# Single-Phase Growth, Stabilization, and Electrical Properties of B Phase VO₂ Films Grown on Mica by Reactive Magnetron Sputtering

E. Ekström,* S. Hurand, M. M. Yildizhan, A. Elsukova, P. O. Å. Persson, B. Paul, G. Ramanath, A. le Febvrier,* F. Eriksson,* and P. Eklund*

The VO₂ metastable (B) phase is of interest for applications in temperature sensing, bolometry, and Li-ion batteries. However, single-phase growth of thin films of this metastable phase is a challenge because vanadium oxide exhibits many polymorphs and the VO₂ stable (M1) phase is usually present as a secondary phase. Additionally, the phase transition at 350 °C in the (B) phase severely narrows the processing window for achieving phase-pure films. Here, single-phase growth of 5-to 50-nm thick VO₂ (B) films on muscovite, mica, by pulsed direct-current reactive magnetron sputtering at 400 °C is demonstrated. The films are phase-pure and exhibit a high density of 4.05 g cm⁻³ and low resistivity of about 50 mΩcm at 30 °C. Increasing the film thickness to 100 nm results in a V₂O₅-capped VO₂ (B) film with a resistivity of 8000 mΩcm. These results indicate that the stability of the VO₂ (B) phase is sensitive to in situ annealing during deposition. These findings should serve as a basis to design processes to exclusively obtain phase-pure VO₂ (B) films.

The presence of numerous phases (e.g., VO, V₂O₃, VO₂ and V₂O₅) and their polymorphs make vanadium oxide attractive for

E. Ekström, M. M. Yildizhan, A. Elsukova, P. O. Å. Persson, B. Paul, A. le Febvrier, F. Eriksson, P. Eklund

Thin Film Physics Division

Department of Physics

Chemistry and Biology (IFM)

Linköping University

Linköping SE-581 83, Sweden

E-mail: erik.ekstrom@ri.se; arnaud.le.febvrier@liu.se;

fredrik.erksson@liu.se; per.eklund@liu.se

S. Hurand

Institut Prime

Université de Poitiers

11 Boulevard Marie et Pierre Curie, Chasseneuil du Poitou 86360, France

G. Ramanath

Department of Materials Science and Engineering

Rensselaer Polytechnic Institute

Troy, NY 12180, USA

<figure><img src="image_1.png" /></figure>

ID The ORCID identification number(s) for the author(s) of this article can be found under https://doi.org/10.1002/apxr.202300032

© 2023 The Authors. Advanced Physics Research published by Wiley-VCH GmbH. This is an open access article under the terms of the Creative Commons Attribution License, which permits use, distribution and reproduction in any medium, provided the original work is properly cited.

DOI: 10.1002/apxr.202300032

a variety of applications including thermochromic coatings, batteries, and bolometry, to name a few.[1-9] For example, VO₂ has a multitude of polymorphs (e.g., (M1), (M2), (B), (R) and (A)) that can be stabilized by doping, stresses and other means.[10-13] The most studied polymorphs among these are the rutile (R) (a = b = 4.55 Å and c = 2.86 Å) and monoclinic (M1) (a = 5.38 Å, b = 4.52 Å, c = 5.74 Å and β = 122.6°) phase, which are the stable phases close to room temperature and exhibit a metal-to-insulator transition at 68 °C.[14] The (A) and (B) phase are also tetragonal (a = b = 8.43 Å and c = 7.68 Å) and monoclinic (a = 12.03 Å, b = 3.69 Å, c = 6.42 Å and β = 106.6°), respectively. It is known that small changes in growth parameters, such as temperature and

oxygen partial pressure, can be utilized for selectively grow different phases and polymorphs. The V-V distance seems to be the key parameter to which polymorph is formed.[15] Additionally, the choice of substrate plays a role in the growth mode and formed phases. We have previously shown that VO₂ (M1) phase can be grown by van der Waals epitaxy on mica.[16] In general, growth on mica may be by van der Waals epitaxy or conventional epitaxy, and it is thus important to prove all features of van der Waals epitaxy, rather than assuming it, as is sometimes done in the literature. There is also a recent report of a third type of epitaxy on mica, substrate-compliant epitaxy.[17] In contrast, the (B) phase can be grown by changing the partial pressure of oxygen, as will be shown in the present paper. Mica is also a flexible substrate, allowing for applications where flexible devices are needed.[18,19]

