Tuning the structural, optical, and electrical properties of Sb₂O₃ thin films

Abstract

Antimony trioxide (Sb₂O₃) thin films were deposited using aerosol-assisted chemical vapor deposition (AACVD). Varying the deposition temperature and precursor concentration increased the film thickness and resulted in a shift from phase-pure senarmontite Sb₂O₃ to mixed Sb₂O₃/Sb⁰ films due to a temperature-driven reductive transformation. These changes altered crystallinity, grain structure, defect chemistry, and light-scattering behaviour, which in turn controlled the electrical, optical, and surface properties of the coatings. Films deposited at higher temperatures produced thicker films containing metallic Sb crystallites that enhanced charge transport but reduced optical clarity. X-ray diffraction (XRD) was used to evaluate the structural properties of the films, while scanning electron microscopy (SEM) revealed distinct changes in surface morphology with varying deposition conditions. X-ray photoelectron spectroscopy (XPS) analysis revealed the oxidation state of Sb in the films, with temperature-dependent variations in residual carbon content. Despite its wide technological relevance, Sb₂O₃ remains largely unexplored using AACVD. This work represents a detailed study of antimony-based thin films prepared via AACVD, demonstrating tunable film properties and potential for use in optoelectronic and sensing applications.

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Introduction

In recent years, semiconducting group V–VI compounds have gained significance in various technological domains owing to their optoelectronic properties and applications.^1–6 Among these materials, antimony trioxide (Sb₂O₃) has emerged as a promising candidate due to its wide band gap (>3.6 eV) and reported thermal and phase stability.^7–9 Thermal phase transitions and stability of Sb₂O₃ polymorphs have previously been investigated using thermal analysis techniques.^9,10 Orman and Holland reported that the senarmontite–valentinite phase transition occurred at elevated temperatures (∼615 °C), while the oxidation behaviour was strongly dependent on the specific Sb₂O₃ polymorph.¹⁰ Sb₂O₃ is a semiconductor with a direct wide band gap of 3.40–4.0 eV.^7,11 Sb₂O₃ has long been used as an effective flame-retardant synergist in the polymer and plastic industry.^12–14 Moreover, it has found uses as a catalyst in photochemistry and organic synthesis.^15–20

Like many other metal-oxide semiconductors, Sb₂O₃ has been employed in optoelectronic devices such as solar cells and UV LEDs.^21–24 It is also used in heat-reflective coatings (heat mirrors), where it enhances the infrared reflectivity of the surface and improves thermal efficiency.^25,26 Antimony trioxide exists in two polymorphic phases: the low-temperature senarmontite phase, exhibiting a cubic shape, and the high-temperature valentinite phase, which displays an orthorhombic shape.^10,27,28

Optoelectronic and structural properties of Sb₂O₃ thin films have widely been investigated using vacuum deposition methods.^{8,25,26,29,30} Tigau et al. reported that substrate temperature strongly influenced the crystallinity, optical transmittance and optical band gap of polycrystalline Sb₂O₃ thin films.⁸ Tigau et al. also demonstrated that post-deposition heat treatment significantly altered the electrical conductivity of Sb₂O₃ films due to thermally induced structural and microstructural changes.²⁵ Divya et al. reported transition-metal-doped Sb₂O₃ thin films exhibited high near-IR transmittance together with anti-reflective and UV-blocking behaviour, highlighting the multifunctional optical properties achievable in Sb₂O₃-based coatings.²⁶ Tigau further investigated the preparation and characterization of Sb₂O₃ thin films, reporting their optical transparency, band gap behaviour and surface morphology.²⁹ In a separate study, Tigau et al. demonstrated that film thickness strongly influenced the crystallinity, optical transmittance and electrical resistivity of Sb₂O₃ thin films.³⁰ Binions et al. deposited insulating Sb₂O₃ films via atmospheric pressure chemical vapor deposition (APCVD) and investigated their gas sensing properties.³¹

While most reported studies have used vacuum-based techniques, these methods often involve high-cost equipment and limited scalability for large-area coatings.³² In contrast, solution-based CVD routes such as AACVD offer a simpler, more controllable and cost-effective alternative for producing high-quality oxide films under atmospheric conditions. In addition, AACVD does not require highly volatile precursors, allowing greater flexibility in precursor and solvent selection compared to many vacuum deposition techniques. This technique has been successfully applied to deposit transparent conducting oxides such as ZnO^33–35 and SnO₂.^36,37 Although film growth in AACVD can be influenced by aerosol transport and local gas flow dynamics, reproducibility is generally less challenging for undoped oxide systems than for multicomponent or heavily doped films, where precise dopant incorporation becomes increasingly important. In the present study, all depositions were repeated three times under identical conditions, with negligible variation observed in the optical and electrical properties of the films. However, to the best of our knowledge, no detailed study has yet reported the deposition and optimisation of Sb₂O₃ thin films via AACVD.

In this work, Sb₂O₃ thin films were deposited at varying substrate temperatures and precursor concentrations. The influence of these parameters on the structural, morphological, optical, and electrical properties of the films was systematically investigated. The study particularly focuses on the relationship between film composition and the resulting film conductivity, as well as the onset of partial reduction to metallic Sb at higher temperatures. This work provides new insights into the controlled synthesis of Sb₂O₃ thin films via AACVD, expanding their potential use in optoelectronic, sensing, and dielectric applications.

Experimental

All chemicals were used as received without further purification: antimony(III) ethoxide, Sb(OC₂H₅)₃ (99.9%, Fischer Scientific) and anhydrous toluene (99.8%, Fischer Scientific). Nitrogen gas was used as supplied by BOC. The glass substrate was 3.2 mm thick, plain float glass with a 50 nm thick SiO₂ barrier layer, supplied by Pilkington NSG.

Synthesis

For all the depositions, the precursor solution was prepared by dissolving Sb(OC₂H₅)₃ in anhydrous toluene. A standard float glass, with 50 nm thick SiO₂ ion-diffusion inhibiting layer coated on the top surface (supplied by NSG Pilkington Ltd), with the dimensions of 15 cm × 5.0 cm × 0.3 cm was used as substrate. The SiO₂ layer acted as a barrier to prevent leaching of ions between the substrate and the film. Prior to depositing, the glass substrate was cleaned with detergent, and isopropanol, respectively.

A horizontal bed AACVD reactor was used, with the glass substrate placed horizontally on a smooth graphite heating block. A top plate, suspended above the substrate, ensured laminar flow of the aerosol for uniform deposition. This entire setup was housed inside a quartz tube. The substrates were heated at different temperatures (480–580 °C). The precursor solution was nebulized using a piezo ultrasonic atomizer from Johnson Matthey, operating at a frequency of 1.6 MHz to generate a mode droplet size of 3 µm. The precursor solution and nebulizer were maintained at ambient laboratory temperature during deposition and no external heating or active temperature stabilisation was applied.

The generated aerosol mist of the precursor solution was transported by a carrier N₂ gas into the deposition chamber. The entire setup was placed in a fume cupboard. The precursor solution was completely misted in 30 ± 5 minutes to deposit the film, after which the substrate was left to cool in reactor under continuous flow of N₂ gas and removed when the temperature of the reactor reached below 50 °C.

