Axel Rouviller
email: axel.rouviller@ensicaen.fr

Moussa Mezhoud

Alex Misiak

Meiling Zhang

Nicolas Chery

Caroline Bonafos

Ulrike Luders

Philippe Marie

Christian Dufour

Xavier Portier

Ulrike Lüders

Fabrice Gourbilleau

Structural, electrical and optical properties of strontium vanadates films grown by magnetron sputtering

teaching and research institutions in France or abroad, or from public or private research centers.

Introduction

Transparent conducting oxides (TCOs) are highly sought materials for the fabrication of devices such as for example photovoltaic pannels or light-emitting diodes. Currently, the most widely used TCO material is ITO (Indium-Tin Oxide). Considering the low abundance of Indium on earth, it is important to find alternatives. In this context, perovskite SrVO 3 (SVO) could be a good challenger to replace ITO since its physical properties are similar. [START_REF] Zhang | Correlated metals as transparent conductors[END_REF] This material belongs to the family of strongly correlated metals and therefore has a high effective mass. Indeed, the greater a material's electronic correlation is, the higher the effective mass m * of the electrons. Strongly correlated metals are ideal candidates to serve as TCOs since their charge carrier density is in the same range as that of typical metals (10 22 cm -1 ), allowing a high electrical conductivity. Moreover, the reflection edge of the light by the free carriers, represented by the screened plasma energy, is out of the visible range due to the high m * value, allowing a significant optical transparency in the region of interest.

Nevertheless, it has been shown that the crystallised phase of SrVO 3 is required to obtain the desired electrical properties since the amorphous one is electrically insulating. [START_REF] Boileau | Optical and electrical properties of the transparent conductor SrVO 3 without long-range crystalline order[END_REF] Unfortu-nately, the obtention of crystallised SrVO 3 needs the use of complex film growth methods and specific substrates such as SrTiO 3 , LaAlO 3 and (LaAlO 3 ) 0,3 (Sr 2 TaAlO 6 ) 0,7 (LSAT), which makes their production hardly compatible with the standards of the microelectonic industry.

In the literature, the growth of SrVO 3 cubic phase was carried out using molecular beam epitaxy (MBE), [START_REF] Zhang | Correlated metals as transparent conductors[END_REF][START_REF] Moyer | Highly Conductive SrVO 3 as a Bottom Electrode for Functional Perovskite Oxides[END_REF] pulsed laser deposition (PLD) 2,4-10 and electron-beam evaporation. [START_REF] Moon | Epitaxial growth of conductive strontium-vanadate films on Si(100) substrates and their electrical resistivities[END_REF][START_REF] Gu | Two-Dimensional Mott Insulators in SrVO 3 Ultrathin Films[END_REF] Only a few articles present the growth of SrVO 3 by sputtering deposition. [START_REF] Jung | Structural and electrical properties of transparent conductor SrVO 3 thin films grown using radio frequency sputtering deposition[END_REF][START_REF] Sharma | Mechanistic insights on electronic properties and electronic/atomic structure aspects in orthorhombic SrVO 3 thin films: XANES-EXAFS study[END_REF] For all these growth approaches, the specific substrates detailed above are required to obtain the cubic phase of SrVO 3 . Their main drawbacks is their price and low compatibility with the microelectronic standards. Thin film growth on low-cost substrates has been attempted by some groups using the sputtering method. For instance, Sharma et al have identified an orthorhombic phase in SrVO 3 films, which are transparent but highly resistive. [START_REF] Sharma | Mechanistic insights on electronic properties and electronic/atomic structure aspects in orthorhombic SrVO 3 thin films: XANES-EXAFS study[END_REF] On the other hand, Jung et al have obtained promising results by reactively sputtering an Sr 2 V 2 O 7 target under an hydrogen rate of 25% in the plasma. They conclude that it is possible to grow SrVO 3 crystallized films on LSAT substrate at a temperature of 400°C. Moreover, they showed that thanks to an annealing, it is possible to obtain conductive SVO films on low cost substrates such as Si/SiO 2 and Si/TiO 2 . At last, Marks et al have shown the possibility to obtain crystallized SrVO 3 on SrTiO 3 substrate by a two steps process. First of all, amorphous SVO has been deposited by RF magnetron sputtering, then, the crystallization was carried out thanks to an heating in a reducing atmosphere. [START_REF] Marks | Solid-phase epitaxial growth of the correlated-electron transparent conducting oxide SrVO 3[END_REF] The growth of SrVO 3 on low-cost substrates using sputtering is therefore an interesting field of research to further investigate.

The main issue with film growth on low-cost substrates is the lower crystallization in comparison to monocrystalline oxide substrates that permit an epitaxial growth. Recent works have shown that the presence of a TiO 2 film between the substrate and the film favors in-situ cubic crystallization of SrVO 3 . [START_REF] Boileau | Buffer layer for the crystal growth of metal oxides of perovskite type in particular on amorphous substrates[END_REF]17 In this work, we report experimental investigations on the growth of transparent and conductive films by sputtering of Sr 2 V 2 O 7 target on Si substrates both with and without a TiO 2 seed layer. During this study, the obtained films have not been annealed. 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 resulting films without any post annealing process.

The aim of this study was to find suitable sputtering parameters in order to obtain optimal physical properties for specific functionalities in the microelectronic and optoelectronic industries.

