Role of substrate effects on the morphological, structural, electrical and thermoelectrical properties of V₂O₅ thin films

Abstract

The present work focuses on the influence of different substrates on the morphological, compositional, phase purity, structural and transport properties of vanadium pentaoxide (V₂O₅) thin films. Thin films of V₂O₅ were fabricated on different substrates: glass, quartz, Si, and alumina (hereafter these films are referred to as V₂O₅:G, V₂O₅:Q, V₂O₅:Si and V₂O₅:A, respectively) using inorganic sol–gel with V₂O₅ powder and hydrogen peroxide (H₂O₂) as precursors by spin coating. Films deposited on glass substrates were found to be amorphous in nature with smooth surfaces, whereas films deposited on quartz, silicon, and alumina substrates exhibited a polycrystalline nature, having an orthorhombic structure with space group Pmmn. The crystallinity improves from quartz to silicon and the best crystalline films were fabricated on alumina. Electrical measurements as a function of temperature and substrate are investigated and characterized by measuring the resistivity, Hall and Seebeck coefficients. Negative values of Hall and Seebeck coefficients reveal that all the films are of n-type semiconductors. Electrical resistivity as well as charge carrier density decreases from the films on glass to quartz to silicon and to alumina. Carrier mobility decreases in the following order V₂O₅:G > V₂O₅:A > V₂O₅:Si > V₂O₅:Q, whereas the Seebeck coefficient varies in the reverse order. Variation of these transport parameters has been understood on the basis of scattering and trapping of charge carriers along the grain boundaries. Furthermore, a model based on thermodynamics is proposed to explain the effect of substrates on the crystallinity of thin films. Interactions between sol and substrate (adhesive forces) determine the thickness, phase purity, structural and morphological properties of thin films. As the magnitude of adhesive forces increases, both film thickness and crystallinity increase.

---

1. Introduction

Vanadium oxides offer many convenient physical effects, which mean they can be used for different microelectronic device applications.¹ Vanadium is a multivalent element, and as a result forms a variety of oxides upon reacting with oxygen, such as VO, V₂O₃, VO₂, V₆O₁₁, V₂O₅ etc. V₂O₅ is the most saturated (highest oxidation state) in the V–O system and is consequently the most stable among these vanadium oxides.^2,3 It has a lamellar, or sheet-like structure.⁴ It is a distorted orthorhombic structure, and this deformation creates its sheet formation. V₂O₅ in the form of thin films has attracted much attention due to its unique electronic, chemical and optical properties.^5–8 Thin films of V₂O₅ have been prepared by different techniques including vacuum evaporation,⁹ sputter deposition,^10–13 thermal oxidation,^14,15 pulsed laser deposition,¹⁶ chemical vapor deposition,^17–19 sol–gel processes²⁰ and spray pyrolysis.²¹ Due to the advantages of high purity, homogeneity, stoichiometry, simplicity of equipment involved and cost effectiveness for large scale production, the sol–gel process has become one of the best approaches to deposit metal oxide thin films.

Recently, it was reported that the metal to insulator transition (MIT) at approximately 280 °C in V₂O₅ is due to lattice distortion and a structural inhomogeneity due to the vanadyl–oxygen vacancies,^22–24 and this makes V₂O₅ a promising material for thermoelectric devices. V₂O₅ exhibits highly anisotropic electrical and optical properties due to its orthorhombic structure²⁵ and these properties have been extensively investigated by many researchers.^25–28 On the other hand, the thermo-physical properties, including thermal conductivity (κ), Seebeck coefficient (s), and dynamic heat capacity, have received less attention.^29–31 One of the important parameters for characterizing the thermo-physical properties of materials is a power factor (s²/ρ), where s is the thermoelectric power (Seebeck coefficient) and ρ is the electrical resistivity. Thus, to be a better thermoelectric material, a large value for s and small values of ρ are required. Surprisingly, the s and electrical resistivity (ρ) vary in opposite directions, preventing the power factor from improving excellently, and both the values are also dependent on each other with carrier density (n). Generally, the s value of a semiconductor or an insulator decreases with increasing temperature. The variation of s with temperature is very important for thermoelectric applications.

In the present work, we investigate the effect of substrates on the microstructural evolution, crystal phase, and surface morphology, electrical and thermo-electrical properties of V₂O₅ thin films fabricated by sol–gel method. Crystallite size, surface morphology and film thickness were observed to vary from substrate to substrate, which in turn affects the transport properties of thin films.

2. Experimental details

A clear yellow solution, containing 0.3 g V₂O₅ (Sigma-Aldrich, purity >99.9%) dissolved in 30 ml, 30% hydrogen peroxide H₂O₂ (Sigma-Aldrich), was formed at room temperature with vigorous stirring. This yellow solution was then heated at 60 °C with continuous stirring to evolve excess oxygen by decomposition of H₂O₂. Immediately, the solution turned into a red brown viscous gel. After aging for 24 hours, the V₂O₅ gel was ready for coating. The mechanism of the reaction is given in reactions (1) and (2) following Ren et al.³²

V₂O₅ + 2H₂O₂ → 2HVO₄ + H₂O (1)

2HVO₄ + (n − 1)H₂O → V₂O₅·nH₂O + O₂ (2)

