL and allowed us to obtain chemical maps of our multilayer structures. The samples were prepared for TEM observations by focused ion beam (FIB) using a Dual beam FEI Helios nanolab 660. Cross sections were prepared such that the TEM observations were done following a [110] direction meaning that the electron beam was parallel to the various interfaces. The surface of the films was investigated by atomic force microscopy (AFM) in tapping mode using a Nanoscope III device from Digital Instruments. The average roughness of the samples were calculated by the obtained AFM images using the Nanoscope software. The compositions of the samples were investigated by Rutherford Backscattering Spectroscopy (RBS) at an energy of 2MeV and normal incidence. Once the measurements acquired, the experimental spectra were fitted using the software SIMNRA in order to determine the elemental concentration of the samples.

For the electrical characterizations, four probe method using a Keithley 2450 sourcemeter allowed us to deter-

mine the resistivity of our samples at room temperature. Moreover, for the most interesting samples, the resistivities were measured from 5 to 300 K using the Van der Pauw configuration and Hall measurements were carried out to determine their carrier concentration and their mobility using a magnetic field which varied from -9 to +9 Tesla. Both measurements were performed with a physical properties measurement system (PPMS) by Quantum Design. In order to eliminate the contribution of the conductivity of the substrate during the electrical characterizations, all the samples used for these experiments were grown on substrates containing an additional 8 µm thick thermal $SiO_2$ layer between the Si substrate and the $TiO_2$ buffer layer. Note that due to its electrical properties and its thickness, this latter does not modify the electrical properties of the SVO layers.

Finally, optical characterizations were carried out by spectroscopic ellipsometry using a UVISEL 2 from Jobin-Yvon. For the measurements, the incident angles of the light were 65, 70, and 75° with respect to the normal incidence, and its energy varied from 1.5 to 6 eV. The experimental spectra were fitted using the software DeltaPsi2 allowing the thickness and optical properties of the films to be obtained.

## III. RESULTS AND DISCUSSION

### A. Structural properties

Figure 1a presents the XRD patterns of strontium vanadate films grown at different substrate temperatures, $T_G$, for a given hydrogen rate of 50%. With the increase of $T_G$, we note the emergence of peaks at the angles 32.5°, 40° and 46.2° which are ascribed to the (110), (111) and (200) lattice planes of cubic $SrVO_3$, respectively. These peaks characteristic of the perovskite phase are only visible starting from a growth temperature of 450°C, which is consistent with the results previously obtained with the $Sr_2V_2O_7$ target¹⁸. Therefore, this demonstrates that the growth temperature is a key parameter for the achievement of layers with the desired crystallographic phase when they are deposited on the anatase $TiO_2$ buffer layer. Figure 1b shows XRD patterns of films deposited with different hydrogen rates, $r_H$, for a $T_G$ of 550°C. We note that no cubic $SrVO_3$ peaks are present when the deposition is carried out without hydrogen in the plasma. These peaks start to appear from a value of $r_H$ of 10% and their intensities increase with the hydrogen introduced. This result highlights the necessity to add hydrogen in the plasma during the growth process to initiate the crystallization of the film even when a $SrVO_3$ target is sputtered. Nevertheless, the use of such a target allows us to obtain the cubic $SrVO_3$ layer by using less hydrogen than in the case of the $Sr_2V_2O_7$ target for which a $r_H$ value of 50% was needed to get the required cubic crystallographic phase¹⁸. This result is very promising since a lower roughness of our films than what has been previously achieved can be therefore expected. However, for all films, peaks at 38° and 44° are also observed. A $Sr_3V_2O_8$ secondary
<figure><img src="image_2.png" /></figure>

phase is observed, mostly independent from the used $T_G$ and $r_H$, and independently from the $SrVO_3$ cubic phase formation.