Experimental details

The silicon substrates used in this study were single crystalline, (100) oriented and p-doped with a resistivity in the 5-15 Ω.cm range. Some of the Si substrates were coated with about 38 nm-thick of anatase TiO 2 (tetragonal structure with a = 3.79 Å and c = 9.51 Å) by atomic layer deposition (ALD). More details on this growth can be found elsewhere. [START_REF] Jolivet | Structural, optical, and electrical properties of TiO 2 thin films deposited by ALD: Impact of the substrate, the deposited thickness and the deposition temperature[END_REF] On these two types of substrates, thin films of SrVO 3 were grown by reactively sputtering 2" target of Sr 2 V 2 O 7 in a hydrogen rich plasma. The introduction of hydrogen allows to reduce the oxygen species in the plasma and therefore to control the oxygen incorporation in the layer as previously demonstrated in Si/SiO 2 multilayers. [START_REF] Gourbilleau | Sirich/SiO 2 nanostructured multilayers by reactive magnetron sputtering[END_REF][START_REF] Ternon | An original approach for the fabrication of Si/SiO 2 multilayers using reactive magnetron sputtering[END_REF] For this study, the radio frequency sputtering power applied on the target was 45 W and the plasma pressure was set at 3 mTorr for all deposits. Three series of samples were prepared. For the first series, the hydrogen rate r H in the plasma was set to 50% while the substrate temperature T G was varied from 400°C to 550°C with a 50°C step in order to study the influence of this parameter on the film properties. For the second series, T G was fixed at 500°C and the hydrogen rate r H varying from 0 to 80%. Finally, the third series was the same as the second one, but with a T G fixed at 550°C. These two last series aim at studying the influence of hydrogen incorporation in the plasma on film characteristics.

The structural properties of the films were studied using grazing incidence X-ray diffraction (GIXRD). The diffractometer used was a D8 Bruker Discover, equipped with a Cu source (λ k α1 = 1, 5406Å). Nanoscale observations and crystallographic analyses were made by transmission electron microscopy (TEM) (Cold FEG double corrected JEOL ARM200 operated at 200 kV) using the DigitalMicrograph software by GATAN for the image processing. Chemical maps were performed owing to STEM EDX data. The EDX spectrometer was a CENTURIO equipment from JEOL. For TEM observations, the thin films were prepared by focused ion beam (FIB) (Dual beam FEI Helios nanolab 660) method such that cross sectional observations of the films were following a [110] direction of the Si substrate were possible. The surface topographies of our films were investigated by atomic force microscopy (AFM) in tapping mode using a Nanoscope III device from Digital Instruments. The AFM images were corrected for flatness using the Nanoscope software and the RMS roughness was estimated from AFM digitized images with a scan area of 4 µm 2 .

The resistivity at room temperature was measured by the four probe method with a keithley 2450 sourcemeter. Moreover, the resistivities of the most interesting samples were measured from 5 to 300 K using the Van der Pauw configuration with a physical properties measurement system (PPMS) by Quantum Design. Hall measurements were carried out to determine the carrier concentration and their mobility using a magnetic field which varied from -9 to +9 Tesla. It is important to note that all the samples used for electrical experiments were grown on substrates containing an additional 8 µm thick thermal SiO 2 layer between the Si substrate and the top layer in order to eliminate the contribution of the conductivity of the substrate during measurements.

At last, optical measurements were carried out by using spectroscopic ellipsometry with an incident light energy which varied from 1.5 to 6 eV and under incident angles of 65, 70, and 75°with respect to the normal incidence (UVISEL 2 Jobin-Yvon). Once the measurements acquired, the software DeltaPsi2 allowed us to simulate the data and obtain the thickness and optical properties of the films.

Results and discussion

Structural properties

Figure 1 displays the GIXRD patterns of as grown SVO films resulting from the sputtering of Sr 2 V 2 O 7 target onto (100) Si substrates with and without a TiO 2 buffer layer. Figure 1a shows the GIXRD patterns of SVO films grown at a substrate temperature T G of 500°C and with a hydrogen rate r H of 50%, on both substrates. For the sample without TiO 2 , no diffraction peak can be observed in the range of 30 to 50°, whereas three peaks characteristics of crystallized SrVO 3 can be observed when TiO 2 anatase is present. This demonstrates that the TiO 2 buffer layer favors in-situ crystallization of SrVO 3 . The use of a TiO 2 buffer layer seems to be essential in order to obtain crystallized SVO layers on low-cost substrate without annealing. This better in-situ crystallization is explained by the weak mismatch between the two materials. [START_REF] Boileau | Buffer layer for the crystal growth of metal oxides of perovskite type in particular on amorphous substrates[END_REF] This demonstrates that the TiO 2 buffer layer favors in-situ crystallization of SrVO 3 . Figure 1b highlights the increasing crystallization of the films thanks to the rise of T G . From T G = 450°C, peaks around 32.6°and then for higher temperature also around 46.1°, which correspond to the (110) and (200) lattice planes of cubic SrVO 3 , respectively, can be observed, whereas they are not present for T G = 400°C. This result emphasizes the importance of the growth temperature for the achievement of crystallised cubic SrVO 3 . However it should be noted that the angles obtained are slightly lower than those expected. Indeed, in the literature, these peaks are found at 33.0°and 47.3°, respectively. This shift to lower angles could be explained by the stress exerted by the anatase TiO 2 buffer layer on the lattice of the SrVO 3 film. Indeed, anatase TiO 2 has a lattice parameter of 3.79 Å whereas that of bulk cubic SrVO 3 is 3.84 Å leading to a mismatch of -1.33%. Thus such a buffer layer creates compressive stress on the growing film which results in greater out-of-plane lattice parameters for the SrVO 3 thin films. Another explanation for this behaviour on the GIXRD pattern, for SVO films, could be a non-perfect stoichiometry, which can also lead to larger lattice parameters as repoted in literature. [START_REF] Moyer | Highly Conductive SrVO 3 as a Bottom Electrode for Functional Perovskite Oxides[END_REF][START_REF] Boileau | Tuning of the Optical Properties of the Transparent Conducting Oxide SrVO 3 by Electronic Correlations[END_REF] The increase of T G also causes the appearance of a peak around 42.2°and of another large peak with a lower intensity around 38.1°. These peaks are characteristic of the Sr 3 V 2 O 8 phase. Note that Sr 3 V 2 O 8 has also a peak at 40.5°, which corresponds to the SrVO 3 (111). Figure 1c presents the influence of r H on the crystallization of the films for T G = 500°C. We note that no peaks corresponding to crystallized SrVO 3 and Sr 3 V 2 O 8 are detected under a value of r H of 50%. Such a result indicates that if there is not enough H 2 during the growth, the amount of oxygen originating from Sr 2 V 2 O 7 will be too high to allow the formation of crystallized phase of SrVO 3 . This is confirmed by the XRD spectra achieved with higher hydrogen rate ranging from 50% to 80% and reported in Figure 1d in which one can observe the characteristic peaks of SrVO 3 and Sr 3 V 2 O 8 for all three films.