Analysis

Grazing-incident X-ray diffraction (GIXRD) measurements were recorded utilizing a Panalytical Empyrean diffractometer. X-rays were generated using a Cu Kα source with a wavelength of 1.5406 Å, a voltage of 40 kV, and an emission current of 40 mA. The incident beam angle was maintained at 1°, and patterns were recorded within a range from 10 to 70° (with 0.05° steps at 0.5° per step). Peak positions were determined by comparison with standard data from the Inorganic Crystal Structure Database (ICSD). All the recorded patterns were analysed to assess crystallinity and preferred growth orientation and subsequently plotted using OriginPro.

X-ray photoelectron spectroscopy (XPS) analyses of the films were conducted using a Thermo Scientific spectrophotometer equipped with a monochromatic Al-Kα radiation source. The chemical compositions were determined through high-resolution surface scans of the Sb 3d, O 1s, and C 1s regions. These scans were carried out with a pass energy of 50 eV and a spot size of 400 µm. Identification and quantification of peaks were performed using Avantage Thermo Fisher scientific software and plotted using OriginPro. The XPS spectra were quantified to determine the O/Sb ratios in Sb₂O₃ films. Binding energies of the peaks were calibrated relative to the adventitious carbon peak at 284.5 eV to correct for charging effects.

UV-Vis transmittance and reflectance spectra were recorded using Shimadzu 3600i plus spectrometer over a wavelength range of 250–2500 nm. The band gap of films was calculated from acquired spectra in OriginPro using Tauc plot method.

Scanning electron microscopy (SEM) was employed to examine the surface morphology of the films in top-down configuration and to estimate the film thickness using cross-sectional configuration. JEOL IT-700HR was used to analyse the samples. All measurements were performed at an accelerating voltage of 5–15 kV. Prior to imaging, the films were vacuum sputtered with a thin layer of carbon to minimize surface charging.

Electrical properties were measured by Hall effect measurements using a four-point probe in van der Pauw geometry on an Ecopia HMS-3000 setup. Recordings were performed on 1 cm² samples with an applied current of 1 nA–1 µA under a calibrated magnetic field of 0.58 T. It provided the values of resistivity, mobility, and carrier concentration.

Finally, static water contact angles were measured using a Krüss DSAE droplet shape analyser employing the sessile drop technique. Measurements were carried out at ambient conditions using 5 µL droplets of ultrapure water. The contact angle values were obtained by fitting the droplet profiles from the recorded images using the Krüss DSAE analysis software.

Results and discussion

Sb₂O₃ films were deposited via AACVD using antimony(III) ethoxide dissolved in anhydrous toluene as the precursor solution. Antimony(III) ethoxide was selected as the precursor due to its good solubility in organic solvents and suitability for solution-based AACVD processing. In contrast to chloride-based precursors, alkoxide precursors can reduce the possibility of halogen incorporation and corrosive by-products during deposition. In addition, precursor volatility is not a strict requirement in AACVD, making Sb(OEt)₃ a suitable precursor for controlled oxide film growth.

The films were deposited on float glass substrates positioned on the graphite block in the horizontal-bed AACVD reactor. All films were deposited on the bottom substrate plates. Deposition parameters including substrate temperature (480–580 °C) and precursor molarity (0.4–0.6 M) were systematically varied to study their influence on the structural, morphological, optical, and electrical properties of the resulting films. The specific deposition conditions for each sample are summarized in Table 1.

Table 1 Deposition conditions used for the preparation of Sb₂O₃ thin films by AACVD, including deposition temperature, precursor composition and precursor molarity

Film	Deposition temp. (°C)	Antimony(III) ethoxide (mL)	Anhydrous toluene (mL)	Precursor molarity (M)
S1	480	3.1	24	0.4
S2	500	3.1	24	0.4
S3	540	4.7	24	0.6
S4	580	4.7	24	0.6

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A wider range of deposition temperatures was initially explored during preliminary optimisation experiments. Films deposited below 450 °C were found to be very thin and powdery, making them unsuitable for further analysis. The upper deposition temperature was also limited by the AACVD setup and glass substrate stability, as temperatures above ∼590 °C were not suitable for the substrates used in this study. The selected temperature range therefore enabled systematic investigation of the transition from phase-pure Sb₂O₃ to mixed Sb₂O₃/Sb films at elevated deposition temperatures.

To confirm reproducibility, all films were deposited three times under identical conditions. The reproduced samples showed negligible variation in optical transparency and electrical conductivity. The Sb₂O₃ films (S1–S2) adhered strongly to the substrate, passing the Scotch tape test. No film removal or edge peeling was observed after 20 test cycles, confirming the mechanical stability of the coatings. However, sample S3 and S4 deposited at 540 °C and 580 °C showed the formation of metallic Sb crystallites due to partial thermal reduction. The bright, faceted Sb crystallites were scratchable with a metal spatula, indicating poorer adhesion of the metallic phase compared to the oxide matrix. This reduced stability likely resulted from differences in thermal expansion and weaker interfacial bonding of metallic Sb to the underlying oxide and glass substrate.

All the films were stored in ambient air for 14 months, during which no measurable change in their electrical or optical properties was observed. This long-term stability indicates that the films properties were stable under normal atmospheric conditions.

The normalized powder X-ray diffraction (XRD) patterns of the as-deposited films (S1–S4) are shown in Fig. 1. The reference diffraction patterns of senarmontite (cubic, ICSD No. 1011201) and Sb (metal, ICSD No. 2310880) phases were obtained from the Inorganic Crystal Structure Database (ICSD) and plotted using the VESTA software package³⁸ (Fig. 1). Films S1 and S2 were found to be phase-pure and crystallized in the senarmontite phase, which is the most reported polymorph for Sb₂O₃ when deposited below 600 °C.¹⁰ In contrast, S3 and S4 showed additional diffraction peaks corresponding to metallic Sb, indicating partial reduction of Sb³⁺ species at higher deposition temperatures (540–580 °C). This confirmed that partial thermal reduction of Sb₂O₃ occurred at the highest deposition temperature, leading to the formation of mixed Sb₂O₃–Sb phases. Such oxide-to-metal transitions at high deposition temperatures have previously been reported when antimony(III) ethoxide was used as a precursor.^39,40

Fig. 1 Collated XRD diffractograms of Sb₂O₃ thin films (S1–S4) deposited by AACVD under the conditions listed in Table 1. Reference patterns of cubic senarmontite Sb₂O₃ and metallic Sb are included for comparison.

While Sb₂O₃ is widely reported to exhibit excellent thermal stability as an oxide, the partial formation of metallic Sb observed at higher deposition temperatures in this study arises from the specific AACVD environment and precursor chemistry. At high deposition temperatures, the decomposition of antimony(III) ethoxide in an organic solvent can generate locally reducing conditions, promoting partial reduction of Sb³⁺ to Sb⁰. The use of toluene solvent together with N₂ carrier gas may further contribute to a relatively low oxygen partial pressure within the AACVD reactor, favouring the formation of metallic Sb during deposition.