<figure><img src="image_3.png" /><figcaption>FIG. 1. XRD patterns of thin films grown by sputtering of an $SrVO_3$ target a) for different values of $T_G$ and with $r_H = 50\%$, b) for different values of $r_H$ and with $T_G = 550°C$. The vertical orange lines correspond to the bulk peak positions of $SrVO_3$.</figcaption></figure>

We can note on these two XRD patterns that the positions of the peaks (110), (111) and (200) lattice planes, characteristic of $SrVO_3$ are shifted of about 1.52%, 1.48% and 2.33%, respectively, in comparison to what is usually reported (33.0°, 40.6° and 47.3°, respectively)$^{17}$. This shift towards smaller angles can be explained by the stress exerted

by the $TiO_2$ buffer layer on the strontium vanadate film due to the difference in the coefficient of thermal expansion of these materials$^{31–33}$. Another explanation of this greater out-of-plane lattice parameter could be due to a non-perfect stoichiometry of the SVO layer as reported in literature$^{3,12}$. Note also that for any growth parameters, the characteristic peak at around 48.1° of the anatase $TiO_2$ is not observed which may be explained by the fact that the buffer layer is not thick enough. This hypothesis is discussed in the following with the analysis of the TEM images.

Figure 2a highlights the microstructure of a strontium vanadate film grown at 550°C with a hydrogen rate of 50% on a 6 nm $TiO_2$ buffer layer. One can note the presence of a 2 nm thick $SiO_2$ layer at the interface between the Si substrate and the $TiO_2$ layer probably due to the oxidation of the substrate during the growth of this layer by ALD. Considering that the thickness of this $TiO_2$ buffer layer was about 10 nm before the sputtering step, this observation shows that this sublayer suffered an etching from the hydrogen species of the plasma. The lack of diffraction peak from this buffer layer in the XRD pattern can therefore be explained by its very small thickness. However, some lattice planes observed in this thin layer are consistent with the anatase form but some rutile grains cannot be excluded. Regarding the SVO layer, we can observe that the film obtained under these growth conditions is polycrystalline with grains reaching a size of about 20 nm. The HRTEM image presented in Figure 2b is a [001] projection of a cubic structure as depicted by the Fast Fourier Transform shows in inset. The crystallographic phase is then consistent with the cubic $SrVO_3$, and therefore with the XRD measurements. This HRTEM image clearly shows that the crystallite is tilted with respect to the substrate. This phenomenon is consistent with the fact that the strontium vanadate film is grown on polycrystalline $TiO_2$ buffer layer with differently oriented grains as shown by Jolivet et al$^{30}$. Therefore, the $SrVO_3$ grains obtained are also randomly oriented. This observation is consistent with the XRD measurements which do not show texturing for our strontium vanadate films, indicating that the film is polycrystalline with no preferential orientation of the crystallites. Figure 2c is a STEM-HAADF image of the cross sectional view of the structure and the Z contrast reveals clearly the various layers and more particularly one can distinguish the bottom $TiO_2$ film from the SVO layer. This latter shows a brighter contrast due to a higher mean Z value. Finally, Figures 2d-g illustrate the O, Sr, Ti and V chemical maps in our structure. It reveals the presence of V atoms in the $TiO_2$ buffer layer indicating a diffusion of V atoms during the growth of the SVO film. This diffusion is emphasized in Figure 2h which presents the depth profile of Ti and V atoms. Similar diffusions have been observed before in literature in the presence of an anatase $TiO_2$ buffer layer$^{17,18}$.

The surfaces of the samples grown with different $r_H$ at 550°C were investigated by AFM (Figures 3a, b, c and d). These observations reveal that the roughness is homogeneous over the entire surface of the films and that it increases with
<figure><img src="image_4.png" /></figure>

This is the author's peer reviewed, accepted manuscript. However, the online version of record will be different from this version once it has been copyedited and typeset.
PLEASE CITE THIS ARTICLE AS DOI: 10.1063/5.0253471