Most of the GIXRD patterns present a peak around 48.1°that corresponds to the anatase TiO 2 (200) crystallographic plane and whose intensity decreases with the increase of T G and / or r H . This evolution with the deposition parameters could be explained by two coupled effects. The first one is an etching process due to the presence of H-species in the plasma which reduces the volume of the crystalline buffer layer. The other one is the crystallization of the SVO layer, which results in a higher density and therefore reduces the penetration of the X-rays in the film.

To complete our data on the structural properties of the films, TEM observations were carried out. Figure 2 shows an HRTEM image of a 35 nm thick SVO film grown at 500°C, with a hydrogen rate r H of 50%, on a (100) silicon substrate. Most of the film is amorphous which confirms the results provided by XRD. However, some small regions in the upper part 7 of the film are found to be crystalline with the presence of nanocrystalline grains about 10 nm wide as attested by the enlarged region. The corresponding Fast Fourier Transform (FFT) is consistent with a [010] projection of the SrVO 3 cubic structure (a = 3.84 Å). The low fraction of the crystalline material in the film explains the lack of SrVO 3 characteristic peaks in the XRD pattern of 1a. But this is a promising result that shows that under very specific parameters to be carefully defined, it could be possible to grow crystalline SrVO 3 film without the use of a crystallized buffer layer. However, in the following, only the experimental results for the films grown with a TiO 2 buffer layer are presented.

Figure 3 is a summary of TEM results obtained from a TiO 2 /SVO structure grown on (100) Si, at 550°C and r H = 60%. Figure 3a is a typical bright field TEM image showing the whole structure with the stacked layers. The thicknesses are about 28 nm for TiO 2 and SrVO 3 sublayers. This observation evidences that the TiO 2 sublayer has been submitted to the etching by the hydrogen species of the plasma since its original thinckness was estimated to be about 38 nm by ellipsometry measurements. Such process is well known and has been evidenced in Si/SiO 2 system by Ternon et al. [START_REF] Ternon | An original approach for the fabrication of Si/SiO 2 multilayers using reactive magnetron sputtering[END_REF] These authors have shown that the presence of hydrogen in the plasma favors an etching phenomena of the growing layer which is dependent on the level of the hydrogen rate in the plasma. We observe, in the same way as these authors, a decrease of the average deposition rate as the function of the hydrogen rate for growing layer. Note that in some part of the sample, the TiO 2 layer thickness has lost about 20 nm due probably to a selective crystallographic etching by the hydrogen species (not shown here). One can also notice the native thin SiO 2 layer (2-3 nm) between the Si substrate and the buffer layer. Figures 3(b Sr and V chemical maps of the whole structure. Interestingly, the V map clearly shows a contrast in the TiO 2 sublayer. This reveals a V diffusion from the SVO film toward the TiO 2 layer, which can also contribute to the loss of crystallinity of the TiO 2 layer evidenced on the XRD spectra (Figure 1). The V concentration detected in the TiO 2 layer is less than 1%. The same interdiffusion has been also observed for a growth of SrVO 3 on TiO 2 using PLD. 17

In order to study the effect of the hydrogen rate on the films, the surface roughness was investigated by AFM (Figure 4). These 3D images clearly show that the films grown with r H = 20% (Figure 4a) and r H = 50% (Figure 4b) at T G = 500°C present a homogeneous roughness over the entire surface of the films. For T G = 500°C and r H = 80% (Figure 4c) or T G = 550°C and r H = 50% (Figure 4d), the AFM scanning reveals the presence of several irregularly distributed bigger domains over the entire surface of the films. Figures 5a and5b show the average roughness obtained by AFM measurements, with images of 2.5 µm x 2.5 µm images, and the corresponding value deduced from ellipsometric modeling as a function of T G for r H = 50% and as a function of r H for a substrate temperature T G of 550°C. At first glance, one can notice the consistency of the evolution between the two curves of Figure 5a showing an increase of the roughness with the growth temperature. This result highlights that the higher the deposition temperature is, the greater the etching of the film. For T G = 550°C, we observe a slight evolution of the roughness up to r H = 70% followed by an important increase between r H = 70% and r H = 80% (Figure 5b). This trend is also oberved for a growth temperature value of 500°C. Here again, the evolutions of the roughnesses deduced from either AFM or ellipsometric measurements are similar demonstrating that the model carried out thanks to the ellipsometriy measurements is accurate. These investigations reveal that if a high hydrogen rate in the plasma is necessary to achieve an in-situ crystallization of the films, it is also leads to an important surface roughness of them and therefore may have consequences on their physical properties. Indeed, concerning the electrical properties, the diffuse scattering of conduction electrons of a rough surface results in an increase of the electrical resistivities of the films. [START_REF] Luo | Surface roughness and conductivity of thin Ag films[END_REF][START_REF] Timoshevskii | The influence of surface roughness on electrical conductance of thin Cu films: An ab initio study[END_REF][START_REF] Ketenoglu | Influence of surface roughness on the electrical conductivity of semiconducting thin films[END_REF]