The crystallite sizes (D) were calculated using the Scherrer formula,⁴¹ as shown in eqn (1):

where k is shape factor (0.9), λ is the wavelength (1.5406 Å) of X-ray used, θ_B is the Bragg diffraction angle, and β is the full width of half maximum (FWHM) of the diffraction peak. It should be noted that the Scherrer equation provides only an approximate estimation of crystallite size, as it assumes that peak broadening arises solely from crystallite size and does not account for contributions from lattice strain or instrumental broadening. Therefore, the calculated crystallite sizes are primarily useful for comparative analysis between the deposited samples.

The crystallite size was found to grow from 31.3 nm to 45.0 nm as the substrate temperature increased from 480 °C (S1) to 540 °C (S3). This trend can be attributed to the enhanced surface and grain boundary mobility at elevated temperatures, which can facilitate the migration and merging of adjacent grains.⁴² As a result, smaller crystallites can merge to form larger ones, leading to an increase in crystallite size. However, for S4 (580 °C, 0.6 M), the crystallite size decreased to 29.0 nm. This reduction is likely due to excessive nucleation and partial reduction of Sb₂O₃ to metallic Sb at elevated temperature and higher precursor concentration. The coexistence of oxide and metallic phases can introduce lattice strain and grain boundary pinning, which can disrupt uniform grain growth and lead to smaller crystallites.

The film thickness values were found to be between 440–1020 nm, as shown in Table 3. It is evident that the film was thinner at lower deposition temperature which may be attributed to the incomplete decomposition of the precursor. At higher temperature the decomposition became more complete, and the films grew thicker. The film thickness was found to increase linearly with substrate temperature. Increasing the precursor concentration also yielded thicker film which can be due to the greater flux of reactive species reaching the substrate.

The increase in film thickness with deposition temperature and precursor concentration significantly influenced the resulting film properties. Thicker films deposited at higher temperatures showed lower optical transmittance and improved electrical conductivity. In contrast, thinner films deposited at lower temperatures remained more transparent but showed higher resistivity. These results demonstrate that film thickness played an important role in controlling the balance between optical transparency and electrical transport in Sb₂O₃-based films.

X-ray photoelectron spectroscopy (XPS) was used to investigate the chemical state and composition of the antimony oxide films. XPS survey spectrum of S2, shown in Fig. S1, indicated the presence of only Sb, O, and C peaks. The Sb 3d XPS spectra of S1–S4 are shown in Fig. 2a. Within the energy range of 520–540 eV, spectra exhibited complex features due to the overlapping of O1s and Sb 3d_5/2 photoemission lines at around 530 eV.⁴³ As a result, in most of the reported literature, the complete deconvolution of O1s and Sb 3d_5/2 peaks have not been conducted.^44,45

Fig. 2 High-resolution XPS spectra of Sb₂O₃ thin films S1–S4 showing (a) Sb 3d and (b) C 1s regions.

In this study, spectral deconvolution was carried out while considering spin–orbit coupling, setting the Sb 3d_3/2/Sb 3d_5/2 area ratio to be 2 : 3, and fixing the splitting energy of the doublet at 9.39 eV.⁴³ A deconvoluted Sb 3d spectrum is shown in Fig. S2. This method allowed for the independent quantification of the contributions from the Sb 3d and O 1s peaks.⁴⁶ The Sb 3d_3/2 peak, which is free from O 1s overlap, was used to determine the antimony oxidation state. Sb(III) reportedly appears at 539.3–539.7 eV and Sb(V) at 540.3–540.8 eV.^44,47,48 In all the studied samples (Fig. 2a), the Sb 3d_3/2 peak was located at approximately 539.5 eV, which confirmed the presence of Sb(III) in all the films. Moreover, no contribution from Sb(V) was found in any of the films, which further confirmed the deposition of Sb₂O₃.

The XPS spectrum of sample S3 and S4 revealed the coexistence of Sb₂O₃ and metallic Sb species. The Sb 3d_5/2 peak of metallic antimony was observed at approximately 528 eV, consistent with previously reported values for Sb⁰ in the literature.⁴⁹ The simultaneous presence of these features confirmed partial thermal reduction of Sb₂O₃ at 580 °C, consistent with the XRD observations. This mixed chemical state suggested that metallic Sb crystallites formed within the Sb₂O₃ matrix during high-temperature depositions, resulting in a composite oxide–metal film.

The O 1s peak consisted of two components, corresponding to distinct oxygen bonding states present in the samples. The deconvolution of O 1s peaks in Sb 3d spectrum is shown in Fig. S2. The first peak centred at approximately 530.3 eV was assigned to lattice oxygen corresponding to Sb–O bonds. Similar O 1s signals at approximately same binding energies have been previously reported and linked to the metal–oxygen bond.^45,50–54 However, the higher binding energy peak at around 532.2 eV was attributed to surface hydroxyl group (–OH).^55,56

The C 1s spectra of all the films are shown in Fig. 2b. C 1s was deconvoluted into three components. The peak at 284.8 eV corresponds to C–C from adventitious hydrocarbons. The component at 286.5 eV is associated with C–O species (alcohol/ether or alkoxy groups). The high-binding-energy peak at 288.9 eV is due to O–C=O groups.⁵⁷

XPS spectra were quantified to determine the O/Sb atomic ratio and the carbon content (at%), as shown in Table 3. The films contained 4–26 at% carbon, depending on the deposition parameters. The brown colouration and reduced transparency are consistent with elevated carbon content. Use of toluene as the solvent likely increased carbon incorporation into the films. Hassan et al. reported a similar brown tinge when toluene was used and attributed it to carbon contamination (>5 at% by EDS).⁵⁸ In a broader solvent study on AZO, Potter et al. reported higher carbon contamination when toluene was used instead of methanol.⁵⁹ In the present study, methanol was avoided due to precursor instability, and toluene was selected for its stable solubility of Sb(OC₂H₅)₃. Brown colouration of the films can also due to appearance of metallic Sb⁰ in the films.

Although residual carbon was detected by XPS, no direct correlation between carbon content and enhanced conductivity was observed. In fact, higher carbon concentration is reported to introduce localized trap states at grain boundaries, which can limit carrier mobility.^60,61 The significant reduction in resistivity observed for S3 and S4 is therefore primarily attributed to the formation of metallic Sb⁰ inclusions rather than carbon-related conduction pathways.

The XPS depth profiling (Fig. 3) spectra of sample S4 are summarized in Table 2. At the surface, the C 1s spectrum was dominated by the C–C/C–H component (79.1 at%), accompanied by minor oxidized carbon species, C–O–C (12.9 at%) and O–C=O (7.9 at%). The carbonate-related O–C=O contribution decreased steadily with sputtering time and became less than 6.0 at% beyond 100 s, confirming that these species are confined to the near-surface region and arise from air exposure or CO₂ adsorption rather than bulk incorporation. In contrast, the C–C intensity remained nearly constant throughout the etching process, consistent with hydrocarbon contamination. XPS depth profiling of S1 (Fig. S3) also showed a systematic decrease in the C–C, C–O–C, and O–C=O contributions with depth profiling time, indicating that the carbon-related species were primarily confined to the near-surface region.

Fig. 3 XPS depth profiling spectra of sample S4 showing the evolution of (a) Sb 3d and (b) C 1s regions with increasing sputtering time.