<figure><img src="image_5.png" /><figcaption>FIG. 2. (a) Bright field TEM image of a SVO film grown on (100) Si on which a TiO₂ buffer layer was grown beforehand. TiO₂ was grown by ALD and SVO was grown by sputtering at 550°C and using rH = 50%. (b) HRTEM image of a SVO region with the corresponding FFT in inset. This latter is in agreement with a [010] direction of SrVO₃. (c) STEM HAADF image. (d), (e), (f) and (g) are respectively the corresponding STEM EDX chemical maps of O, Sr, Ti and V. (h) depth profile of Ti and V atoms.</figcaption></figure>

the hydrogen rate. This evolution has been quantified by two methods consisting on one hand by calculating the average roughness from AFM images and on the other hand, by deducing it from modelling of ellipsometry measurements. First of all, the evolution of the average roughness with the two methods is similar, which demonstrates that the ellipsometric modeling describes accurately the behavior of the SVO layers (Figure 3e). One can observe an increase of the average roughness as a function of r<sub>H</sub> explained by the etching phenomenon that appears during a reactive sputtering with hydrogen species in the plasma<sup>28,34</sup>. However, this roughness is significantly less important than what was observed using an Sr<sub>2</sub>V<sub>2</sub>O<sub>7</sub> target, for where layers with

irregularly distributed domains over their surfaces were observed<sup>18</sup>. Such a result gives us reason to expect an enhancement of the TCO properties since a rougher surface results in an increase of the diffuse scattering of conduction electrons, and as a consequence in a decrease of the electrical conductivity<sup>35-37</sup>. The etching phenomenon has also an influence on the growth rate as it is shown on Figure 3f in which the decrease of the average growth rate as a function of the r<sub>H</sub> for a T<sub>G</sub> value of 550°C can be observed. Indeed, we note that the average growth rate for the film grown without hydrogen is about 0.225 nm per minute and it decreases to about 0.086 nm per minute for the one deposited with 50% of hydrogen in the plasma.
<figure><img src="image_6.png" /></figure>

<figure><img src="image_7.png" /><figcaption>FIG. 3. AFM images (5 µm x 5 µm) of the films obtained by sputtering of an $SrVO_3$ target on Si substrates with an anatase $TiO_2$ buffer layer at a growth temperature of 550°C and a) $r_H$ = 0%, b) $r_H$ = 10%, c) $r_H$ = 30%, d) $r_H$ = 50%. e) Evolution of the average roughness and of the roughness derived from the ellipsometric model as a function of $r_H$ for $T_G$ = 550°C. f) Evolution of the average growth rate as a function of $r_H$ for $T_G$ = 550°C.</figcaption></figure>

## B. Electrical properties

Figure 4a displays the evolution of the electrical resistivity at room temperature obtained by the 4 probe method as a function of the $T_G$ for $r_H$ equal to 50%. The $T_G$ range is 350 to 550°C with a step of 50°C. We observe that between 350 and 400°C, the electrical resistivity of the films is higher than $10^{-1} \Omega.cm$ before spectacularly decreasing of about 2 orders of magnitude when $T_G$ reaches 450°C. The electrical resistivity continues to decrease until it reaches a value of $7.89 \times 10^{-4} \pm 0.20 \times 10^{-4} \Omega.cm$ for a growth temperature of 550°C. This trend is consistent with the structural properties of these films since this evolution is concomitant with the emergence of the XRD peaks characteristic of cubic $SrVO_3$ and the evolving crystallization of the layer with the growth temperature (Figure 1a). The gain in using a $SrVO_3$ target is demonstrated by analyzing this first set of resistivity values. In fact, one can note that the electrical resistivity obtained for the film deposited with the couple of parameters ($T_G$ = 450°C, $r_H$ = 50%) is the same, within the uncertainties, as the resistivity of the film obtained with the $Sr_2V_2O_7$ target for the same hydrogen rate but with a growth temperature of 550°C. Moreover, the film grown with the $Sr_2V_2O_7$ target and exactly the same growth parameters ($T_G$ = 450°C and $r_H$ =

50%) has an electrical resistivity of about $2.41 \times 10^{-2} \Omega.cm$, which is almost one order of magnitude higher than the value found with the $SrVO_3$ target¹⁸. The evolution of the electrical resistivity at room temperature as a function of the hydrogen rate for films grown at 550°C is reported in Figure 4b. One can notice that the electrical resistivity decreases with $r_H$, ranging from an insulating behavior when the plasma is free from hydrogen to lower than $1 \times 10^{-3} \Omega.cm$ for $r_H$ equal to 50%. This result confirms the essential role of hydrogen in the plasma for the crystallization of the layer (Figure 1b) and consequently on its electrical properties. We have reported for comparison the best value of resistivity achieved for a layer grown using the $Sr_2V_2O_7$ target with the same $T_G$ and a $r_H$ value of 80% (in violet Figure 4b. One can see that, with the $SrVO_3$ target, we almost reach the same value with only 30% of hydrogen in the plasma and even improve the electrical resistivity when we incorporate 50% of hydrogen during the sputtering process.