The spin coater (Spin-NXG-P1: made by Apex Instruments, India) was used to deposit thin films on glass, quartz, Si, and alumina substrates. Prior to deposition, these substrates were cleaned in dilute sulphuric acid for 30 minutes and then thoroughly rinsed with ethanol, acetone and de-ionized water. Five coatings of V₂O₅ gel were performed on each substrate at a rate of 3000 rpm for 30 seconds and after each coating the films were dried at 80 °C for 15–20 minutes. Finally these V₂O₅ gel films were crystallized by annealing at 400 °C for 4 hours under ambient atmosphere in a programmable tubular furnace (Nabertherm GmbH Tube furnace: RHTC80), with a heating and cooling rate of 3 °C per minute. Rutherford backscattering (RBS) with He⁺ ions of energy 2 MeV was used to characterize the stoichiometry, thickness and uniformity of the films. X-ray diffraction (XRD) measurements at room temperature in the 2θ range 20–60° was carried out to identify the crystalline phases and structure of the films using a Bruker D8 advance diffractometer with Cu Kα (0.15406 nm) X-ray source at a scan speed of 0.5° min⁻¹. Raman studies of these films were carried out and the spectra were measured in back scattering geometry using an Ar excitation source, with a wavelength of 488 nm coupled with a Labram-HR800 micro-Raman spectrometer equipped with a 50× objective, appropriate notch filter, and a Peltier cooled charge-coupled device detector. The surface morphology of films was characterized by SEM (model: MIRA II LMH, TESCAN). The electrical resistivity, Seebeck coefficients of the films were measured in the temperature range 300–400 K using a standard DC four probe technique and bridge method, respectively. The Hall effect measurements were carried out in the same temperature range by Ecopia HMS-3000 Hall Measurement System to evaluate charge carrier density and mobility.

3. Results and discussion

3.1. Morphological studies

Fig. 1 shows the surface morphology of V₂O₅ thin films deposited on glass, quartz, Si, and alumina substrates investigated by SEM, and these films hereafter will be referred to as V₂O₅:G, V₂O₅:Q, V₂O₅:Si, and V₂O₅:A, respectively. As evident, there is a significant change in the surface morphologies. Films deposited on glass substrates have relatively smooth surfaces, with few compact structures revealing the amorphous nature of these films; also some cracks are observed on the surface, which may be due to shrinkage of sol during the annealing of these film. Films deposited on quartz possess irregular nanostructures with varying sizes. The bigger particles grow due to the agglomeration of smaller ones to reduce the surface energy under a thermodynamic driving force, when two particles merge at high temperature. The surfaces of films deposited on Si and alumina possess cubic growth characteristics resembling spiral-like features, which may have emanated from screw dislocations. The average crystallite size was found to be increased while moving from quartz through silicon to alumina. Further grain size distribution was found to be more random for V₂O₅ films deposited on quartz than those deposited on silicon and alumina substrates. This may be due to the difference in sticking coefficients and strains developed in the films grown on an amorphous background. The average grain sizes of V₂O₅:Q, V₂O₅:Si and V₂O₅:A are approximately 70 nm, 135 nm and 180 nm, respectively. Larger grain size was observed in V₂O₅:A compared to V₂O₅:Si due to lattice mismatches of V₂O₅ and Al₂O₃ then V₂O₅ and Si.

Fig. 1 SEM images of (a) V₂O₅:G, (b) V₂O₅:Q, (c) V₂O₅:Si and (d) V₂O₅:A. Films deposited on glass are amorphous in nature with smooth surfaces. Films deposited on quartz, Si and alumina are crystalline in nature and average crystallite size increases from quartz to Si to alumina.

3.2. Compositional analysis

As mentioned above, oxides of vanadium can exist in a wide range of stoichiometries. The RBS measurements with SIMNRA Program were used to determine composition and thickness of films. RBS spectra (with SIMNRA simulation) of the four samples (V₂O₅:G, V₂O₅:Q, V₂O₅:Si, and V₂O₅:A) are shown in Fig. 2(a)–(d), respectively. The simulated results are summarized in Table 1. This simulation suggests that the V/O ratios for all the samples are 2/5 within simulation errors, indicating good stoichiometries of all these four thin films. The thickness was found to be ∼46.7 nm, 85.5, 118.4 and 130.3 nm for films V₂O₅:G, V₂O₅:Q, V₂O₅:Si, V₂O₅:A, respectively. Increase in film thickness from glass to quartz to Si to alumina may be due to the increase in magnitude of cohesive forces in comparison to adhesive forces from V₂O₅:G, V₂O₅:Q, V₂O₅:Si, and V₂O₅:A. It is also clear that the RBS results indicated that there was no contamination during film deposition.

Fig. 2 RBS spectra of (a) V₂O₅:G, (b) V₂O₅:Q, (c) V₂O₅:Si, and (d) V₂O₅:A. Thickness of film increases from glass (46.7 nm) to quartz (85.5 nm) to silicon (118.4 nm) to alumina (130.3 nm).