The (B) phase, which is of potential use in Li-ion batteries, has promising electrical properties and can detect small increases in temperatures. The latter is of importance in such diverse areas as bolometers, temperature management in microsystems and precise temperature sensing. It is also of potential use in optical sensors. While these features provide versatility in properties and applications, exclusive synthesis of desired phase(s) is a challenge. Hydrothermal methods allow growth of B-phase nanostructures such as nanorods, but phase-pure growth of epitaxial films or single crystals is challenging. For instance, although the metastable (B) phase (space group C2/m) can be stabilized by stress, its exclusive growth is constrained by its conversion to the stable rutile

<figure><img src="image_2.png" /></figure>
(R) phase at 350 °C.[20] While pulsed laser deposition has demonstrated the growth of the (B) phase at higher temperatures,[21,14] the (M1) phase is also usually present in these films.[22,23] Hence, there are continued efforts to devise new methods to grow phase-pure films of the VO₂ (B) phase at a sufficiently high temperature to obtain good film quality, without activating the B-to-R phase transition.

Here, we demonstrate single-phase growth of 5- to 50-nm thick films VO₂ (B) films on muscovite mica substrates by pulsed DC (direct current) reactive magnetron sputtering at 400 °C. These phase-pure films exhibit a low resistivity of 50 mΩcm at 30 °C. Increasing the film thickness to 100 nm results in a V₂O₅-capped VO₂ (B) film with a resistivity of 8000 mΩcm. These results indicate that the stability of B-phase VO₂ is sensitive to in situ annealing during deposition. Our findings should serve as a basis to design processes to exclusively obtain phase-pure VO₂ (B) films.

The VO₂ films were grown on freshly tape-cleaved muscovite mica (001) substrates obtained from Oxford Instruments using pulsed DC reactive magnetron sputtering. A 99.7% pure 2-inch vanadium target (Plasmaterials, Livermore, CA, USA) was sputtered using an 89:11 Ar-O₂ mixture at 0.53 Pa (4 mTorr). Prior to deposition, the substrates were heated to 400 °C, and held for 10 min for degassing at a 6.7 × 10⁻⁶ Pa (5 × 10⁻⁸ Torr) base pressure. The chamber is described elsewhere.[24] The substrate temperature was maintained at 400 °C throughout the deposition. For the depositions, a 160 W power source was pulsed at 50 Hz, at a 90% duty cycle, a 30 µs crowbar length, and a 10% reversed voltage with a floating substrate bias. Film thickness tfilm measured by x-ray reflectivity (XRR) [see Figure S1, Supporting Information] revealed that the planar growth rate of the films was 1 nm per min. Accordingly, the deposition time was varied from 5 to 100 mins to obtain films with 5 nm ≤ tfilm ≤ 100 nm.

X-ray diffractometry (XRD) was performed using a X"Pert PRO powder diffractometer with a copper anode source (Cu Kα, λ = 1.54 Å), operated at 45 kV and 40 mA. The incident beam path included a 0.5° divergence slit followed by Bragg-Brentano^HD optical module and a 0.5° anti-scatter slit. The diffracted beam path comprised of a 5 mm anti-scatter slit followed by a 0.04 rad Soller slit, and a Ni-filter. An X"Celerator detector operated in scanning line mode with a 2.122° active length was used to quantify the diffracted X-rays. XRR measurements to determine the film thickness and density were carried out in a PANalytical Empyrean diffractometer with a copper anode source (Cu Kα, λ = 1.54 Å), operated at 45 kV and 40 mA. The incident beam path consisted of a 1/32° divergence slit and a hybrid mirror, and in the diffracted beam path a triple axis Ge 220 analyzer was used together with a PIXcel3D detector operated in open detector mode.

A LEO Gemini 1550 Zeiss instrument was used for scanning electron microscopy (SEM) of the film surfaces. A Carl Zeiss crossbeam 1540 Focused Ion Beam system was used to prepare samples for transmission electron microscopy (TEM) for which a 200 kV FEI Tecnai G2 instrument or the Linköping monochromated, double-corrected 300 kV FEI Titan³ 60–300 (S)TEM were employed. Electron energy loss spectroscopy was carried out with a 0.2 eV energy resolution using a Gatan GIF Quantum ERS postcolumn energy filter.

Room-temperature electrical resistivity of the films were determined from four-point probe measurements in a Jandel RM3000 instrument together with XRR-determined film thicknesses. The

temperature dependence of the resistivity from 30 °C to 120 °C was measured in a four-probe van der Pauw configuration with an input current of 10 µA on the high-temperature module AHT55T5 of a Hall Measurement System HMS-5300 provided by the ECOPIA company.