Table 2 Relative atomic percentages of carbon (C–C/C–H, C–O–C, O–C=O), oxygen (O–Sb, Sb–OH), and antimony (Sb⁰, Sb³⁺) species obtained from XPS depth profiling of S4. Samples (S1–S4) correspond to the deposition conditions listed in Table 1

Etching time (s)	C–C (rel. at%)	C–O–C (rel. at%)	O–C=O (rel. at%)	O–Sb (rel. at%)	Sb–OH (rel. at%)	Sb^0 (rel. at%)	Sb^3+ (rel. at%)
0 (Surface)	79.1	12.9	7.9	56.9	43.1	36.9	63.1
50	81.9	12.2	5.9	56.9	43.1	29.7	70.3
100	81.6	12.8	5.6	56.1	43.9	27.2	72.8
150	82.7	12.4	4.9	55.1	44.9	27.5	72.5
200	81.8	12.7	5.5	55.1	44.9	27.9	72.1

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Table 3 Film thickness, electrical properties, optical properties, and carbon content of all the films prepared by AACVD. Samples (S1–S4) correspond to the deposition conditions listed in Table 1

Film	Film thickness (±50 nm)	Resistivity (×10^−1 Ω cm)	Mobility (cm^2 V^−1 s^−1)	Carrier conc. (cm^−3)	Sheet resistance (Ω □^−1)	F.O.M. (×10^−6 Ω^−1)	Band gap (eV)	Tλ550 (%)	Tλ380–750 (%)	Tλ750–1400 (%)	Tλ1400–2500 (%)	C (at% from XPS data)
S1	440	2100 ± 200	8	3.7 × 10^15	4 800 000	0.02	3.63 ± 0.03	79.2	78.6	82.8	86.7	4.0
S2	490	590 ± 50	10	1.0 × 10^16	1 200 000	0.07	3.72 ± 0.03	77.6	77.2	82.9	85.6	7.0
S3	750	8.2 ± 0.8	14	5.5 × 10^17	11 000	1.98	3.59 ± 0.04	68.2	67.7	75.4	78.9	26.0
S4	1020	1.1 ± 0.1	11	5.2 × 10^18	1100	0.03	3.47 ± 0.06	36.1	36.2	54.2	63.7	21.0

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The lattice oxygen (O–Sb) and surface hydroxyl oxygen (Sb–OH) maintained a nearly constant ratios throughout the sputtering time. However, the relative fraction of Sb⁰ decreased from 36.9 at% to 27.9 at%, while Sb³⁺ increased from 63.0 at% to 72.1 at% with sputtering time. These results indicated that surface carbonates were confined to the outermost layer, while the OH-related oxygen species persisted with depth. The sustained presence of the Sb–OH component suggested that hydroxyl-related oxygen remained stable even after surface carbon removal.

The SEM images (Fig. 4) show a systematic evolution in surface morphology of the Sb₂O₃ films with changes in deposition temperature and precursor concentration. At lower deposition temperature (S1, 480 °C) the film showed small spherical grains. The film coverage improved at 500 °C (S2), producing smoother and compact surface composed of spherical grains of around 100 nm. This morphology reflects limited surface diffusion and slow grain growth, typical of kinetically controlled oxide deposition at lower temperatures.^33,62,63 The high resistivity and low mobility of these samples can be attributed to their less dense surface morphology.

Fig. 4 Top-view SEM images of Sb₂O₃ thin films deposited under the conditions listed in Table 1: (a) S1 (b) S2 (c) S3 (d) S4, showing the evolution of surface morphology with increasing deposition temperature and precursor concentration.

With increasing precursor concentration (S3, 540 °C, 0.6 M), the grains became larger, and the surface roughness increased due to accelerated particle agglomeration and presence of Sb⁰ as suggested by XRD and XPS. The improved grain connectivity and densification in this sample along with the presence of metallic antimony was consistent with the observed reduction in resistivity and higher carrier concentrations. While the increased surface roughness and presence of Sb⁰ contributed to lower optical transmittance.

At the highest deposition temperature (S4, 580 °C, 0.6 M), the surface morphology changed markedly. The film exhibited large, faceted grains with bright metallic contrast, indicating partial reduction of Sb₂O₃ to elemental Sb. These metallic crystallites were dispersed within the oxide matrix but did not form a continuous conducting network, resulting in a mixed-phase structure. The presence of these metallic domains explained the sharp decrease in resistivity compared to the oxide samples. However, the film still retained semiconducting characteristics due to the non-percolating nature of the Sb inclusions.

Large, faceted crystallites with well-defined polygonal geometry dominated the surface, consistent with the metallic phase observed in XRD. These micrometre-sized grains exhibited sharp edges and smooth faces, typical of the trigonal crystal structure of elemental antimony. This confirmed the thermal reduction of Sb₂O₃ to metallic Sb at higher deposition temperature. The loosely packed crystallites suggested localized metallic nucleation and growth within the antimony(III) oxide matrix. This observation is fully consistent with the XRD and XPS results that showed both Sb⁰ and Sb³⁺ states.

UV-Vis-NIR spectroscopy was used to evaluate the optical properties of the films. Fig. 5 shows the transmittance and reflectance spectra recorded in the wavelength range of 250–2600 nm. The average transmittance progressively decreased with increasing deposition temperature and precursor concentration.

Fig. 5 Optical transmittance (T) and reflectance (R) spectra together with photographs of Sb₂O₃ films S1–S4.

Samples S1 (480 °C, 0.4 M) and S2 (500 °C, 0.4 M) showed the highest transparency, with visible-range average transmittances (T_λ380–750) of approximately 78.6% and 77.2%, respectively. Their high transparency is consistent with their phase-pure Sb₂O₃ composition, smaller grain sizes, and the absence of metallic antimony. The relatively low film thicknesses (440–490 nm) can also contribute to reduced scattering and absorption losses. The obtained transmittance values are in good agreement with the reported literature of Sb₂O₃ films.^31,64,65 However a transmittance of 47% was reported for Sb₂O₃ film deposited by spray pyrolysis.⁶⁶ The decline in transparency with increasing temperature and concentration can be correlated with higher film thickness, increased sub-bandgap absorption due to defect states, and enhanced light scattering from larger grain structures at elevated temperatures.^31,64,65

However, the transmittance decreased significantly for samples S3 (540 °C, 0.6 M) and S4 (580 °C, 0.6 M). These films exhibited noticeably reduced visible transmittance values of 50% (S3) and 37% (S4). This loss of transparency might be linked to the combination of structural and compositional factors. The substantially greater film thicknesses can contribute towards the optical scattering. This can also be linked with the larger grain sizes and higher surface roughness as observed in SEM images. Additionally, the formation of metallic Sb inclusions (confirmed by both XRD and XPS) can introduce free-carrier absorption and plasmonic losses that further suppress transmittance.

The optical band gap (E_g) values were calculated using the Tauc method,⁶⁷ and shown in Fig. 6. In this method, the linear region of a plot of (αhν)² versus energy E_g (eV) near the high-energy absorption edge was extrapolated to the x-axis to provide an estimate of the band gap energy. In the Tauc equation (eqn (2)),⁶⁷ α represents the linear absorption coefficient, A is a constant, and hν denotes the photon energy, where h is Planck's constant and ν is the frequency of the incident light.

(αhν)² = A(hν − E_g) (2)

Fig. 6 Tauc plots used for optical band gap determination and haze spectra of Sb₂O₃ films S1–S4.