Table 1 Simulation results on composition of thin films on various substrates

Substrates	Composition of film		Thickness (nm)
	V	O
Glass	0.286	0.714	46.7
Quartz	0.286	0.714	85.5
Si	0.286	0.714	118.4
Alumina	0.286	0.714	130.3

---

3.3. Phase study

Fig. 3 shows the XRD pattern of the V₂O₅ thin films deposited on glass, quartz, Si, and alumina substrates. The X-ray diffraction spectrum of V₂O₅ films deposited on a glass substrate indicates amorphous nature, as the diffraction pattern is diffused and non characteristic. Films fabricated on quartz, Si, and alumina substrates are comprised of the V₂O₅ phase within the XRD detection limit, without any other phases of vanadium oxide. The intensity of the diffraction peaks demonstrates that crystallinity increases from quartz through Si to alumina. The peaks are indexed according to the standard pattern [JCPDS file no. 85-0601] for polycrystalline orthorhombic V₂O₅. The relatively high intensity of the (001) peaks demonstrates the growth of films oriented along the c-axis perpendicular to the surface of the substrate. In addition to the (001) Bragg reflection, the subsequent appearance of other characteristic orientations, such as (101), (110) and (002), reveals the existence of in-plane organization of V–O–V chains. The evaluated lattice parameters on the basis of measured d spacings for films on quartz, Si, and alumina are found to be a = 0.3564 ± 0.001, b = 1.151 ± 0.001, and c = 0.433 ± 0.001 nm; these are in good agreement with previously reported values in the literature.³³ Even though the (001) peak is the strongest in a polycrystalline XRD powder pattern of V₂O₅, the intensity ratios (I(001)/I(hkl)) in the sol–gel prepared films are larger than in polycrystalline V₂O₅ powder, indicating a strong (001) texture. This preferred orientation in the sol–gel films can be understood from the properties of the starting material. The gel is formed by the hydrolysis and condensation of molecular precursors. The chemical control of these reactions allows the formation of V₂O₅ gels directly from the solutions at lower temperature than by the standard solid state process.^34–36 Therefore, the sol–gel films are comprised of V₂O₅·nH₂O, before annealing at 400 °C.³⁷ These have a V₂O₅ layered structure with trapped water molecules and are characterized by a strong structural anisotropy.

Fig. 3 XRD Spectra of (a) V₂O₅:G, (b) V₂O₅:Q, (c) V₂O₅:Si, and (d) V₂O₅:A. The ‘*’ indexed peaks correspond to substrates.

3.4. Raman measurements

Fig. 4 displays the Raman spectra of V₂O₅ thin films fabricated on glass, quartz, silicon, and alumina substrates in the wavelength range 100–1100 cm⁻¹. The broad and non-characteristic Raman spectrum reveals the amorphous nature of V₂O₅ when deposited on the glass substrate. The Raman spectra of films deposited on quartz, silicon and alumina substrates exhibit distinguishable and characteristic assignable Raman peaks, indicating the polycrystalline nature of these films. Peak intensities clearly demonstrate that crystallinity increases from films fabricated on quartz through silicon to alumina. Raman measurements of the films can be described using the shape and frequency of the peaks.

Fig. 4 Raman scattering spectra of (a) V₂O₅:G, (b) V₂O₅:Q, (c) V₂O₅:Si and (d) V₂O₅:A. The ‘*’ indexed peaks correspond to the substrate.

V₂O₅ crystallizes in the orthorhombic system with the group Pmmn (D¹³_2h) according to Bachmann et al.³³ The structure of V₂O₅ is built up from VO₅ square pyramids sharing edges to form (V₂O₄)_n zigzag double chains along [001] and cross linked along [100] by corner sharing and thus forming sheets in the xz plane.^38,39 Thus in each layer V is five-fold coordinated; with three V–O bonds involving three fold coordinated oxygen (O_c) belonging to (V₂O₄)_n chains, one V–O bond involving two fold coordinated oxygen (O_B) constituting bridges between two chains and one involving vanadyl oxygen (O_V). The successive layers are kept together by an equal number of weak van der Waals bonds and much stronger double bonds.⁴⁰ The unit cell contains two formula units (V₄O₁₀) yielding a total number of 39 optical zone centers. Due to the inversion symmetry and according to D_2h factor group analysis, the modes are split into IR (odd) and Raman (even) active modes. All 21 g modes are Raman active while 15 B_u modes are IR active. Beattie and Gilson⁴¹ indicated that a convenient way to describe vibrations in a binary oxide lattice is to treat the oxygen atom as vibrating against an array of immobile metal atoms. Each oxygen species is then considered first as vibrating according to its site symmetry and, following this, the unit cell modes are generated by symmetry operations.