SEM images from the as-deposited film with thickness tfilm = 5 nm and tfilm = 10 nm (see **Figure 1a**) shows no detectable topographical contrast from grains or surface features and suggest sub-nm surface roughness. Films with tfilm = 21 nm and tfilm = 50 nm reveal mild contrast indicating 100–200 nm grains (see Figure 1b). Films with tfilm = 100 nm (see Figure 1c) exhibit distinct contrast arising from randomly oriented faceted grains.

X-ray diffractograms (**Figure 2a**) reveal a strong 001 peak, whose width w₀₀₁ decreases with increasing film thickness (Figure 2b). Thickness fringes are visible around the 001 peak for tfilm = 21 nm and tfilm = 50 nm, indicating good crystal quality. For films with tfilm ≤ 50 nm, the 001 peak shifts to higher angles with increasing film thickness, attributable to the variation of film stress with film thickness.[25,26] Diffractograms from 100 nm thick films exhibit additional peaks from a secondary phase corresponding to the 001 and 102 Bragg reflections from the V₂O₅ structure.

Film density analysis from XRR data (see Figure S1, Supporting Information) indicates that the surface of the 100 nm thick film has a lower density of 3 g cm⁻³ as compared to 4 g cm⁻³ for the tfilm ≤ 50 nm films. We propose, given that V₂O₅ has a lower density than VO₂ (B) and the fact that 100 nm thick films exhibit underdense surface morphology with randomly oriented faceted grains, that the surface of the 100 nm thick films consists of the V₂O₅ phase. Thus, the higher density films with tfilm ≤ 50 nm consist exclusively of phase-pure VO₂ (B) films, consistent with featureless SEM images. These results suggest that at extended deposition times, the VO₂ (B) formation gives way to V₂O₅ formation, underscoring the delicate balance between kinetic, thermodynamic, and other factors such as possible strain, that determine phase selection in this system.

TEM micrographs from cross-sections of films with tfilm = 50 nm indicate a continuous and uniform film (see **Figure 3a**) with high crystallinity (Figure 3b) throughout the film thickness. The film thickness is determined to be ≈50 nm throughout the entire lamella. A lattice resolved image of the sample is illustrated in Figure 3c and a Fast Fourier Transform (FFT) obtained from the same region, provided in the inset, reveal the main reflections. The pattern is in good agreement with (11̄0) projection of VO₂–(B) [ICDD PDF 01-081-2392].

High angle annular dark field scanning TEM (HAADF-STEM) images from films with tfilm = 100 nm revealed a layered film microstructure (**Figures 4a, 5a**). The chemical structure of the film layers was identified using near edge fine structure of the electron energy loss (EEL) spectra[27] to chemically distinguish the VO₂ and V₂O₅ regions. This approach circumvents the complication of overlapping V L and O K edges in both EDX and EEL spectra for these phases. EEL spectra acquired from various film layers were compared with reference spectra from the vanadium oxide VO₂ and V₂O₅ phases.[14] The difference in near-edge fine structures of the two phases is manifested by the presence of a shoulder in the VO₂ spectrum (marked by a red arrow in Figure 4c), shapes of the oxygen O-K(P) and O-K(Q) peaks and their ratio to the vanadium L₂,₃ edges. The vanadium oxidation states were
<figure><img src="image_3.png" /><figcaption>**Figure 1.** SEM images from VO₂ films with a) $t_{film} = 5$ nm, b) $t_{film} = 50$ nm, and c) $t_{film} = 100$ nm.</figcaption></figure>

<figure><img src="image_4.png" /><figcaption>**Figure 2.** 2θ-θ X-ray diffractograms of the VO₂ films are shown in (a) and zoomed in diffractogram of the 001 film peak is shown in (b), highlighting the peak shift to higher diffraction angles for thicker films. Substrate peaks are marked by an asterisk, while indexed peaks correspond to the V₂O₅ structure.</figcaption></figure>

mapped by fitting experimental spectrum derivatives with a linear combination of the derivatives of VO₂ and V₂O₅ reference spectra (Figure 4d). The map shows that the bottom layer exhibits a VO₂ oxidation state, and the upper layers are composed either of V₂O₅ or a mix of oxides. Another area was analyzed in the same way and shows similar results, see Figure S3 (Supporting Information) for more details.