Electrical properties of the films grown with a TiO 2 buffer layer

The resistivity values at room temperature obtained by the 4 probe method are reported in Figure 6. Figure 6a displays the evolution of the resistivity as a function of the growth temperature for r H = 50%. We can see that the resistivity is high and constant until 450°C before decreasing significantly when the deposition temperature reaches 500°C. The high resistivity value achieved for the lowest deposition temperature (about 23 × 10 -3 Ω.cm) is attributed to the mainly amorphous nature of the film whereas its decrease to 7.5 × 10 -3 Ω.cm is correlated with the appearance of diffraction peaks characteristic of cubic SrVO 3 (Figure 1b). By increasing the deposition temperature by 50°C, the resistivity achieved the value of about 3 × 10 -3 Ω.cm for an hydrogen rate in the plasma of 50%. This evolution of the electrical properties can therefore be explained by the crystallization of the deposited films with the growth temperature. For a given growth temperature, the same trend can be observed with the increased hydrogen rate (Figure 6b). Without hydrogen during the deposition, the resulting layer is amorphous and has the characteristics of an insulating layer as expected with a resistivity as high as 550 Ω.cm. But a rate of 20% of H 2 in the plasma allows to improve this value by 4 orders of magnitude to reach 5.2 × 10 -3 Ω.cm, demonstrating the reactive effect in the reduction of the sputtering of the Sr 2 V 2 O 7 insulating target. By continuously increasing the hydrogen rate in the plasma, we observe an improvement of the resistivity of the lowest resistivity around 1.1 × 10 -3 Ω.cm for r H = 80% at 550°C. This is to our knowledge the best resistivity value ever reported in the litterature for such a SrVO 3 layer deposited by sputtering on a low cost substrate and in a large scale approach. [START_REF] Jung | Structural and electrical properties of transparent conductor SrVO 3 thin films grown using radio frequency sputtering deposition[END_REF] Nevertheless, this value remains higher than the resistivities of monocrystalline SrVO 3 films grown on LSAT, due to the polycristalline nature and the non-optimal crystallization template of the films. To confirm the essential role of the TiO 2 buffer layer, 4 probes measurements have been performed on the layer grown directly on the Si substrate at T G = 500°C and with r H = 50%. The resistivity achived for this layer is ρ = 226.5 Ω.cm, whereas the same layer grown on the same substrate but with TiO 2 buffer layer presents a resistivity value of 9.2 × 10 -3 Ω.cm. This confirms again the essential role played by the TiO 2 buffer layer in the in-situ crystallization process of the SrVO 3 films and consequently on its physical properties.

The evolutions of the resistivities of the different samples as a function of temperature are presented in Figure 7 : Figure 7a for different T G at a hydrogen rate r H of 50% and in Figure 7b for different r H % values at a fixed T G of 550°C. These evolutions are not corresponding to a cubic SrVO 3 material which is supposed to have a metallic behaviour with its resistivity increasing with temperature. [START_REF] Boileau | Tuning of the Optical Properties of the Transparent Conducting Oxide SrVO 3 by Electronic Correlations[END_REF][START_REF] Mirjolet | Electron-Phonon Coupling and Electron-Phonon Scattering in SrVO 3[END_REF][START_REF] Inoue | Band-width control in a perovskite-type 3d 1 correlated metal Ca 1-x Sr x VO 3 . I. Evolution of the electronic properties and effective mass[END_REF] Indeed, Figures 7a and7b clearly show the resistivities of the films decreasing with the increasing temperature. In contrast to theoretical analyses based on density functional theory, which showed that cubic SrVO 3 can only, by nature, be metallic, such behaviour is typical of a semiconductor material. [START_REF] Varignon | Origin of band gaps in 3d perovskite oxides[END_REF] Considering that Sr 2 V 2 O 7 is an insulating material and SrVO 3 is a conductive one, we can hypothesize that the films obtained are over-oxidized SrVO 3 or under-oxidized Sr 2 V 2 O 7 . This means the stoichiometry of these films might actually be SrVO 3+x with x comprised between 0 and 0.5, which could confirm the interpretation for the shift of the characteristic peaks of SrVO 3 on GIXRD pattern.