The band gap was found to vary systematically with deposition temperature and precursor concentration, decreasing from 3.72 eV to 3.47 eV across the series. Samples S1 (3.63 eV) and S2 (3.72 eV), both deposited at lower temperatures (480–500 °C) and at lower precursor concentration, exhibited the highest band gap values. Their wider band gaps are consistent with phase-pure senarmontite Sb₂O₃, relatively low defect densities, and the absence of metallic Sb.

A notable decrease in the band gap was observed for S3 (3.59 eV) and S4 (3.47 eV). These films were deposited at higher temperature and higher precursor concentration and both XRD and XPS confirmed the presence of Sb⁰ species in these samples.

This reduction in the apparent optical band gap with increasing film thickness can be associated with enhanced structural disorder and partial incorporation of metallic Sb at higher deposition temperatures. Thicker films also exhibited larger grain structures and increased optical scattering, which can influence the absorption edge.

The appearance of metallic Sb inclusions in S3 and more prominently in S4 could led to enhanced free carrier absorption and localized plasmonic effects that can lower the band gap values. The progressive reduction from 3.72 eV (S2) to 3.47 eV (S4) therefore reflects the combined influence of increased defect density and metallic antimony incorporation at elevated deposition temperatures. A decrease in band gap from 3.98 eV to 3.71 eV when the film thickness was increased from 0.20 µm to 1.10 µm was previously reported for Sb₂O₃ films deposited by thermal evaporation method.³⁰ Although metallic Sb inclusion was observed in S3 and S4, the oxide matrix remained predominantly Sb₂O_3. Therefore, the fundamental band gap of these samples was found to be close to the reported Sb₂O₃ films.⁶⁵

The haze values in the visible region (380–750 nm), shown in Fig. 6, increased progressively from S1 (7.8%) to S4 (11.9%), indicating a rise in light-scattering behaviour with higher deposition temperature. The low haze of S1 and S2 reflects their smoother, compact microstructures, which minimise scattering. In contrast, S3 shows slightly higher haze (9.07%) due to increased grain growth and surface roughening at 540 °C. The highest haze in S4 corresponds to the formation of large, faceted grains and metallic Sb domains, which strongly scatter incident light. Overall, haze evolution confirms that microstructural coarsening and partial oxide reduction at elevated temperatures degrade optical clarity.

The Sb₂O₃ thin films deposited via AACVD exhibited n-type semiconducting behaviour across all deposition conditions. A clear trend was observed in the electrical transport properties as a function of deposition temperature and precursor concentration. The resistivity decreased systematically with increasing temperature and precursor molarity.

At lower deposition temperatures (S1: 480 °C and S2: 500 °C), the films displayed high resistivity and low carrier concentrations. This behaviour is consistent with their compact but fine-grained microstructure. XPS and XRD showed that those films were purely Sb₂O₃ with no inclusion of Sb⁰. The carbon content of 4–7% originated from incomplete combustion of toluene, leading to carbonaceous residues at grain boundaries that act as trap centres and scattering sites, thereby limiting carrier mobility.⁶¹ The properties of previously reported Sb₂O₃ films deposited by various techniques are summarized in Table 4. Resistivity values lower than those obtained in this work have been reported for Sb₂O₃ films prepared by thermal evaporation³⁰ and spray pyrolysis.⁶⁴

Table 4 Properties of Sb₂O₃ films previously reported by various deposition techniques

Film	Resistivity (Ω cm)	Thickness (nm)	Transmittance (%)	Band gap (eV)	Precursor	Solvent	Deposition method	Ref.
Sb2O3	1.0 × 10^−2	298	74	3.6	SbCl3	Ethanol	Spray pyrolysis	64
Sb2O3	—	—	70	—	SbCl5	Ethyl acetate	APCVD	31
Sb2O3	—	—	—	—	Sb(BuO)3	—	APCVD	39
Sb2O3	6.5 × 10^−2	300	85	—	Sb2O3 powder	—	Thermal evaporation	30
Sb2O3	—	—	47	4.0	SbCl3	Water	Spray pyrolysis	66
Sb2O3	26 × 10^2	189	75	3.43	SbCl3	Acetic acid	Spray pyrolysis	65

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A noticeable reduction in resistivity was observed for films deposited at 540 °C (S3). A decrease in resistivity values with the increase in temperature has also been widely reported.^65,68 While excess carbon content can introduce additional trap states, the conductivity of S3 improved by several orders of magnitude compared to S1 and S2. This increase can be attributed to the first appearance of metallic Sb⁰ features (detected in both XRD and XPS), which suggests partial thermal reduction of Sb₂O₃. Although metallic Sb might not have formed a continuous conductive pathway at this stage, its presence increased local carrier density that led to enhanced electrical transport.

The film deposited at the highest temperature (S4: 580 °C) exhibited the lowest resistivity among the series. This can be linked to its greater film thickness, larger and faceted grain morphology, and extensive formation of metallic Sb inclusions. XPS depth profiling confirmed the coexistence of Sb³⁺ and Sb⁰ throughout the film, with the Sb⁰ fraction decreasing with etching depth. These metallic domains acted as electron-rich regions within the oxide matrix and significantly increased carrier concentration. Despite the large fraction of Sb⁰, the conductivity did not reach metallic behaviour, indicating that the Sb⁰ crystallites remained non-percolating. The mobility of S4 remained modest (∼11 cm² V⁻¹ s⁻¹), likely due to enhanced scattering at interfaces between Sb₂O₃ and metallic Sb, as well as increased microstructural roughness, consistent with SEM observations.

Superior transport properties such as high carrier concentration are often achieved at the cost of reduced optical transmittance. To assess the optoelectronic performance of TCO materials the figure of merit (FOM) proposed by Haacke in 1976⁶⁹ was employed in this study. The FOM (Φ) is defined in eqn (3).

The low-temperature films S1 and S2 exhibited very low FOM values (0.02 × 10⁻⁶ and 0.07 × 10⁻⁶ Ω⁻¹, respectively), reflecting their high sheet resistances despite high visible transmittance. A pronounced improvement was observed for S3 (540 °C, 0.6 M), which achieved the highest FOM of 1.98 × 10⁻⁶ Ω⁻¹ due to its significantly reduced sheet resistance (11 000 Ω □⁻¹) and moderate transparency. In contrast, S4 exhibited a markedly reduced FOM (0.03 × 10⁻⁶ Ω⁻¹), as its substantial loss in transmittance (caused by large metallic Sb crystallites formed at 580 °C) outweighed the benefits of its low sheet resistance. These trends demonstrate that although higher deposition temperature improved conductivity, excessive thermal reduction diminished transparency. Such oxide-to-metal transitions at high deposition temperatures have previously been reported when antimony(III) ethoxide was used as a precursor.^39,40

Surface wettability

The water contact angles of the films (Fig. 7) increased from S1 (44 ± 2°) to S3 (65 ± 2°), reflecting the combined effect of surface chemistry and morphological changes with increasing deposition temperature. Films S1 and S2, which exhibited the lowest carbon content (4–7 at%), showed moderately hydrophilic behaviour. However, S3 showed the highest contact angle (65 ± 4°), correlating with its significantly higher surface carbon (26 at%) and the presence of metallic Sb crystallites. Higher surface carbon content has long been reported to reduce surface energy and promote hydrophobicity.^70,71 Interestingly, S4 displayed a slightly lower angle (59 ± 5°) despite its high roughness and metallic Sb inclusions. This can be attributed to its lower carbon content relative to S3 (21 at%), and the more complex mixture of oxide and metal facets that can introduce local polar sites capable of hydrogen bonding.^72,73 Although hydroxyl surface species generally promote hydrophilic behaviour due to their polar nature, the wettability behaviour observed in S3 and S4 appeared to be dominated by the increased surface carbon content and metallic Sb inclusions. The higher carbon concentration likely masked the hydrophilic contribution from surface hydroxyl groups, resulting in the observed increase in water contact angle.