Table 2 depicts the principle Raman active modes with the corresponding motion of V₂O₅ thin films. Each spectrum consists of two groups of peaks located in the high-frequency region, known as internal modes, and in the low frequency region, known as external modes. The internal modes that are observed in the high frequency region are assigned to the stretching and bending of V–O bonds. The high-frequency Raman peak at 1006 cm⁻¹ gives the structural quality of the films and can be ascribed to the stretching mode related to the A_g symmetry vibrations of the shortest vanadium oxygen bond, which is V=O_V. Unlike the other O atoms this atom is strongly bonded to only one V atom and for this reason is called terminal oxygen.⁴² The frequency shift of this mode measures the deviations from stoichiometry. The frequency shift to lower values of this mode is due to softening of the V⁵⁺=O bond in oxygen deficient V₂O₅ films, resulting from vacancies created by removing O_V, with some of the V⁵⁺ reduced to V⁴⁺ in order to balance the charge. Negligible frequency shift of this mode to 995 cm⁻¹ demonstrates good stoichiometry of all three crystalline samples, which is in good agreement with the RBS and XRD results. This is also evident as no peak corresponding to V⁴⁺=O near 932 cm⁻¹ is observed as reported by Lee et al.⁴³ The second peak at 702 cm⁻¹ is assigned to the doubly coordinated oxygen (V₂–O_B) stretching mode, which results from the corner-shared oxygen common to two pyramids. The third peak at 529 cm⁻¹ is assigned to the triply coordinated oxygen (V₃–O_C) stretching mode, which results from the edge-shared oxygen atoms common to three pyramids. The two peaks located at 404 and 283 cm⁻¹ are assigned to the bending vibration of the V=O_V bonds. The peaks located at 481 cm⁻¹ and 304 cm⁻¹ are assigned to the bending vibrations of the bridging V–O_B–V (doubly coordinated oxygen) and V₃–O (triply coordinated oxygen) bonds, respectively.

Table 2 Raman active modes of orthorhombic V₂O₅ thin film^a

Wave number (cm^−1)	Symmetry	Assignment	Remarks
a Symmetry and assignation of modes are mainly based on ref. 42 and 43. Frequency values are derived from present measurements.
146	B3g, B2g	Ty, Rz	Relative motions of V2O5 layers with respect to each other
198	B1g	Tx, Ry	Relative motions of V2O5 layers with respect to each other
283	B2g, B3g	δ(V=O)b	Bending vibration of the V=OV bonds
304	Ag	δ(V3–O)b	Bending vibrations of the V3–O (triply coordinated oxygen) bond
404	Ag	δ(V=O)b	Bending vibration of the V=OV bonds
481	Ag	δ(V–O–V)b	Bending vibrations of the bridging V–OB–V (doubly coordinated oxygen)
529	Ag	ν(V3–O)s	Triply coordinated oxygen (V3–OC) stretching mode
710	B2g, B3g	ν(V–O–V)s	Doubly coordinated oxygen (V2–OB) stretching mode
1006	Ag	ν(V=O)s	Stretching mode vibrations of the shortest vanadium oxygen bond (V=OV)

---

The external modes can be viewed as relative motions of structural units with respect to each other, i.e. translations T¹_x, T¹_y, T¹_z and librations R¹_x, R¹_y, R¹_z. These vibrations occur at low frequencies because each unit is considered heavier than its constituent atoms, while the restoring force has the same order of magnitude.⁴⁴ The external low frequency Raman modes at 146 and 198 cm⁻¹ correspond to the relative motions of V₂O₅ layers with respect to each other.⁴² These two peaks at 146 cm⁻¹ and 198 cm⁻¹ are strongly associated with the layered structure and only appear when there is long range structural order. The presence of these low frequency modes in the films deposited on quartz, silicon and alumina substrates suggest that these films possess a layered structure and are well crystallized. The films grow preferentially, with the c-axis oriented perpendicular to the substrate plane.⁴⁵

3.5. Transport measurements

The temperature dependence of resistivity (ρ) of V₂O₅:G, V₂O₅:Q, V₂O₅:Si and V₂O₅:A thin films in the temperature range 300–400 K is shown in Fig. 5(a). All the samples show semiconducting behavior as electrical resistivity decreases with temperature. Further electrical resistivity decrease occurs in the following order V₂O₅:G > V₂O₅:Q > V₂O₅:Si > V₂O₅:A. Hall measurements of samples demonstrate that all these films are n type semiconductors. Fig. 5(b) displays variation of carrier density (n) and carrier mobility (μ) in the temperature range 300–400 K of the samples. Carrier concentration increases with temperature but mobility decreases with temperature. Decrease in μ values may be caused by increase in carrier scattering due to thermal phonons. Comparison of n and μ of V₂O₅:G, V₂O₅:Q, V₂O₅:Si, and V₂O₅:A films demonstrates that carrier concentration increases in the following order V₂O₅:G < V₂O₅:Q < V₂O₅:Si < V₂O₅:A, whereas mobility increases in the following order: V₂O₅:Q, V₂O₅:Si, and V₂O₅:A to V₂O₅:G in the whole temperature range. Furthermore the differences become more enhanced with temperature. Variation of resistivity, carrier concentration and mobility in these orders can be explained on the basis of thickness and crystallinity differences of the films. Both film thickness and crystallinity increase from V₂O₅:G to V₂O₅:Q to V₂O₅:Si to V₂O₅:A. As film thickness decreases, stacking defects increase, and these defects trap the carriers and lower the free carrier concentration,^46,47 and thus carrier concentration increases as we go from V₂O₅:G to V₂O₅:Q to V₂O₅:Si to V₂O₅:A. In addition, the film grown on the glass substrate is amorphous in nature, whereas films deposited on alumina, silicon and quartz are crystalline in nature. The crystallites formed in V₂O₅:A, V₂O₅:Si and V₂O₅:Q produce a number of grain boundaries and, these boundaries acts as scattering centers for the flow of charge carriers and thus cause reduction in the μ values, as we go from amorphous film fabricated on glass to the crystalline films fabricated on alumina, silicon and quartz. Furthermore, as one goes from V₂O₅:A to V₂O₅:Si to V₂O₅:Q, grain size decreases and thus the number of grain boundaries increases, which results in increase in carrier mobility from V₂O₅:A to V₂O₅:Si to V₂O₅:Q. Fig. 5(c) displays the Seebeck coefficient (s) of V₂O₅:G, V₂O₅:Q, V₂O₅:Si and V₂O₅:A thin films in the temperature range 300–400 K. The negative value of s also demonstrates that all films are n type semiconductors. The s value increases significantly over the entire temperature range from amorphous V₂O₅:G to crystalline V₂O₅:Q, V₂O₅:Si and V₂O₅:A thin films. The s value increases with decrease in crystallite size from V₂O₅:A to V₂O₅:Si to V₂O₅:Q. This can be understood as follows: the energy barrier formed at the grain boundaries acts as an additional trap center and traps low energy charge carriers; this filtering of charge carriers by grain boundaries enhances the average energy of carriers taking part in the transport mechanism.⁴⁸ The value of thermoelectric power depends on the mean carrier energy relative to the Fermi level.⁴⁹ Therefore, this phenomenon of filtering of charge carriers also leads to an increase in s value from V₂O₅:A to V₂O₅:Si, to V₂O₅:Q. Furthermore, with increase in temperature, the average energy of charge carriers taking part in the transport mechanism increases, resulting in an increase in s value with temperature.