Another explanation might be the presence of the Sr 3 V 2 O 8 phase observed by GIXRD, which could modify the electrical properties of the films. In the Rath et al paper, SVO films grown by PLD on glass/TiO 2 showed a metallic behavior. 17 This means that the semiconducting behavior observed here is not caused by the TiO 2 buffer layer but by the specific conditions of the sputtering growth process. Note that to our knowledge, no data have been reported on the physical properties of such a material grown with this process in the literature. The films grown with optimized values of r H and T G in this study are more conductive than undoped thin films of semiconductors used as TCOs, such as ZnO [START_REF] Lai | Investigation of optical and electrical properties of ZnO thin films[END_REF][START_REF] Lu | Effect of RF power on optical and electrical properties of ZnO thin film by magnetron sputtering[END_REF][START_REF] Cho | Effects of Growth Temperature on the Properties of ZnO Thin Films Grown by Radio-frequency Magnetron Sputtering[END_REF] and SnO 2 [START_REF] Brousseau | Electrical properties and topography of SnO 2 thin films prepared by reactive sputtering[END_REF][START_REF] Kim | Effects of hydrogen plasma treatment on the structural and electrical properties of sputter-deposited SnO 2 thin films[END_REF][START_REF] Kim | Structural, electrical, and optical properties of reactively sputtered SnO 2 thin films[END_REF] with resistivity values of 4.3 × 10 -3 Ω.cm and 4.5 × 10 -3 Ω.cm, respectively. The obtained layers are also more conductive than thin films of Al doped ZnO before annealing for which the electrical resistivity is around 3.3 × 10 -3 Ω.cm. [START_REF] Kim | Optical and electronic properties of post-annealed ZnO:Al thin films[END_REF][START_REF] Youssef | Structural, optical and electrical properties of ZnO:Al thin films for optoelectronic applications[END_REF][START_REF] Lennon | Effects of Annealing in a Partially Reducing Atmosphere on Sputtered Al-Doped ZnO Thin Films[END_REF] Consequently, our vanadate films with a resistivity between 1 and 2 × 10 -3 Ω.cm offer potential applications in the microelectronics field since our approach is fully compatible with the standard of the microelectronic industry.

In Figure 8, the charge carriers charateristics at room temperature as a function of the growth parameters are presented. Figure 8a displays the evolution of the carrier concentration (orange) and the carrier mobility (purple) as a function of the growth temperature for a hydrogen rate of 50%. We can observe a spectacular increase of the carrier concentration as a function of T G . Indeed, the carrier concentration is around 5 × 10 19 cm -3 for T G = 400°C and it increases by approximately 2 orders of magnitude for T G = 550°C. We can also notice a decrease of the carrier mobility with T G with a mobility 6 times lower at 550°C than at 400°C. These evolutions are explained by the crystallization of the layers, as evidenced on the GIXRD patterns (Figure 1b). The crystallization of SVO leads to the correlated conduction regime characterized by high charge densities and low mobilities related to the high effective mass. In Figure 8b, the evolutions of the carrier concentration and of the mobility with r H from 50 to 80%, for a given growth temperature of 550°C, are shown. We can note an increase of the carrier concentration in this range until a maximum value is reached around 1.45 × 10 22 cm -3 . This value is consistent, but slighly lower, when compared to the values found in the literature for cubic SrVO 3 grown on monocrystalline substrates. [START_REF] Zhang | Correlated metals as transparent conductors[END_REF][START_REF] Boileau | Tuning of the Optical Properties of the Transparent Conducting Oxide SrVO 3 by Electronic Correlations[END_REF][START_REF] Mirjolet | Independent Tuning of Optical Transparency Window and Electrical Properties of Epitaxial SrVO 3 Thin Films by Substrate Mismatch[END_REF][START_REF] Jung | Structural and electrical properties of transparent conductor SrVO 3 thin films grown using radio frequency sputtering deposition[END_REF] This result demonstrates the importance of the hydrogen rate level to reach a carrier concentration value characteristic of crystallized SrVO 3 . The fact that the carrier concentration is low under non-optimal growth conditions underlines that the oxygen-stoichiometry is likely too high. This leads to an overoxidation to the V 5+ state, and therefore low charge concentrations. In terms of carrier mobility, we note a decrease between the r H values of 50 and 60% and after 60% an almost constant value of approximately 0.47 cm 2 .V -1 .s -1 .The films obtained in this study are polycristalline which entails carrier scattering by the grain boundaries and limits their mobility. This explains why our obtained values are lower than those found in the literature for monocristalline SrVO 3 thin films, situated between 1 and 2 cm 2 .V -1 .s -1 . [START_REF] Zhang | Correlated metals as transparent conductors[END_REF][START_REF] Boileau | Tuning of the Optical Properties of the Transparent Conducting Oxide SrVO 3 by Electronic Correlations[END_REF][START_REF] Mirjolet | Independent Tuning of Optical Transparency Window and Electrical Properties of Epitaxial SrVO 3 Thin Films by Substrate Mismatch[END_REF][START_REF] Jung | Structural and electrical properties of transparent conductor SrVO 3 thin films grown using radio frequency sputtering deposition[END_REF] Another reason is that our layers are rougher than those of literature.

Therefore, as explained earlier, the charge carriers of our films can also be diffused by their rough surfaces. However, with the data presented in this article we cannot conclude on the predominance of one of the two phenomena.

Based on these findings, the evolution of the carrier mobility as a function of the measurement temperature has been investigated for layers deposited at 550°C with a 60 and 70%

hydrogen rate Figure 9a. The evolution unveils the existence of a low and of a high temperature regime between 50 and 300 K with an increase of the mobility between 50 and 200 K followed by its decrease between 200 and 300 K. In the literature, theoretical investigations have been reported aiming at understanding the electron scattering by ionized impurities in semiconductors. [START_REF] Conwell | Theory of Impurity Scattering in Semiconductors[END_REF][START_REF] Chattopadhyay | Electron scattering by ionized impurities in semiconductors[END_REF] They have shown that, at low temperature, the mobility follows a T 3/2 law due to the charge carrier diffusion. Whereas, other articles have reported that, the mobility follows a T -3/2 law in the high temperature regime caused by the diffusion of the carriers by the acoustic phonons. [START_REF] Tang | The role of acoustic phonon scattering in charge transport in organic semiconductors: A first-principles deformation-potential study[END_REF][START_REF] Bardeen | Deformation Potentials and Mobilities in Non-Polar Crystals[END_REF][START_REF] Seitz | On the Mobility of Electrons in Pure Non-Polar Insulators[END_REF] In Figure 9a, dashed lines following T 3/2 and T -3/2 are provided as a guide to the eye, in the low and high temperature ranges, to evidence the good agreement between our measurements and the theoretical predictions, which confirms the semiconductor behaviour of the films. Concerning the evolution of the charge carrier concentration with temperature, a high and a low temperature regime is also observed (Figure 9b).