<figure><img src="image_8.png" /><figcaption>FIG. 4. Electrical resistivities at room temperature obtained by the 4 probe method for the samples grown with the $SrVO_3$ target as a function of a) $T_G$ with $r_H$ = 50%, b) $r_H$ with $T_G$ = 550°C</figcaption></figure>
These results highlight that the use of a SrVO$_3$ target is more appropriate than a Sr$_2$V$_2$O$_7$ target since the electrical characteristics achieved at room temperature are better with a lower thermal budget and/or with a lower hydrogen rate. Moreover, this reactive sputtering approach demonstrates that one can manage the electrical properties of the SVO films depending on the desired electronic applications. This could be promising for future industrial development compatible with the standards of the microelectronic industry.

The evolution of the electrical resistivity as a function of temperature, measured with Van der Pauw configuration, for different $r_H$ and a $T_G$ value of 550°C is shown in Figure 5a. As already stated, the higher the hydrogen rate is, the lower the electrical resistivity would be. The graph reveals a semiconducting behavior for all the samples with a decrease of the resistivity with measured temperature. This result is surprising since cubic SrVO$_3$ is a metal and therefore its electrical resistivity is supposed to increase following the Fermi liquid model as is reported in the literature both by DFT and experimental studies$^{11,38-43}$. RBS measurements performed on this set of samples show a systematic depletion of Sr with respect to V in the main part of the film leading to a composition Sr$_{1-x}$VO$_3$ with x ranging from 0 to 0.25 depending on the hydrogen rate. This depletion of Sr could be attributed to a different sputtering yield between Sr and V. Note that the diffusion of V into the TiO$_2$ is also noticed. This Sr deficiency, which does not prevent the formation of cubic grains of cubic SrVO$_3$ as evidenced by TEM and XRD experiments, could be at the origin of the semi-conducting properties of the layer, but also of the observed enhanced lattice parameters.

The evolution of the carrier concentration and the mobility, at room temperature, of the samples as a function of $r_H$ is shown in Figure 5b. One can notice an important decrease of the mobility from 14.69 $cm^2.V^{-1}.s^{-1}$ to 0.26 $cm^2.V^{-1}.s^{-1}$ for a hydrogen rate of 10 and 50%, respectively. This is accompanied by a concomitant increase of the carrier concentration which starts from $7.9 \times 10^{19}cm^{-3}$ to reach $3.0 \times 10^{22}cm^{-3}$, respectively. Such results demonstrate the interest of our approach since for the films deposited with 10% of hydrogen, the values achieved for both the carrier concentration and the mobility are in the same range than those of classical TCO materials reported in the literature. Indeed, for Al doped ZnO, $In_2O_3$, Sb doped SnO$_2$, and Nb doped SnO$_2$ the carrier concentrations are around $1 \times 10^{20}cm^{-3}$, and the mobilities vary from 6.5 to $25 cm^2.V^{-1}.s^{-144-52}$. Note that, for highest value of $r_H$, we obtain a film for which the transport properties grow closer to those of ideal SrVO$_3$ for which the electrical conductivity is mainly insured by a large carrier concentration$^{1,3,5,13}$. However, even if the carrier concentration of this film is in the same order of magnitude as that of cubic SrVO$_3$, the mobility is 4 to 5 times weaker compared to monocrystalline films$^{1,3,5,13}$. This could be explained by several phenomena such as the presence of grain boundaries observed by TEM, the presence of the Sr$_3$V$_2$O$_8$ phase, revealed by XRD, which is electrically insulating, and

the non-stochiometry deduced from the RBS measurements. We can thus conclude that this low mobility value is the reason why this film is less conductive than those grown on a monocrystalline substrate$^3$.