Fig. 5 (a) Electrical resistivity, (b) carrier mobility and concentration, and (c) Seebeck coefficient of V₂O₅:glass (b) V₂O₅:Q (c) V₂O₅:Si and (d) V₂O₅:alumina thin films as a function of temperature.

4. Model for crystallization of V₂O₅ thin films on different substrates

Lattice mismatch between the substrates and V₂O₅ and the underlying surface strongly affects nucleation and growth processes. In this section we model for the effect of substrates on the crystallization of V₂O₅ thin films. Fig. 6 shows a schematic diagram of spherical nuclei nucleated on the hetero-substrate. The driving force for nucleation is Gibbs free energy.^50,51 The Gibbs free energy change for the nucleation of spherical nuclei of radius r on the hetero-substrate⁵⁰ is given by:

ΔG_Het. = −VΔG_v + ∑A_Jγ_J (3)

= −VΔG_v + (A_ELγ_EL + A_ESγ_ES − A_SLγ_SL) (4)

= −VΔG_v + ΔG_S, (5)

where V is the volume of the spherical cluster, ΔG_v the Gibbs free energy of molar volume, A_EL and γ_EL the interface area and the interface energy of the embryo–liquid interface, A_ES and γ_ES the interface area and energy of the embryo–substrate interface, and A_SL and γ_SL the interface area and the energy between the substrate–liquid interfaces. According to geometry

A_SL = 2πr²(1 − m) (7)

A_HL = πr²(1 − m²), (8)

where r is the radius of the spherical cluster and m describes the interaction between the liquid embryo and the solid substrate and this has been described in the classical model with the help of the Young equation:⁵¹where θ is the contact or wetting angle between the substrate and the liquid embryo as shown in Fig. 6. Thus, we can obtainwhere the factoris called heterogeneous factor, and its value lies in the range 0–1. Its value is 1, when clusters nucleated on the homo-substrates. The critical nucleus radius r* found by setting ∂ΔG/∂r = 0, with the result

Fig. 6 Schematic diagram displaying the spherical nuclei nucleated on the hetero-substrate.

It is important to note that the critical radius r* remains unchanged for heterogeneous nucleation and homogeneous nucleation. However, the volume (V) can be significantly lower for heterogeneous nucleation, due to the wetting angle affecting the shape of the nucleus. The associated energy barrier for nucleation is found by substituting r* in eqn (8)

= ΔG^*_Homo.f(θ) (17)

The nucleation rate J, which is the number of critical nuclei formed per unit time per unit volume, is usually written in Arrhenius activation form as

where D is a kinetic pre-exponential factor, correlated to encounter frequency between the molecules, and is consequently dependent on their diffusion coefficients, K_B is the Boltzmann constant, and ΔG* is the reversible work of formation of the critical nucleus. The diffusion coefficient D is observed to depend upon temperature as⁵²where E_d is the activation energy for diffusion, K_B is Boltzmann’s constant, T is the absolute temperature, and D₀ is a temperature-independent factor that depends upon the choice of material and deposit. Eqn (18), together with eqn (17) and (19), defines the rate of heterogeneous nucleation of liquid embryos on the heterogeneous substrate. Thus the rate of nucleation depends upon the nature of material of the substrate, the temperature of the substrate and the contact angle between the substrate and the liquid embryo. Hence, maintaining all the substrates at fixed temperature during the deposition of V₂O₅ and post annealing at the same temperature, nucleation of V₂O₅ varies from substrate to substrate. The main parameter that may determine the growth of V₂O₅ thin film on a substrate at fixed temperature is the contact angle, and this contact angle characterizes the interactions between the sol and substrate known as adhesive forces, measuring the wettability of the solid surface by liquid. The work of adhesion, W_ad, is defined as⁵³