For the low temperature regime we note an atypical decrease of the carrier concentration as a function of the temperature as already observed for several different semiconductors such as GaAs, In 2 O 3 , In 2 O 3 -ZnO and Mg 3 Sb 2 , and attributed either to the presence of impurities or vacancies. [START_REF] Kunets | Isotropic Hall effect and "freeze-in" of carriers in the InGaAs self-assembled quantum wires[END_REF][START_REF] Bierwagen | High electron mobility In 2 O 3 (001) and (111) thin films with non degenerate electron concentration[END_REF][START_REF] Makise | Activation like behaviour on the temperature dependence of the carrier density in In 2 O 3 -ZnO films[END_REF][START_REF] Mao | Anomalous electrical conductivity of n-type Te-doped Mg 3.2 Sb 1.5 Bi 0.5[END_REF] We can hypothesize that a similar mechanism occurs in our films which explains the observed decrease of carrier concentration with the measurement temperature.

By contrast, the charge carrier concentration increases as a function of the temperature in the high temperature regime. Therefore, it appears that the amount of charge carriers in the conduction band increases thanks to their increased thermal energy. It is possible to

show that, when charge carriers pass from a donor level towards the conduction band, the evolution of the conductivity (σ) as a function of the temperature follows [START_REF] Weller | An introduction to principles of the solid state. Extrinsic semiconductors[END_REF] :

ln(σ) = ln(A) - ∆E a 2k B T (1)
Where A is a constant, ∆E a an activation energy and k B the Boltzmann constant. Figure 9c presents the evolution of the plot of ln(σ) as a function of 1000 T , in the high temperature regime, for samples grown at 550°C with a hydrogen rate of 60 and 70%. We can observe that, for these two samples, the linear adjustments using the equation 2 fit well the experimental data well, which confirms our hypothesis about the increase of the number of charge carriers in the conduction band at high temperatures. Activation energies of approximately 18.7 meV and 12.5 meV were obtained for the samples grown with r H = 60% and r H = 70% respectively. However, another model is currently used to describe carrier transport in disordered semiconductors. This model is called Mott's variable range hopping (MVRH) and in this model, the electrical conductivity evolves according to the temperature as [START_REF] Chanda | Magnetic ordering and conduction mechanism of different electroactive regions in Lu 2 NiMnO 6[END_REF][START_REF] Hill | Variable-range hopping[END_REF][START_REF] Rosenbaum | Crossover from Mott to Efros-Shklovskii variable-range-hopping conductivity in In x O y films[END_REF][START_REF] Rudra | Existence of nearest-neighbor and variable range hopping in Pr 2 ZnMnO 6 oxygen-intercalated pseudocapacitor electrode[END_REF] :

ln(σ) = ln(B) -     T m T     1/4 (2)
Where B is a constant and T m the Mott temperature. Figure 9d presents the evolution of the plot of ln(σ) as the function of 1000 T 1/4 , in the high temperature regime, for exactly the same samples as those of Figure 9c. We can observe that, here again, the linear adjustments fit the experiment data well. Considering that these two models fit the experiment data with R 2 = 0.999 for the two samples, we cannot conclude about the conductivity mechanism which takes place in the obtained layers with the data presented in this article.

Optical properties of the films grown with a TiO 2 buffer layer

The electrical characteristics of the films have been detailed in the previous part, and in order to complete the study of their TCO properties, we will now study their optical properties. Figure 10 presents the evolution of the refractive index at 1.95 eV as a function of the hydrogen rate for a given growth temperature of 550°C. In this figure, we can observe that, without hydrogen during the deposition, the resulting SrVO 3 layer has an optical index of 2.06, decreasing as a function of r H until it reaches a value of 1.85 for the film grown with 80% of hydrogen. Theoretical analyses based on density functional theory have shown that the optical index at 1.95 eV of Sr 2 V 2 O 7 is around 2.10, [START_REF] Schira | DFT and Hybrid-DFT calculations on the electronicproperties of vanadate materials: theory meets experiments[END_REF] and Roth et al in their study on the growth of cubic SrVO 3 by MBE on LSAT substrates have reported a value of 1.75. [START_REF] Roth | Self-regulated growth of [111]-oriented perovskite oxide films using hybrid molecular beam epitaxy[END_REF] Therefore, Figure 10 highlights again the importance of the amount of hydrogen in order to reduce the oxygen of the Sr 2 V 2 O 7 target during the growth process. This result is consistent with the electrical measurements since, as demonstrated earlier, the layer deposited without hydrogen is insulating while the one grown with 80% of hydrogen has a carrier concentration close to the values presented in the literature for SrVO 3 .