For semiconducting materials, two models are currently used to describe the evolution of the conductivity as a function of temperature. The first one considers that the decrease of electrical resistivity with the temperature is explained by the charge carriers passing from a donor level toward the conduction band according to Fermi-Dirac statistics. Consequently, the electrical conduction with temperature follows the law$^{53}$:

$$ln(\sigma) = ln(A) - \frac{\Delta E_a}{2k_B T} \quad (1)$$

where A is a constant, $k_B$ the Boltzmann constant and $\Delta E_a$ the activation energy. The second model is Mott's variable range hopping (MVRH), which is currently used to describe transport properties of disordered semiconductors. In this model, the electrical conductivity is governed by the phonon-assisted hopping of charge carriers from occupied to unoccupied localized states, and its evolution as a function of temperature follows$^{54-57}$:

$$ln(\sigma) = ln(B) - \left(\frac{T_m}{T}\right)^{1/4} \quad (2)$$

where B is a constant and $T_m$ the Mott temperature. For the most conductive sample of this study, grown at 550°C with a hydrogen rate of 50%, the plots of $ln(\sigma)$ as the function of $1000/T$ and $(1000/T)^{1/4}$ are presented in Figure 5c for a temperature range of 5 K - 25 K and in Figure 5d for a temperature range of 200 K - 300 K. For both temperature regimes, the MVRH model is the most accurate to describe the evolution of the experimental data.

The electrical characterizations have shown that a $T_G$ of at least 500°C is necessary to obtain films with a resistivity at room temperature of about $1 \times 10^{-3} \Omega.cm$. Therefore, in the following, we mainly focus the discussion on the influence of the parameter $r_H$ on the optical properties of the films grown at the highest temperatures, i.e. 500 and 550°C. Moreover, in order to complete the electrical characterization, the transmittance in the visible range of the sample grown with a hydrogen rate of 50% and a growth temperature of 450°C will be presented and be compared with that of the sample of our previous study grown with the same growth parameters and the Sr$_2$V$_2$O$_7$ target$^{18}$.

### C. Optical properties

Figure 6a presents the evolution of the real ($\varepsilon_1$) and the imaginary parts ($\varepsilon_2$) of the dielectric permittivity ($\varepsilon = \varepsilon_1 - i\varepsilon_2$) as a function of the incident light energy for the film deposited with $r_H$ values of 0 and 50% and a $T_G$ of
<figure><img src="image_9.png" /><figcaption>FIG. 5. a) Evolution of the electrical resistivity as a function of the temperature for different values of $r_H$ with $T_G = 550°C$, b) Evolution of the carrier concentration and the mobility as a function of $r_H$ with $T_G = 550°C$. Plot of $\ln(\sigma)$ as a function of both 1000/T and $(1000/T)^{1/4}$ for the sample grown with $T_G = 550°C$ and $r_H = 50\%$ c) for T between 5 and 25 K and d) for T between 200 and 300 K.</figcaption></figure>

550°C. The behaviour of the imaginary part of the dielectric function at low energy is drastically different for the two samples with a weak and almost constant value in the infrared range for the film deposited with a plasma free of hydrogen, while an important increase of $\epsilon_2$ is observed for the other layer. Such a feature has been attributed in literature to the reflection and absorption of the light by the free carriers in cubic $SrVO_3$¹,⁵,⁸. This evolution of $\epsilon_2$ is therefore consistent with the structural and electrical measurements, described above, since the increase of $\epsilon_2$ in the infrared region is observed for the crystallized and conductive layer ($r_H = 50\%$), having an important carrier concentration, and not for the insulating one ($r_H = 0\%$). However, the curve of the real part of the dielectric function at low energy for the film grown with 50% of hydrogen does not match with those reported for $SrVO_3$ for which $\epsilon_1$ is negative in the infrared region¹,⁷,⁵,⁸. This evolution is ascribed to the fact that the films present a semiconducting behaviour and not that of a strongly correlated metal¹⁸ and confirms again the electrical measurements.