W_ad = σlv(1 + cos θ) (20)

A high work of adhesion indicates good wetting, whereas a low work of adhesion indicates poor wetting. Fig. 7(a)–(d) show the wetting and spreading phenomena of water on glass, quartz, silicon, and alumina substrates, respectively. Glass substrates are the most wetted and alumina substrates are the least wetted. The magnitude of the contact angle increases from 19° for glass to 35.4° for quartz to 38.5° for silicon to 46.5° for alumina substrates. A similar trend of increasing contact angle from glass to quartz to silicon to alumina substrate is expected for V₂O₅ sol. This indicates that the magnitude of adhesive forces between sol and the substrate decreases from glass to quartz to Si to alumina substrate. Thus V₂O₅ sol spreads more on glass and the least on alumina. As a consequence of this, the nucleation barrier increases from glass to quartz to silicon to alumina, as shown in Fig. 8, and thus the nucleation rate J, which is the number of critical nuclei formed per unit time per unit volume, decreases from glass to quartz to Si to alumina substrate. Thus, the number of critical nuclei that is formed on the surface of the substrates decreases from glass to quartz to Si to alumina, and hence the thickness and crystallinity of the V₂O₅ thin film increases from glass to quartz to silicon to alumina substrate.

Fig. 7 Images of water contact angle on (a) glass, (b) quartz, (c) Si and (d) alumina substrates.

Fig. 8 Schematic diagram showing variations of Gibbs free energy with radius of nuclei.

5. Conclusion

The nature of substrates determines the properties of V₂O₅ thin films. SEM images showed a change in surface morphology of V₂O₅ films from substrate to substrate. Films deposited on glass substrates are amorphous in nature with a smooth surface, whereas films deposited on quartz, Si and alumina substrates are crystalline. The degree of crystallization increases from quartz to Si to alumina. RBS results show that the thickness of the film increases from glass to quartz to silicon to alumina substrate. XRD and Raman results show that films fabricated on a glass substrate are amorphous in nature, whereas films fabricated on quartz, silicon, and alumina are crystalline, and crystallinity increases from quartz to silicon to alumina. Basically, these are the interactions between sol and substrate (adhesive forces) that play a significant role in elucidating the thickness and degree of crystallization of V₂O₅ thin films. As the magnitude of adhesive forces between the sol and substrate increases, both film thickness and crystallinity are enhanced. The electrical and thermoelectrical properties were found to be functions of crystallization and thickness. As film thickness and crystallization increase from V₂O₅:G to V₂O₅:Q to V₂O₅:Si to V₂O₅:A, both carrier concentration and conductivity increase. Carrier mobility decreases from amorphous film to crystalline film and increases with increasing crystallite size. The Seebeck coefficient shows a strong dependence on crystallization. The Seebeck coefficient increases from amorphous film to crystalline film and then decreases with increase in crystallite size. This behavior of μ and s with crystallization is due to the scattering and trapping of charge carriers along the grain boundaries and stacking defects.

Acknowledgements

One of the authors (B.A.) acknowledges the director, NIT-Srinagar and Director, IUAC, New Delhi for their support and encouragements. He also thanks Materials Science group for all characterization facilities.