The real and imaginary parts of the dielectric permittivity ( = 1 -i 2 ) were extracted from the spectroscopic ellipsometry data. Figures 11a and11b present, respectively, the evolution of 1 and 2 in the range of 0.6 to 6.5 eV for different r H values and a growth temperature of 550°C. These two figures show the presence of peaks for 1 and 2 and we can observe that a shift of these peaks takes place for a r H value between 50 and 60%. Indeed, 1 reaches a maximum at 4.20 eV for the layer grown with a hydrogen rate of 50% whereas, for the other films deposited with higher r H , this maximum is achieved for 3.80 eV. For 2 , the peak is shifted from 4.80 to 4.50 eV. However, the shape of the permittivity curves for the obtained films does not match with those reported in the literature. [START_REF] Zhang | Correlated metals as transparent conductors[END_REF][START_REF] Mirjolet | Independent Tuning of Optical Transparency Window and Electrical Properties of Epitaxial SrVO 3 Thin Films by Substrate Mismatch[END_REF][START_REF] Mirjolet | Electron-Phonon Coupling and Electron-Phonon Scattering in SrVO 3[END_REF][START_REF] Boileau | Highly Transparent and Conductive Indium-Free Vanadates Crystallized at Reduced Temperature on Glass Using a 2D Transparent Nanosheet Seed Layer[END_REF] Indeed, contrary to what we observe, 1 is usually negative for energies between 0.6 and 1.3 eV and is supposed to go through 0 before reaching a maximum at a value of 3.10 eV, while 2 should present a maximum around 3.30 eV. Mirjolet et al explained that for SrVO 3 , the screened plasma energy E ω * p can be determined thanks to the relation 1 (E ω * p ) = 0. 8 However, this approach only works when SVO is a strongly correlated metal for which the Drude model accurately describes the contribution of the free electrons to the permitivity at low energy, but does not work for semiconductors since 1 does not go through 0. In conclusion, in agreement with the electrical measurements, the permittivity spectra confirm that our films lack the properties of strongly correlated metals.

In order to study the optical absorption of the light by our samples, Tauc plots were depicted by plotting (αhν) x as the function of hν where α is the absorption coefficient (extracted from the ellipsometric models) and hν the incident photon energy (Figure 12a). For a direct optical transition, x is equal to 2 whereas it is equal to [START_REF] Zhang | Correlated metals as transparent conductors[END_REF] 2 in the case of an indirect transition. [START_REF] Tauc | Optical Properties and Electronic Structure of Amorphous Semiconductors[END_REF][START_REF] Ghobadi | Band gap determination using absorption spectrum fitting procedure[END_REF] We have therefore plotted two graphs for the layer grown with a hydrogen rate of 50% at 550°C in Figure 12a. The dashed orange curve represents the case of an indirect optical transition and the dashed purple curve represents the case of a direct optical transition. These two curves present linear parts and we can consequently determine the energy bandgap for a direct and an indirect absorption (orange and purple lines). Bandgap values between 4.11 eV and 3.80 eV were obtained for direct transitions. However, considering that the visible range is between 1.55 and 3.25 eV, this direct transition will not have any impact on the TCO properties of our films. The evolution of the energy bandgap for an indirect transition as a function of r H , for a given growth temperature of 550°C, is shown in Figure 12b. A decrease of this indirect bandgap as a function of the hydrogen rate can be noted. Indeed, for a r H of 50% this energy is equal to 2.97 eV while we obtained a value of around 2.72 eV for a r H of 80%. In literature, a value of 2.70 eV is found for the indirect transition of SrVO 3 and is associated to an excitation of an electron from the oxygen 2p band to the empty V 3d-t 2g band. [START_REF] Mirjolet | Optical Plasmon Excitation in Transparent Conducting SrNbO 3 and SrVO 3 Thin Films[END_REF][START_REF] Mohammadi | Tailoring Optical Properties in Transparent Highly Conducting Perovskites by Cationic Substitution[END_REF] We can therefore conclude once again that, the higher the hydrogen rate is during the growth process, the more the resulting layer has physical properties which tend towards those of SrVO 3 .

To determine the optical transparency of our films, we used the following formula :

T = (1 -R -dα) × 100 (3)
Where d and α are respectively the thickness layer and the absorption coefficient extracted from the model obtained thanks to the ellipsometry measurements. The value of the reflectivity (R) for a perpendicular light incidence has been calculated using the relation :

R =     N 2 -N 1 N 2 + N 1     2 (4)
Where N 2 is the optical index of the vanadate layer, also extracted from the ellipsometric model, and N 1 is the optical index of air (N 1 = 1). Figure 13a presents the evolution of the optical transparency for different values of T G with r H = 50%. We note a sharp increase of the optical transparency in the visible range starting from 500°C. This increase of the optical transmittance in the visible range is concomitant with the emergence of the peaks characteristic of cubic SrVO 3 in the GIXRD patterns (Figure 1b). The growth temperature is therefore a key parameter in order to improve the optical transparency of the samples in the visible range, thanks to the improved crystallization of the films. Figure 13b presents the evolution of the transmittance for different values of r H with a growth temperature of 550°C. The layer formed without hydrogen is the sample with the best transparency, and as r H increases, we see a drop in transmittance. Stated differently, the optical transmittance decreases with increasing electrical conductivity of the resulting layer (Figure 6b). Using

Hall measurements, we have demonstrated that the carrier concentration increases with the hydrogen rate (Figure 8b). Futhermore, as shown by the Tauc plots (Figure 12b), films produced with a high hydrogen content exhibit an indirect bandgap transition in the visible region. Thus, the decrease of the transmittance can be explained physically by the fact that the visible light can be absorbed and reflected by a layer's charge carriers to a greater extent as it grows with a significant amount of hydrogen. Gu et al. had previously noted in their literature the difference in optical transparency between the conductive and insulating vanadates layers. [START_REF] Gu | Redox controllable switch of crystalline phase and physical property in SrVO x epitaxial films[END_REF] These two physical behaviours can therefore explain the decrease of the transmittance as a function of r H . We note that the obtained films are less transparent than thin films of semiconductors used as TCOs, such as ZnO and SnO 2 which present optical transmittance of about 85% in the visible range. [START_REF] Lai | Investigation of optical and electrical properties of ZnO thin films[END_REF][START_REF] Kim | Structural, electrical, and optical properties of reactively sputtered SnO 2 thin films[END_REF] However, even the layer grown with a hydrogen rate of 80% has a minimum transparency higher than the maximum transparency of some SVO films reported in the literature (63%), 1,10 which were grown by PLD and have a thickness of around 45 nm.