This Figure 6a shows also a shift towards lower energies

of the peak maximum of the real and the imaginary parts of the dielectric permittivity with the introduction of hydrogen in the plasma. This result can be extended to all the values of $r_H$ from 0 to 50% as well as for a $T_G$ of 500°C as reported in Figures 6b) and 6c), which show the evolution of the position of the maximum reached by $\epsilon_1$ and $\epsilon_2$ as a function of $r_H$ and for both $T_G$ (500 and 550°C). We can observe in these two graphs that the positions of these $\epsilon_i$ peaks continually decreases as a function of $r_H$ for both temperatures. For $\epsilon_1$, the maximum position decreases from about 4.05 eV down to around 3.25 eV, whereas this value for $\epsilon_2$ goes from 4.5 eV to 3.6 eV. The maximum $\epsilon_i$ values found in the literature for films deposited by PLD technique are reported in orange on the graphs and vary between 3.00 and 3.20 eV for the real part and between 3.25 and 3.50 eV for the imaginary part¹,⁵,⁵,⁸,⁵,⁹. Therefore, the evolution of the position of these peaks shows that the introduction of hydrogen in the plasma during the growth process allows us to obtain films with optical properties closer to those of cubic $SrVO_3$. The evolution of the optical index at 1.95 eV, deduced from spectroscopic ellipsometry data, as a function of $r_H$ for the two $T_G$ is presented in Figure 6d. We can observe that the
This is the author's peer reviewed, accepted manuscript. However, the online version of record will be different from this version once it has been copyedited and typeset.
PLEASE CITE THIS ARTICLE AS DOI: 10.1063/5.0253471

<figure><img src="image_10.png" /><figcaption>FIG. 6. a) Plot of ε₁ (blue) and ε₂ (red) for films grown with T<sub>G</sub> = 550°C and r<sub>H</sub> = 0% (dashed lines) or r<sub>H</sub> = 50% (solid lines), b) Evolution of the peak position of ε₁ as a function of r<sub>H</sub> for films grown at T<sub>G</sub> = 500°C and T<sub>G</sub> = 550°C, c) Evolution of the peak position of ε₂ as a function of r<sub>H</sub> for films grown at T<sub>G</sub> = 500°C and T<sub>G</sub> = 550°C, d) Plot of the optical index at 1.95 eV as a function of r<sub>H</sub> for films grown at T<sub>G</sub> = 500°C and T<sub>G</sub> = 550°C.</figcaption></figure>

optical index decreases with r<sub>H</sub> for both temperatures and reaches a value of 1.77 for a hydrogen rate of 50% and a growth temperature of 550°C. This value is closed to the experimentally one of 1.75 reported by Roth et al. and achieved on high quality monocrystalline <i>SrVO<sub>3</sub></i> grown by MBE on LSAT substrates<sup>60</sup>. Considering now the work of Schira et Latouche, they found using density functional theory that the optical index of <i>Sr<sub>2</sub>V<sub>2</sub>O<sub>7</sub></i> is around 2.10<sup>61</sup>. Therefore, for the two presented substrates temperatures, when the films are grown without hydrogen, their optical indexes are closer to that of <i>Sr<sub>2</sub>V<sub>2</sub>O<sub>7</sub></i> and the increase of the amount of hydrogen allows this value to tend towards that of the <i>SrVO<sub>3</sub></i> crystallographic phase. Note that the same trend has been noticed by using a <i>Sr<sub>2</sub>V<sub>2</sub>O<sub>7</sub></i> target. These results demonstrate first, consistently with the structural and electrical characterizations, that the use of a <i>SrVO<sub>3</sub></i> target for growing this new generation of TCO films requires a smaller amount of hydrogen to obtain layers with satisfactory physical properties. The second important point is that, with

this reactive approach, and whatever the oxide target used, i.e. *SrVO*<sub>3</sub> or *Sr<sub>2</sub>V<sub>2</sub>O*<sub>7</sub>, we are able to monitor a large range of optical and/or electrical properties of strontium vanadate thin films which is promising for future industrial development layer.