Fig. 7 Static water contact angle measurements of Sb₂O₃ films S1–S4.

Conclusions

Antimony(III) oxide films were successfully deposited by AACVD. The structure and properties of films were strongly influenced by the deposition temperature and precursor concentration, which also affected the film thickness. At lower deposition temperatures (480–500 °C) the films were purely Sb₂O₃ and showed weaker charge transport properties. At 540–580 °C the films underwent a reductive phase transformation that generated metallic Sb alongside Sb₂O₃. The inclusion of metallic Sb improved electrical conductivity, as seen by the drop in resistivity from 2.1 × 10² Ω cm (S1) to 1.1 × 10⁻¹ Ω cm (S4) and the rise in carrier concentration from 3.7 × 10¹⁵ to 5.2 × 10¹⁸ cm⁻³. Mobility values remained within 8–14 cm² V⁻¹ s⁻¹, consistent with transport limited by grain boundaries and defect scattering. The optical performance followed the structural changes. Increasing film thickness at higher deposition temperatures and precursor concentrations contributed to larger grain structures, enhanced conductivity, lower optical transparency and increased haze. Increased grain size, film thickness, and metallic Sb inclusions reduced transparency from ∼79% to ∼36% and raised haze from ∼8% to ∼12% due to stronger light scattering. Despite these losses, S3 (540 °C) offered the best compromise between transparency and conductivity, giving the highest Haacke figure of merit. Water contact angles also reflected the combined influence of surface chemistry and morphology. Its values increased from 44° to 65° due to higher carbon content and metallic Sb features that enhanced hydrophobicity. Overall, intermediate deposition temperatures produced the most balanced films, combining good crystallinity, improved conductivity, and acceptable transparency. Controlling the degree of thermal reduction to Sb° is crucial, since limited Sb formation enhanced charge transport, whereas excessive metallic content reduced optical clarity and altered surface chemistry. This balance enabled effective optimisation of Sb₂O₃ films suitable for optoelectronic and sensing applications.

Author contributions

IR prepared and characterised the samples and contributed to the writing and editing of the manuscript. CJC and IPP supervised the project and contributed to scientific discussions, writing and editing of the manuscript. All authors participated in scientific discussions throughout the study.

Conflicts of interest

The authors declare no conflict of interest.

Data availability

The data that support the findings of this study are included within the article and the supplementary information (SI). The supplementary information contains additional XPS characterisation data, including the XPS survey spectrum of sample S2, deconvoluted Sb 3d spectra and XPS depth-profiling spectra of sample S1. See DOI: https://doi.org/10.1039/d6ma00123h.

Additional data are available from the corresponding author upon reasonable request.

Acknowledgements

Iqra Ramzan is thankful for the funding received for this research from NSG Pilkington Glass Ltd. and UCL Dean's prize. A part of this work was also supported by the Engineering and Physical Sciences Research Council (EPSRC) (Grant Number: EP/W010798/1). The authors would also like to thank NSG Pilkington Glass Ltd. for providing glass substrates for research conducted in this work. Thanks to Jamie Gould for the valuable discussions and insights on XRD.