Conclusion

In conclusion, the growth of vanadate films on Si both with and without a TiO 2 seed layer by magnetron sputtering using sputtering of a Sr 2 V 2 O 7 target was carried out successfully.

Their structural, electrical and optical properties were examined. The first conclusion that could be drawn from these measurements is that the TiO 2 seed layer is necessary to obtain a fully crystallised SVO film with TCO physical properties. Secondly, even with the seed layer, T G and r H are both crucial growth parameters to consider in order to optimize the electrical and optical properties of the obtained films. These properties were optimized for a growth temperature of 550°C and an important amount of hydrogen in order to have a conductive film. Thirdly, the optical properties of SVO films grown on low cost substrates by magnetron sputtering were presented, which complements the existing literature. Finally, with these optimized growth parameters, the vanadate films are semiconductors which possess a higher electrical conductivity than those of the semiconductors currently used as TCOs and with optical transparency higher than 63% in all the visible range.

Several options can be considered to further improve the physical properties of these films. Heat treatments after deposition could be attempted to favor the formation of the crystalline phase and then to enhance the optical and electrical properties. In-situ annealing under vaccum or a reducing atmosphere could also be carried out in order to avoid the oxidation of the films (and the appearance of Sr 2 V 2 O 7 ) and a potential decrease of the electrical conductivity. Furthermore, a passivation of the films under H 2 could heal damaged atomic bonds and thus increase the electrical conductivity, as it is the case with Si-based microelectronic devices. At last, the doping of these films with luminescent ions could also be attempted in order to obtain luminescent materials for optoelectronic functionalities.

A study by sputtering of an SrVO 3 target could be also interesting in order to reduce the H 2 amount in the plasma during the growth and obtain less rough films.

  ) and 3(c) are HRTEM images of enlarged regions of TiO 2 and SrVO 3 , respectively. The FFT pattern in the inset of Figure 3c is in agreement with the cubic structure of SrVO 3 along the [010] direction. At last, Figure 3d depicts Ti,

Figure 1 :

 1 Figure 1: GIXRD patterns of thin film grown with sputtering of Sr 2 V 2 O 7 target a) with and without a TiO 2 seed layer b) for different values T G and r H = 50%, c) for different values of r H and T G = 500°C, d) for different values of r H and T G = 550°C.

Figure 2 :

 2 Figure 2: Typical HRTEM image of a SrVO 3 film grown directly on a (100) Si substrate, at 500°C with r H = 50%. SrVO 3 small grains are cleary visible in the film.

Figure 3 :

 3 Figure 3: (a) Bright field TEM image of a TiO 2 /SVO film grown on (100) Si. TiO 2 was grown by ALD and SVO by sputtering at 550°C and using r H = 60%. (b) HRTEM image of a TiO 2 region. The FFT in inset is consistent with the anatase structure projected along the [-1-11] direction. (c) HRTEM image of a SVO region with the corresponding FFT in inset. This latter is in agreement with a [010] direction of SrVO 3 . (d) STEM HAADF image and the corresponding STEM EDX chemical maps of Ti, Sr and V.

Figure 4 :

 4 Figure 4: AFM images of the films obtained by sputtering of Sr 2 V 2 O 7 target on Si substrates and anatase TiO 2 seed layer. a) for T G = 500°C and r H = 20%, b) for T G = 500°C and r H = 50%, c) for T G = 500°C and r H = 80%, d) for T G = 550°C and r H = 50%. AFM images are 2 µm x 2 µm

Figure 5 :

 5 Figure 5: Evolution of the RMS roughness and the roughness derived from ellipsometric model as a function of a) T G with r H = 50%, b) r H for T G = 550°C.

Figure 6 :

 6 Figure 6: Electrical resistivities at room temperature obtained by the 4 probe method. a) for different values of T G with r H = 50%, b) for different values of r H with T G = 550°C.

Figure 7 :

 7 Figure 7: Electrical resistivities as a function of the temperature, a) for different values of T G with r H = 50%, b) for different values of r H with T G = 550°C.

Figure 8 :Figure 9 :

 89 Figure 8: Transport properties at 300K. a) carrier concentration evolution at room temperature for different values of T G and r H , b) carrier mobility at room temperature for different values T G and r H

Figure 10 :

 10 Figure 10: Evolution of the optical index for an incident photon energy of 1.95 eV as a function of r H for T G = 550°C.

Figure 11 :

 11 Figure 11: a) Evolution of the real part of the dielectric permittivity as a function of the energy for different values of r H with T G = 550°C, b) Evolution of the imaginary part of the dielectric permittivity as a function of the energy for different values of r H with T G = 550°C.

Figure 12 :

 12 Figure 12: a) Tauc plots of SVO film grown at 550°C with 50% of hydrogen rate, plotted for direct (purple) and indirect (orange) optical transitions, b) Energy bandgap for an indirect transition as a function of r H with T G = 550°C.

Figure 13 :

 13 Figure 13: Evolution of transmittance as a function of the wavelengh for, a) different values of T G with r H = 50%, b) different values of r H with T G = 550°C.

Acknowledgement

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.

Xavier Portier thanks the ANR (Agence Nationale de la Recherche) institution and the Normandy Region for their financial support in acquiring the FIB setup (ANR-11-EQPX-0020).

The authors thank Cédric Frilay for the preparations of the TiO 2 seed layer by ALD and Franck Lemarié to the AFM mesurements and the FIB preparations.