References

1. R. B. Darling and S. Iwanaga, Structure, properties, and MEMS and microelectronic applications of vanadium oxides, Sadhana, 2009, 34(4), 531–542.
2. F. J. Morin, Oxides which show a metal-to-insulator transition at the Neel temperature, Phys. Rev. Lett., 1959, 3(1), 34–36.
3. R. L. Smith, G. S. Rohrer, K. S. Lee, D. K. Seo and M. H. Whangbo, A scanning probe microscopy study of the (001) surfaces of V₂O₅ and V₆O₁₃, Surf. Sci., 1996, 367(1), 87–95.
4. F. Hulliger, in Structural Chemistry of Layer-Type Phases, ed. F. Lévy and D. Reidel, Dordrecht, 1976, Fig. 100, with kind permission of Springer Science and Business Media.
5. Y. Wang and G. Cao, Synthesis and enhanced intercalation properties of nanostructured vanadium oxides, Chem. Mater., 2006, 18(12), 2787–2804.
6. V. Petkov, et al. Structure beyond Bragg: Study of V₂O₅ nanotubes, Phys. Rev. B: Condens. Matter Mater. Phys., 2004, 69.8, 085410.
7. G. T. Kim, J. Muster, V. Krstic, J. G. Park, Y. W. Park, S. Roth and M. Burghard, Field-effect transistor made of individual V₂O₅ nanofibers, Appl. Phys. Lett., 2000, 76(14), 1875–1877.
8. C. V. Ramana, R. J. Smith, O. M. Hussain, C. C. Chusuei and C. M. Julien, Correlation between growth conditions, microstructure, and optical properties in pulsed-laser-deposited V₂O₅ thin films, Chem. Mater., 2005, 17(5), 1213–1219.
9. Z. Lu, M. D. Levi, G. Salitra, Y. Gofer, E. Levi and D. Aurbach, Basic electroanalytical characterization of lithium insertion into thin, well-crystallized V₂O₅ films, J. Electroanal. Chem., 2000, 491(1), 211–221.
10. K. West, B. Zachau-Christiansen, S. V. Skaarup and F. W. Poulsen, Lithium insertion in sputtered vanadium oxide film, Solid State Ionics, 1992, 57(1), 41–47.
11. Y. J. Park, K. S. Ryu, K. M. Kim, N. G. Park, M. G. Kang and S. H. Chang, Electrochemical properties of vanadium oxide thin film deposited by RF sputtering, Solid State Ionics, 2002, 154, 229–235.
12. S. C. Mui, J. Jasinski, V. J. Leppert, M. Mitome, D. R. Sadoway and A. M. Mayes, Microstructure effects on the electrochemical kinetics of vanadium pentoxide thin-film cathodes, J. Electrochem. Soc., 2006, 153(7), A1372–A1377.
13. A. Benayad, H. Martinez, A. Gies, B. Pecquenard, A. Levasseur and D. Gonbeau, XPS investigations achieved on the first cycle of V₂O₅ thin films used in lithium microbatteries, J. Electron Spectrosc. Relat. Phenom., 2006, 150, 1–10.
14. T. Szörényi, K. Bali and I. Hevesi, Structural characterization of amorphous vanadium pentoxide thin films prepared by chemical vapour deposition/CVD, J. Non-Cryst. Solids, 1980, 35, 1245–1248.
15. K. Inumaru, T. Okuhara, M. Misono, N. Matsubayashi, H. Shimada and A. Nishijima, EXAFS analysis of vanadium oxide thin overlayers on silica prepared by chemical vapour deposition, J. Chem. Soc., Faraday Trans., 1992, 88(4), 625–630.
16. R. Lindström, V. Maurice, S. Zanna, L. Klein, H. Groult, L. Perrigaud, C. Cohen and P. Marcus, Thin films of vanadium oxide grown on vanadium metal: oxidation conditions to produce V₂O₅ films for Li-intercalation applications and characterisation by XPS, AFM, RBS/NRA, Surf. Interface Anal., 2006, 38(1), 6–18.
17. J. Światowska-Mrowiecka, V. Maurice, S. Zanna, L. Klein, E. Briand, I. Vickridge and P. Marcus, Ageing of V₂O₅ thin films inducedby Li intercalation multi-cycling, J. Power Sources, 2007, 170(1), 160–172.
18. G. J. Fang, Z. L. Liu, Y. Wang, Y. H. Liu and K. L. Yao, Synthesis and structural, electrochromic characterization of pulsed laser deposited vanadium oxide thin films, J. Vac. Sci. Technol., A, 2001, 19(3), 887–892.
19. A. Mantoux, H. Groult, E. Balnois, P. Doppelt and L. Gueroudji, Vanadium oxide films synthesized by CVD and used as positive electrodes in secondary lithium batteries, J. Electrochem. Soc., 2004, 151(3), A368–A373.
20. D. B. Le, S. Passerini, A. L. Tipton, B. B. Owens and W. H. Smyrl, Aerogels and xerogels of V₂O₅ as intercalation hosts, J. Electrochem. Soc., 1995, 142(6), L102–L103.
21. A. Bouzidi, N. Benramdane, A. Nakrela, C. Mathieu, B. Khelifa, R. Desfeux and A. Da Costa, First synthesis of vanadium oxide thin films by spray pyrolysis technique, J. Mater. Sci. Eng. B, 2002, 95(2), 141–147.
22. M. Kang, I. Kim, S. W. Kim, J. W. Ryu and H. Y. Park, Metal–insulator transition without structural phase transition in V₂O₅ film, Appl. Phys. Lett., 2011, 98(13), 131907.
23. R. P. Blum, H. Niehus, C. Hucho, R. Fortrie, M. V. Ganduglia-Pirovano, J. Sauer and H. J. Freund, Surface metal–insulator transition on a vanadium pentoxide (001) single crystal, Phys. Rev. Lett., 2007, 99(22), 226103.
24. A. L. Pergament, G. B. Stefanovich and A. A. Velichko, Oxide Electronics and Vanadium Dioxide Perspective: A Review, Journal on Selected Topics on Nano Electronics and Computing, 2013, 1, 24.