References

1. A. Fernandez and M. Merino, Preparation and characterization of Sb₂Se₃ thin films prepared by electrodeposition for photovoltaic applications, Thin Solid Films, 2000, 366(1–2), 202–206.
2. A. Salem and M. S. Selim, Structure and optical properties of chemically deposited Sb2S3 thin films., J. Phys. D: Appl. Phys., 2001, 34(1), 12.
3. H. Li, K. Wang, M. Zhou, W. Li, H. Tao, R. Wang, S. Cheng and K. Jiang, Facile tailoring of multidimensional nanostructured Sb for sodium storage applications., ACS Nano, 2019, 13(8), 9533–9540.
4. R. Kudrawiec and D. Hommel, Bandgap engineering in III-nitrides with boron and group V elements: Toward applications in ultraviolet emitters., Appl. Phys. Rev., 2020, 7(4), 041314.
5. Z. Wu and J. Hao, Electrical transport properties in group-V elemental ultrathin 2D layers, npj 2D Mater. Appl., 2020, 4(1), 4.
6. A. Nagaoka, S. K. Swain and A. H. Munshi, Review on Group-V Doping in CdTe for Photovoltaic Application., IEEE J. Photovolt., 2024, 397–413.
7. A. Jabeen, A. Majid, M. Alkhedher, S. Haider and M. S. Akhtar, Impacts of structural downscaling of inorganic molecular crystals-A DFT study of Sb₂O₃., Mater. Sci. Semicond. Process., 2023, 166, 107729.
8. N. Tigau, V. Ciupina and G. Prodan, The effect of substrate temperature on the optical properties of polycrystalline Sb₂O₃ thin films., J. Cryst. Grow., 2005, 277(1–4), 529–535.
9. D. Zeng, X. Chen, R. Jiang, C. Xie, B. Zhu and W. Song, Thermal stability of Sb/Sb₂O₃ composite nanoparticles., Mater. Chem. Phys., 2006, 96(2–3), 454–458.
10. R. Orman and D. Holland, Thermal phase transitions in antimony(III) oxides., J. Solid State Chem., 2007, 180(9), 2587–2596.
11. H. Shiel, T. D. Hobson, O. S. Hutter, L. J. Phillips, M. J. Smiles, L. A. Jones, T. J. Featherstone, J. E. Swallow, P. K. Thakur and T.-L. Lee, Band alignment of Sb₂O₃ and Sb₂Se₃., J. Appl. Phys., 2021, 129(23), 235301.
12. H. Qu, W. Wu, Y. Zheng, J. Xie and J. Xu, Synergistic effects of inorganic tin compounds and Sb₂O₃ on thermal properties and flame retardancy of flexible poly (vinyl chloride)., Fire Safety J., 2011, 46(7), 462–467.
13. L. Niu, J. Xu, W. Yang, J. Zhao, J. Su, Y. Guo and X. Liu, Research on nano-Sb₂O₃ flame retardant in char formation of PBT., Ferroelectrics, 2018, 523(1), 14–21.
14. A. Garcia, Transparent antimony based flame retardants for vinyl coated fabrics., J. Coat. Fabrics, 1982, 11(3), 137–142.
15. Y. Akinay, H. C. Kazici, I. N. Akkuş and F. Salman, Synthesis of 3D Sn doped Sb₂O₃ catalysts with different morphologies and their effects on the electrocatalytic hydrogen evolution reaction in acidic medium., Ceram. Int., 2021, 47(20), 29515–29524.
16. D.-N. Liu, G.-H. He, L. Zhu, W.-Y. Zhou and Y.-H. Xu, Enhancement of photocatalytic activity of TiO2 nanoparticles by coupling Sb₂O₃., Appl. Surf. Sci., 2012, 258(20), 8055–8060.
17. M. Abdellatif, Y. Louafi, D. Nunes, T. Freire, E. Fortunato, R. Martins, S. Kabouche and M. Trari, Studies on photocatalytic degradation of Rhodamine B using the valentinite Sb₂O₃., React. Kinet., Mech. Catal., 2023, 136(3), 1643–1655.
18. R. Makhloufi, S. E. Hachani, Z. Zekri and W. Tair, Solvothermal Synthesis of Antimony Trioxide Sb₂O₃ Used as a Photocatalyst for Crystal Violet Dye Degradation., Moscow Univ. Chem. Bull., 2022, 77(2), 111–116.
19. H. Liu and Y. Iwasawa, Unique performance and characterization of a crystalline SbRe2O6 catalyst for selective ammoxidation of isobutane., J. Phys. Chem. B, 2002, 106(9), 2319–2329.
20. H. Liu, H. Imoto, T. Shido and Y. Iwasawa, Selective ammoxidation of isobutylene to methacrylonitrile on a new family of crystalline Re–Sb–O catalysts., J. Catal., 2001, 200(1), 69–78.
21. L. Zuo, X. Jiang, L. Yang, M. Xu, Y. Nan, Q. Yan and H. Chen, Sb₂O₃ anode buffer induced morphology improvement in small molecule organic solar cells., Appl. Phys. Lett., 2011, 99(18), 183306.
22. M. I. Tareq and B. A. Hasan, Characterization and Photovoltaic Effect of (Sb₂O₃: Metal Oxides)/c-Si Heterojunctions, Iraqi J. Phys., 2023, 21(4), 92–102.
23. H. R. Yeom, S. Song, S. Y. Park, H. S. Ryu, J. W. Kim, J. Heo, H. W. Cho, B. Walker, S.-J. Ko and H. Y. Woo, Aesthetic and colorful: Dichroic polymer solar cells using high-performance Fabry-Perot etalon electrodes with a unique Sb₂O₃ cavity., Nano Energy, 2020, 77, 105146.
24. C. Song, N. Zhang, J. Lin, X. Guo and X. Liu, Sb₂O₃/Ag/Sb₂O₃ multilayer transparent conducting films for ultraviolet organic light-emitting diode., Sci. Rep., 2017, 7(1), 41250.
25. N. Tigau, V. Ciupina, G. Prodan, G. Rusu, C. Gheorghies and E. Vasile, The influence of heat treatment on the electrical conductivity of antimony trioxide thin films., J. Optoelectron. Adv. Mater., 2003, 5(4), 907–912.
26. K. Divya and K. Abraham, Multifunctional transition metal doped Sb₂O₃ thin film with high near-IR transmittance, anti-reflectance and UV blocking features., Appl. Surf. Sci., 2019, 493, 1115–1124.
27. C. Svensson, Refinement of the crystal structure of cubic antimony trioxide, Sb₂O₃., Acta Crystallogr. Sect. B: Struct. Crystallogr. Cryst. Chem., 1975, 31(8), 2016–2018.
28. C. Svensson, The crystal structure of orthorhombic antimony trioxide, Sb₂O₃., Acta Crystallogr. Sect. B: Struct. Crystallogr. Cryst. Chem., 1974, 30(2), 458–461.
29. N. Tigau, Preparation and characterization of Sb∼ 2O∼ 3 thin films., Romanian J. Phys., 2006, 51(5/6), 641.
30. N. Tigau, V. Ciupina and G. Prodan, Structural, optical and electrical properties of Sb∼ 2O∼ 3 thin films with different thickness., J. Optoelectron. Adv. Mater., 2006, 8(1), 37.
31. R. Binions, C. J. Carmalt and I. P. Parkin, Antimony oxide thin films from the atmospheric pressure chemical vapour deposition reaction of antimony pentachloride and ethyl acetate., Polyhedron, 2006, 25(15), 3032–3038.
32. M. N. Chaudhari, R. B. Ahirrao and S. D. Bagul, Thin film deposition methods: A critical review., Int. J. Res. Appl. Sci. Eng. Technol., 2021, 9(6), 5215–5232.
33. I. Ramzan, J. Borowiec, I. P. Parkin and C. J. Carmalt, Transparent Conductive Titanium and Fluorine Co-doped Zinc Oxide Films, ChemPlusChem, 2024, 89(8), e202400073.
34. D. S. Bhachu, G. Sankar and I. P. Parkin, Aerosol assisted chemical vapor deposition of transparent conductive zinc oxide films., Chem. Mater., 2012, 24(24), 4704–4710.
35. A. Jiamprasertboon, S. C. Dixon, S. Sathasivam, M. J. Powell, Y. Lu, T. Siritanon and C. J. Carmalt, Low-cost one-step fabrication of highly conductive ZnO: Cl transparent thin films with tunable photocatalytic properties via aerosol-assisted chemical vapor deposition., ACS Appl. Electron. Mater., 2019, 1(8), 1408–1417.
36. Z. Jarzebski and J. Marton, Physical properties of SnO2 materials: I. preparation and defect structure., J. Electrochem. Soc., 1976, 123(7), 199C.
37. M. Epifani, J. D. Prades, E. Comini, E. Pellicer, M. Avella, P. Siciliano, G. Faglia, A. Cirera, R. Scotti and F. Morazzoni, The role of surface oxygen vacancies in the NO2 sensing properties of SnO2 nanocrystals., J. Phys. Chem. C, 2008, 112(49), 19540–19546.