From ellipsometric data and by using the Tauc model, the indirect optical transition of the layers can be deduced from the linear part of (αhν)<sup>1/2</sup> plotted as a function of hν<sup>62,63</sup> as reported in the insert of Figure 7 which presents the case of the film grown at 500°C for a r<sub>H</sub> of 50%. Figure 7 presents the evolution of the deduced energy of the indirect transition, as a function of r<sub>H</sub> for both temperatures. We can note that for this two growth temperatures, the energy of this indirect transition decreases with the increase of hydrogen content in the plasma. This decrease is most pronounced between 0 and 10%, which, again, can be explained by the crystallization of the layers (Figure 1b). Moreover, for both T<sub>G</sub>, we can note that, starting from a r<sub>H</sub> of 10%,
the values of the indirect transition are close to the one reported in the literature which corresponds to the transition between the oxygen 2p band and the vanadium 3d-t2g band at 2.7 eV⁷,⁶⁴. Therefore, we can conclude, once again, that the introduction of hydrogen in the sputtering chamber during the growth process is mandatory to grow strontium vanadate films with optical properties close to those of SrVO₃.

<figure><img src="image_11.png" /><figcaption>FIG. 7. Energy value of the indirect transition as a function of rH for films grown at Tg = 500°C and Tg = 550°C with a typical Tauc plot in inset.</figcaption></figure>

In order to determine the optical transparency of the films, their reflectivity for a perpendicular incidence has been determined using the following formula :

$$R = \left( \frac{N_2 - N_1}{N_2 + N_1} \right)^2 \qquad (3)$$

where N2 is the optical index of the SVO film, obtained thanks to the ellipsometry measurements, and N1 is the optical index of air (N1 = 1). R being known, the transmittance of the layers has been calculated using the following formula :

$$T = (1 - R - d\alpha) \times 100 \qquad (4)$$

where d and α are the thickness of the layer and the absorption coefficient, respectively. These latter were also extracted from the ellipsometric model. In Figure 8a, the evolution of the optical transparency as a function of the wavelength for films grown at 550°C with different hydrogen rates is presented. We can observe that the higher the rH value, the lower the optical transmittance in the visible range. This evolution is concomitant with the increase of carrier concentration as a function of rH described above (Figure 5b). Therefore, this decrease of the transmittance can be explained by the fact that when the hydrogen content in the plasma is increased, the number of free carriers increases favoring the

absorption or the reflexion of the incident light. This result highlights the effect of the amount of free carriers on the transmittance properties of SVO thin films. Such a trend has already been observed in our previous study as well as Gu et al. who evidenced a lower optical transparency of their SVO layer in the visible range when it became conductive¹⁸,⁶⁵. An additional result that we can note is the shape of the curves of the conductives layers at low wavelengths. Indeed, for these three samples we note a significant decrease of the transmittance in the UV range. This decrease is characteristic of the behavior of cubic SrVO₃ and, it is therefore consistent with all the experimental results described above on the achievement of this phase. This decrease can be associated to the charge transfer transitions between the occupied oxygen 2p states and the unoccupied vanadium 3d states since both theoretical and experimental investigations on the optical conductivity of SrVO₃ have shown that this optical transition has a large peak centered around 3.5 eV (350 nm)⁶⁶,⁶⁷. Moreover, Rath et al have also interpreted this decrease of the optical transmittance near of the UV range as due to this optical transition¹⁷. If we compare the curve of the most conductive film of our previous study, using the Sr₂V₂O₇ target, with the most conductive layer obtained using the SrVO₃ target (rH = 50%), we can notice that the film obtained with this later is slightly more transparent in the visible range. This result allows us to conclude that the use of this new target is more relevant in order to obtain transparent and conductive stontium vanadate layers on low cost substrates using magnetron sputtering since the produced films reach better electrical and optical properties while also having a lower roughness. Finally, we have demonstrated that it is possible to tune the TCO properties of the films by introducing more or less hydrogen according to the physical properties that we wish to favor between the transparency in the visible range or the electrical conductivity.