High resolution HAADF-STEM imaging was used to investigate the crystal structure of the 100 nm film (Figure 5a-c) FFT analysis confirmed the presence of V₂O₅ phase in a middle layer

<figure><img src="image_5.png" /><figcaption>**Figure 3.** Cross-section TEM micrographs from a film with $t_{film} = 50$ nm. a) Low magnification overview, b) high magnification of the film thickness and c) lattice resolved micrograph of the same film. The FFT obtained from (c) is shown in the inset.</figcaption></figure>

of the film (Figure 5b and insets). However, the FFT pattern of the bottom image layer could not be unequivocally matched to a VO₂ (B) phase (Figure 4c and inset). To further analyze the structure of the bottom layer, we obtained a selected area diffraction pattern calibrated using an oriented gold thin film (Figure 5d,e). Experimental diffraction pattern was matched to all possible vanadium oxide phases. None of the phases matched the obtained diffraction precisely, with phases (B) and (M2) providing the closest matches. In Figure 5e, simulated corresponding diffraction patterns are overlayed with diffraction image. The experimental diffraction spots lie between two reflections from two phases (Figure 5e, inset). This result indicates that during the film growth the bottom layer undergoes an unfinished transition from (B) to (M2) phase.

The electrical resistivity of the VO₂ (B) films does not vary appreciably at ≈50 mΩcm for film thicknesses in the 5 nm ≤ $t_{film}$ ≤ 50 nm range. In comparison, literature reports resistivity values in a wide range, 4 mΩcm – 500 mΩcm.[6,15,16] The film with $t_{film} = 100$ nm that contained both the VO₂ and V₂O₅ phases exhibited a resistivity of 8000 mΩcm, showing the strong influence of the high resistivity V₂O₅ phase. The resistivity of films with 5 nm ≤ $t_{film}$ ≤ 50 nm decreases with temperature, typical of semimetal behavior expected for the VO₂ (B) around room temperature.[15,16] There are no observable sharp changes in the resistivity between
<figure><img src="image_6.png" /><figcaption>Figure 4. a) HAADF-STEM image of three-layer vanadium oxide thin film. b,c) EEL spectra (green area) obtained from points 1 and 2 marked in image (a) plotted together with reference VO₂ and V₂O₅ spectra (red and blue lines, respectively). d) HAADF-STEM image of three-layer vanadium oxide thin film with corresponding map of vanadium oxidation state.</figcaption></figure>

<figure><img src="image_7.png" /><figcaption>Figure 5. a–c) HAADF-STEM images from a 100 nm thick film. Insets in (b) show high resolution lattice images and a FFT from V₂O₅ in the middle layer, while the inset in (c) indicates the FFT from the VO₂ (B) layer closer to the substrate. d) Overview bright field TEM image of a 100 nm thick film with e) a corresponding selected area diffraction pattern obtained from the area marked with red circle. Simulated diffraction patterns for (B) and (M2) vanadium oxide phases (red and blue dots, respectively) are overlaid on experimental image.</figcaption></figure>

30 °C and 120 °C, indicating unlikelihood of presence of any significant amounts of VO₂ (M1), known to exhibit a metal-insulator transition[17] at 68 °C. These results again are consistent with our XRD and TEM results.

The temperature coefficient of resistance (TCR), α, of the films are calculated by using Equation (1),

$$ \frac{d\rho}{\rho dT} = \alpha \qquad (1) $$

where ρ and T are resistivity and temperature, respectively. The films with thicknesses 10 nm to 50 nm have similar TCR values of −1.5% K⁻¹ at 30°C to −1.0% K⁻¹ at 120°C. The 5 nm thick film follows the same trend with an overall 0.5 percentage points higher (−2.0% K⁻¹ at 30°C to −1.0% K⁻¹ at 120°C). These values are slightly lower at low temperatures compared to Kil et al. who reports −3.5% K⁻¹ at 25 °C[6] and Wan et al. with −2% K⁻¹.[5] However, at higher temperatures our films are comparable to what Kil et al. and Wada et al. report.[6,28] Nevertheless, the low resistivity of the films in this study makes them attractive for reducing
signal noise, making them interesting for bolometric applications even though the TCR is lower.