38. K. Momma and F. Izumi, VESTA 3 for three-dimensional visualization of crystal, volumetric and morphology data., J. Appl. Crystallogr., 2011, 44(6), 1272–1276.
39. P. W. Haycock, G. A. Horley, K. C. Molloy, C. P. Myers, S. A. Rushworth and L. M. Smith, Atmospheric pressure CVD of antimony oxides., Chem. Vap. Depos., 2001, 7(5), 191–193.
40. S. He, A. Bahrami, X. Zhang, J. Julin, M. Laitinen and K. Nielsch, Low-temperature ALD of highly conductive antimony films through the reaction of silylamide with alkoxide and alkylamide precursors., Mater. Today Chem., 2023, 32, 101650.
41. B. D. Cullity, Elements of X-ray Diffraction, Addison-Wesley Publishing, 1956.
42. N. Tigau, V. Ciupina, G. Prodan, G. Rusu and E. Vasile, Structural characterization of polycrystalline Sb₂O₃ thin films prepared by thermal vacuum evaporation technique., J. Cryst. Grow., 2004, 269(2–4), 392–400.
43. J. Chastain and R. C. King Jr, Handbook of X-ray photoelectron spectroscopy., Perkin-Elmer Corporation, 1992, 40, 221.
44. R. Izquierdo, E. Sacher and A. Yelon, X-ray photoelectron spectra of antimony oxides., Appl. Surf. Sci., 1989, 40(1–2), 175–177.
45. A. Rastogi and K. Reddy, Growth of dielectric layers on the InSb surface., Thin Solid Films, 1995, 270(1–2), 616–620.
46. R. L. McCreery and A. Bard, Electroanalytical chemistry, ed. A. J. Bard, 1991, vol. 17, p. 221.
47. B. Kalkofen, V. M. Mothukuru, M. Klingsporn and E. P. Burte, Investigation of antimony oxide films deposited by atomic layer deposition., ECS Trans., 2012, 45(3), 461.
48. R. Reiche, J. Holgado, F. Yubero, J. Espinos and A. Gonzalez-Elipe, Characterization of Sb₂O₃ subjected to different ion and plasma surface treatments., Surf. Interface Anal., 2003, 35(3), 256–262.
49. D. J. Morgan, Metallic antimony (Sb) by XPS., Surf. Sci. Spectra, 2017, 24(2), 024004.
50. L. Santinacci, G. Sproule, S. Moisa, D. Landheer, X. Wu, A. Banu, T. Djenizian, P. Schmuki and M. Graham, Growth and characterization of thin anodic oxide films on n-InSb (1 0 0) formed in aqueous solutions., Corros. Sci., 2004, 46(8), 2067–2079.
51. H. Bryngelsson, J. Eskhult, L. Nyholm, M. Herranen, O. Alm and K. Edström, Electrodeposited Sb and Sb/Sb₂O₃ nanoparticle coatings as anode materials for Li-ion batteries., Chem. Mater., 2007, 19(5), 1170–1180.
52. O. E. L. Pérez, M. D. Sánchez and M. L. Teijelo, Characterization of growth of anodic antimony oxide films by ellipsometry and XPS., J. Electroanal. Chem., 2010, 645(2), 143–148.
53. W. Qi, S. Guo, H. Sun, Q. Liu, H. Hu, P. Liu, W. Lin and M. Zhang, Synthesis and characterization of Sb₂O₃ nanoparticles by liquid phase method under acidic condition., J. Cryst. Growth, 2022, 588, 126642.
54. G. Fan, Z. Huang, C. Chai and D. Liao, Synthesis of micro-sized Sb₂O₃ hierarchical structures by carbothermal reduction method., Mater. Lett., 2011, 65(8), 1141–1144.
55. F. Larbi, S. Bourahla, S. Kouadri Moustefai and F. Elagra, Impact of fluorine and chlorine doping on the structural, electronic, and optical properties of SnO2: first-principles study., Can. J. Phys., 2023, 101(9), 496–503.
56. S. Akbar, S. Hasanain, O. Ivashenko, M. Dutka, N. Akhtar, J. T. M. De Hosson, N. Ali and P. Rudolf, Defect ferromagnetism in SnO 2: Zn 2+ hierarchical nanostructures: correlation between structural, electronic and magnetic properties., RSC Adv., 2019, 9(7), 4082–4091.
57. X. Chen, L. Wang, F. Ma, T. Wang, J. Han, Y. Huang and Q. Li, Core@ shell Sb@ Sb 2 O 3 nanoparticles anchored on 3D nitrogen-doped carbon nanosheets as advanced anode materials for Li-ion batteries., Nanoscale Adv., 2020, 2(12), 5578–5583.
58. I. A. Hassan, A. Ratnasothy, D. S. Bhachu, S. Sathasivam and C. J. Carmalt, The effect of solvent on the morphology of indium oxide deposited by aerosol-assisted chemical vapour deposition., Aust. J. Chem., 2013, 66(10), 1274–1280.
59. D. B. Potter, I. P. Parkin and C. J. Carmalt, The effect of solvent on Al-doped ZnO thin films deposited via aerosol assisted CVD., RSC Adv., 2018, 8(58), 33164–33173.
60. G. Niu, H.-D. Kim, R. Roelofs, E. Perez, M. A. Schubert, P. Zaumseil, I. Costina and C. Wenger, Material insights of HfO2-based integrated 1-transistor-1-resistor resistive random access memory devices processed by batch atomic layer deposition., Sci. Rep., 2016, 6(1), 28155.
61. S. Kim, S.-H. Lee, I. H. Jo, J. Seo, Y.-E. Yoo and J. H. Kim, Influence of growth temperature on dielectric strength of Al2O3 thin films prepared via atomic layer deposition at low temperature., Sci. Rep., 2022, 12(1), 5124.
62. M. Maleki and S. Rozati, An economic CVD technique for pure SnO 2 thin films deposition: Temperature effects., Bull. Mater. Sci., 2013, 36(2), 217–221.
63. N. Chen, I. Ramzan, S. Li and C. J. Carmalt, Synthesis and Characterization of Optically Transparent and Electrically Conductive Mo-Doped ZnO, F-Doped ZnO, and Mo/F-Codoped ZnO Thin Films via Aerosol-Assisted Chemical Vapor Deposition., Cryst. Growth Des., 2024, 24(24), 10256–10266.
64. A. S. Hassanien and I. El Radaf, Effect of fluorine doping on the structural, optical, and electrical properties of spray deposited Sb₂O₃ thin films., Mater. Sci. Semicond. Process., 2023, 160, 107405.
65. Y. Shinde, P. Sonone, R. Kendale, P. Koinkar and A. Ubale, Engineering of physical properties of spray-deposited nanocrystalline Sb₂O₃ thin films by phase transformation., Nanotechnology, 2020, 32(2), 025602.
66. H. Slimani, A. C. Nkele, S. Dagher, A. Alshoaibi, N. Bessous, B. Akhozheya and F. I. Ezema, Investigating the optical properties of antimony oxide, Sb₂O₃ nanomaterials synthesized by spray pyrolysis., Opt. Mater., 2024, 157, 116405.
67. M. Fadavieslam, H. Azimi-Juybari and M. Marashi, Dependence of O 2, N 2 flow rate and deposition time on the structural, electrical and optical properties of SnO 2 thin films deposited by atmospheric pressure chemical vapor deposition (APCVD)., J. Mater. Sci.: Mater. Electron., 2016, 27, 921–930.
68. R. A. Gonçalves, M. R. Baldan, A. J. Chiquito and O. M. Berengue, Synthesis of orthorhombic Sb₂O₃ branched rods by a vapor–solid approach., Nano-Struct. Nano-Obj., 2018, 16, 127–133.
69. G. Haacke, New figure of merit for transparent conductors., J. Appl. Phys., 1976, 47(9), 4086–4089.
70. Y. Ju, L. Ai, X. Qi, J. Li and W. Song, Review on hydrophobic thin films prepared using magnetron sputtering deposition., Materials, 2023, 16(10), 3764.
71. J. K. Patel, I. Ramzan, C. I. De Leon Reyes, I. P. Parkin and C. J. Carmalt, Tetrakis (2, 2, 6, 6-tetramethyl-3, 5-heptanedionate) Cerium for the Deposition of Hydrophobic Coatings, Langmuir, 2025, 15562–15572.
72. Z. Zhu, Z. Yu, F. F. Yun, D. Pan, Y. Tian, L. Jiang and X. Wang, Crystal face dependent intrinsic wettability of metal oxide surfaces, Natl. Sci. Rev., 2021, 8(1), nwaa166.
73. A. Feng, B. J. McCoy, Z. A. Munir and D. Cagliostro, Wettability of transition metal oxide surfaces., Mater. Sci. Eng., A, 1998, 242(1–2), 50–56.

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