25. C. V. Ramana, O. M. Hussain, B. S. Naidu and P. J. Reddy, Spectroscopic characterization of electron-beam evaporated V₂O₅ thin films, Thin Solid Films, 1997, 305(1), 219–226.
26. M. Zhu, Z. Zhang and W. Miao, Intense photoluminescence from amorphous tantalum oxide films, Appl. Phys. Lett., 2006, 89(2), 021915.
27. M. Losurdo, G. Bruno, D. Barreca and E. Tondello, Dielectric function of V₂O₅ nanocrystalline films by spectroscopic ellipsometry: Characterization of microstructure, Appl. Phys. Lett., 2000, 77(8), 1129–1131.
28. M. I. Kang, I. K. Kim, E. J. Oh, S. W. Kim, J. W. Ryu and H. Y. Park, Dependence of optical properties of vanadium oxide films on crystallization and temperature, Thin Solid Films, 2012, 520(6), 2368–2371.
29. R. Santos, J. Loureiro, A. Nogueira, E. Elangovan, J. V. Pinto, J. P. Veiga and I. Ferreira, Thermoelectric properties of V₂O₅ thin films deposited by thermal evaporation, Appl. Surf. Sci., 2013, 282, 590–594.
30. S. Iwanaga, M. Marciniak, R. B. Darling and F. S. Ohuchi, Thermopower and electrical conductivity of sodium-doped V₂O₅ thin films, J. Appl. Phys., 2007, 101(12), 123709.
31. E. Lazaro, M. Bhamidipati, M. Aldissi and B. Dixon, Thermoelectric Organics, USAR Office, East Falmouth, Massachusetts, USA, 1998.
32. X. Ren, Y. Jiang, P. Zhang, J. Liu and Q. Zhang, Preparation and electrochemical properties of V₂O₅ submicron-belts synthesized by a sol–gel H₂O₂ route, J. Sol-Gel Sci. Technol., 2009, 51(2), 133–138.
33. H. G. Bachmann, F. R. Ahmed and W. H. Barnes, The crystal structure of vanadium pentoxide, Z. Kristallogr., 1961, 115(1–6), 110–131.
34. K. S. Hwang, J. H. Ahn and B. H. Kim, Vanadium-doped nanocrystalline TiO₂ layer prepared by spin coating pyrolysis, J. Mater. Sci., 2006, 41(10), 3151–3153.
35. J. Livage, Vanadium pentoxide gels, Chem. Mater., 1991, 3(4), 578–593.
36. Sol-gel science: the physics and chemistry of solgel processing, ed. C. J. Brinker and G. W. Scherer, Gulf Professional Publishing, 1990.
37. B. Alonso and J. Livage, Synthesis of Vanadium Oxide Gels from Peroxovanadic Acid Solutions: A51V NMR Study, J. Solid State Chem., 1999, 148(1), 16–19.
38. C. V. Ramana, O. M. Hussain, B. S. Naidu and P. J. Reddy, Spectroscopic characterization of electron-beam evaporated V₂O₅ thin films, Thin Solid Films, 1997, 305(1), 219–226.
39. L. Abello, E. Husson, Y. Repelin and G. Lucazeau, Vibrational spectra and valence force field of crystalline V₂O₅, Spectrochim. Acta, Part A, 1983, 39(7), 641–651.
40. J. Haber, M. Witko and R. Tokarz, Vanadium pentoxide I. Structures and properties, Appl. Catal., A, 1997, 157(1), 3–22.
41. I. R. Beattie and T. R. Gilson, Oxide phonon spectra, J. Chem. Soc. A, 1969, 2322–2327.
42. P. Clauws, J. Broeckx and J. Vennik, Lattice Vibrations of V₂O₅. Calculation of Normal Vibrations in a Urey–Bradley Force Field, Phys. Status Solidi B, 1985, 131(2), 459–473.
43. S. H. Lee, H. M. Cheong, M. J. Seong, P. Liu, C. E. Tracy, A. Mascarenhas and S. K. Deb, Raman spectroscopic studies of amorphous vanadium oxide thin films, Solid State Ionics, 2003, 165(1), 111–116.
44. L. Abello, E. Husson, Y. Repelin and G. Lucazeau, Vibrational spectra and valence force field of crystalline V₂O₅, Spectrochim. Acta, Part A, 1983, 39(7), 641–651.
45. G. J. Fang, Z. L. Liu, Y. Q. Wang, H. H. Liu and K. L. Yao, Orientated growth of V₂O₅ electrochromic thin films on transparent conductive glass by pulsed excimer laser ablation technique, J. Phys. D: Appl. Phys., 2000, 33(23), 3018.
46. S. S. Lin, J. L. Huang and D. F. Lii, The effect of thickness on the properties of Ti-doped ZnO films by simultaneous rf and dc magnetron sputtering, Surf. Coat. Technol., 2005, 190(2), 372–377.
47. Y. Qu, T. A. Gessert, K. Ramanathan, R. G. Dhere, R. Noufi and T. J. Coutts, Electrical and optical properties of ion beam sputtered ZnO:Al films as a function of film thickness, J. Vac. Sci. Technol., A, 1993, 11(4), 996–1000.
48. M. Bala, S. Gupta, T. S. Tripathi, S. Varma, S. K. Tripathi, K. Asokan and D. K. Avasthi, Enhancement of thermoelectric power of PbTe:Ag nanocomposite thin films, RSC Adv., 2015, 5(33), 25887–25895.
49. J. Martin, L. Wang, L. Chen and G. S. Nolas, Enhanced Seebeck coefficient through energy-barrier scattering in PbTe nanocomposites, Phys. Rev. B: Condens. Matter Mater. Phys., 2009, 79(11), 115311.
50. G. W. Yang and B. X. Liu, Nucleation thermodynamics of quantum-dot formation in Vgroove structures, Phys. Rev. B: Condens. Matter Mater. Phys., 2000, 61(7), 4500.
51. D. Y. Kwok, C. N. C. Lam, A. Li, A. Leung, R. Wu, E. Mok and A. W. Neumann, Measuring and interpreting contact angles: a complex issue, Colloids Surf., A, 1998, 142(2), 219–235.
52. C. Kittel, Introduction to solid state physics, Wiley, New York, 1968.
53. A. W. Neumann and R. J. Good, Surface and Colloid Science, ed. R. J. Good and R. R. Stromberg, Plenum Press, New York, 1979, vol. II.

---