Some additional findings can be extracted from the temperature-dependent electrical resistivity. The electrical resistivity of VO₂ decreases with temperature, i.e., the films are semiconducting, and the transport properties will be related to the excitation energy or bandgap, ΔE. That means that the electrical resistivity will be an exponential function of ΔE. An Arrhenius plot can be applied to extract the ΔE as shown in **Figure 6c**. With the decrease in film thickness, the slope of the Arrhenius plot varies, resulting in different values of ΔE. In Figure 6c, it is visible that the black line corresponds to the highest slope resulting in a ΔE of 210 meV, which is substantially higher (around 100 meV) than that of the thicker films. This may be due to the increased bandgap or excitation energy ΔE<sub>5nm</sub> of t<sub>film</sub> = 5 nm due to quantum confinement and/or other subtle changes in microstructure or compositional variations.

The temperature stability of the films was investigated by XRD combined with in situ annealing in air, see Figure S4 (Supporting Information) for more details. The results showed that the VO₂ (B) film is stable up to 350 °C. At 375 °C, the V₂O₅ phase was observed and at 400 °C the film was fully transformed to the new structure. The temperature is in similar range as what Tsang and Manthiram observed on a VO₂ (B) sample in differential scanning calorimetry and thermogravimetric analysis in 1997 (performed in N₂ atmosphere).[14] With the distinct difference that the new phase formed in their study was the thermodynamically stable VO₂ in monoclinic rutile (M1) phase. The different phases formed can be explained in the different atmosphere used. In their study, no extra oxygen was available during annealing and thus the formation of V₂O₅ was not possible. While our annealing was performed in air which allowed the VO₂ film to adsorb more oxygen resulting in the more stable V₂O₅ phase to form. It is conceivable that the formation of other phases and polymorphs of vanadium oxide on mica can be achieved by annealing at different gas ambients and partial pressures during growth (see Figure S5, Supporting Information) or post-deposition annealing.

## Conclusions

We have grown 5- to 50-nm thick films of dense metastable VO₂ (B) films of uniform thickness and low surface roughness on mica substrates using pulsed DC magnetron sputtering at relatively high temperatures (400 °C). The phase pure films are semiconducting and exhibit a low resistivity of 50 mΩcm at 30 °C. The electrical transport in these films is strongly dependent on the film thickness in the range 5 nm ≤ t<sub>film</sub> ≤ 21 nm as the characteristic activation energy for carrier transport decreases from ΔE = 210 meV to 110 meV with increasing thickness. The activation energy remains unchanged with further increases in thickness, suggesting quantum confinement effects. A secondary V₂O₅ phase can be formed at films thicker than 50 nm, or by subjecting the films to a post-deposition anneal in air above 350 °C. The formation of V₂O₅ can increase the resistivity by more than two orders in magnitude. The identity of the oxides formed through the decomposition of VO₂ (B) during annealing is sensitive to the gas ambient during deposition and post-deposition annealing. Pulsed DC reactive sputtering offers multiple avenues for

<figure><img src="image_8.png" /><figcaption>**Figure 6.** Resistivity versus thickness is shown in (a). In (b), resistivity versus temperature is shown where red, orange, green and blue lines correspond to 5, 10, 21, and 50 nm, respectively. Arrhenius plots of the t<sub>film</sub> ≤ 50 nm is shown in (c).</figcaption></figure>
phase selection in vanadium oxide thin films, that may not be accessible through other deposition methods.

## Supporting Information

Supporting Information is available from the Wiley Online Library or from the author.

## Acknowledgements

The authors acknowledge funding from the Swedish Research Council (VR) under Project No. 2016-03365, the Knut and Alice Wallenberg Foundation through the Wallenberg Academy Fellows program (grant no. KAW 2020.0196), the Swedish Government Strategic Research Area in Materials Science on Functional Materials at Linköping University (Faculty Grant SFO-Mat-LiU No. 2009 00971), and the Swedish Energy Agency under project 46519-1. The authors also acknowledge access to the Swedish National Infrastructure for Advanced Electron Microscopy, ARTEMIS, supported by the Swedish Research Council (VR) and The Foundation for Strategic Research (SSF), through grants 2021-00171 and RIF21-0026. S. H. acknowledges the French government program "Investissements d'Avenir" (EUR INTREE, Reference No. ANR-18-EURE-0010, and LABEX INTERACTIFS, Reference No. ANR-11-LABX-0017-01).

## Conflict of Interest

The authors declare no conflict of interest.

## Data Availability Statement

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

## Keywords

alEpitaxy, metastable phases, PVD, Thin film, VO₂

Received: March 18, 2023

Revised: May 7, 2023

Published online: June 3, 2023

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