Since the electrical resistivities at room temperature for films grown at 450°C using both targets have been compared earlier, the transmittance in the visible range for these two samples is now presented on Figure 8b. We can observe that between 400 and 625 nm the sample grown with the SrVO₃ target is more transparent whereas after 625 nm and up to 800 nm it is the opposite. In order to compare the transmittance of these two samples over the entire visible range, the normalized integrated transmittance factor has been calculated using the formula¹⁷:

$$\eta_T = 100 \times \frac{\int_{\lambda_{min}}^{\lambda_{max}} T(\lambda) d\lambda}{\int_{\lambda_{min}}^{\lambda_{max}} T_{max}(\lambda) d\lambda} \qquad (5)$$

where Tmax is equal to 100%, λmin = 400 nm and λmin = 800 nm. The values of ηT for both targets are shown on the insert of Figure 8b. This value in the visible range is equal to 73% for the film grown with the SrVO₃ target whereas it is about 70% for the film grown with the Sr₂V₂O₇ target. The average transmittance in the visible range for a film grown with a temperature of 450°C and a hydrogen rate of 50% is therefore
This is the author's peer reviewed, accepted manuscript. However, the online version of record will be different from this version once it has been copyedited and typeset.
PLEASE CITE THIS ARTICLE AS DOI: 10.1063/5.0253471

almost identical using the SrVO₃ target. We can conclude that thanks to this new target it is possible to grow strontium vanadate films with a growth temperature of only 450°C which have a much better electrical conduction than those obtained from the Sr₂V₂O₇ target, while maintaining the same transparency level. The use of an SrVO₃ target appears to be essential for future growth development at low temperatures, which is an important development for the integration of this material in actual electronic devices.

<figure><img src="image_12.png" /></figure>

<figure><img src="image_13.png" /><figcaption>FIG. 8. Evolution of transmittance as a function of the wavelength for, a) different values of $r_H$ with $T_G = 550°C$, b) different targets with $r_H$ = 50% and $T_G = 450°C$. The inset presents the average transmittance of the films between 400 and 800 nm.</figcaption></figure>

#### IV. CONCLUSION

In conclusion, the structural, electrical and optical properties of Strontium-Vanadate layers grown by reactive sputtering using a SrVO₃ target have been examined. First of all, we have shown that the growth temperature and the hydrogen rate are crucial parameters in order to optimize the crystallization and the TCO properties of the obtained SVO₃ layers.

Moreover, we have exhibited that, for a given growth temperature, it is possible to influence the TCO properties of the films by varying the $r_H$ value which makes it possible to tune the properties of the layers according to the desired optoelectronic applications. Then, the results concerning the non-stochiometry of these films in comparison to the SrVO₃ phase allow us to formulate a hypothesis about their semiconducting behavior.

We have also shown throughout this study that the use of a SrVO₃ target is more relevant in order to achieve the growth of conductive and transparent strontium vanadate thin films. Indeed, we first highlighted that it is possible to obtain the SrVO₃ crystallographic phase by introducing a smaller amount of hydrogen in the plasma during the growth process in comparison with the Sr₂V₂O₇ target. This allows us to considerably reduce the roughness of the layers. Second, we demonstrated that using this target, the produced thin films have better TCO properties than those grown with the Sr₂V₂O₇ target. Third, the better TCO properties for a growth temperature of 450°C, using an SrVO₃ target, opens the gate for future growth studies at low temperatures.

#### V. ACKNOWLEDGMENTS

This research was partly funded by l'Agence Nationale de la Recherche (ANR) through the LEDVAN project "ANR-23-CE50-0005-01". This work was also supported by the CNRS Federation IRMA - FR 3095 and the Normandy Region through the BOSEED project. The authors also thank the ANR institution and the Normandy Region for their financial support in acquiring the FIB setup and the X-Ray diffractometer (ANR-11-EQPX-0020).

The authors thank Cédric Frilay and Jérémie Gonçalves for the preparations of the TiO₂ buffer layers by ALD. Franck Lemarié for the AFM measurements and the FIB preparations. Philippe Marie for this help with the ellipsometry measurements.
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