Structural, Optical, Band Edge and Enhanced Photoelectrochemical Water Splitting Properties of Tin-Doped WO₃

Abstract

The substitutional doping of tungsten oxide (WO₃) with metal ions demonstrates a promising approach to enhance its photoelectrochemical (PEC) water splitting efficiency. In this article, the substitutional doping of Sn ions into WO₃ lattice and its effect on optical, electrical, band edge, and PEC water splitting properties are explored. Sn-doped WO₃ thin films were synthesized using a facile hydrothermal method. The characterization data reveal that the doping of Sn alters the morphology, induces multiple crystal phases, effects the crystal orientation, reduces the band gap, and increases the carrier density of WO₃. With the uniform distribution of Sn ions in WO₃ and the decreased charge transfer resistance at the electrode/electrolyte interface, the doped WO₃ show notable enhancement in its PEC activity compared to the undoped WO₃. The band edge study revealed that the introduction of Sn in WO₃ lattice causes an increase in the energy distance between the valence band edge and Fermi level and, at the same time, induces a downward shift in both the valence and conduction band edges towards higher potentials with respect to reversible hydrogen electrode (RHE). Conclusively, this work shows significant and new insights about Sn-doped WO₃ photoanodes and their influence on PEC water splitting efficiency.

1. Introduction

A promising resolution to the global energy requirement of the future, along with the protection of environment, comes from the next generation and sustainable energy source hydrogen and its clean and carbon-free production via photoelectrochemical water splitting [1]. In photoelectrochemical (PEC) water splitting, a semiconductor is used to split water into H₂ and O₂ using abundant and sustainable resources, like sunlight and water [2,3]. The first successful demonstration of solar water splitting using TiO₂ semiconductor was reported by Fujishima and Honda [4], and has initiated significant research attention and contributions that have continued until now. As a result, a variety of semiconductors were explored, researched, and reported as an active material for solar water splitting reaction. Notably, the well-matched optical, electrical properties, and stability of the metal oxides have demonstrated their active participation as useful electrode materials in PEC water splitting systems.

Among the various metal oxide materials utilized in PEC water splitting, tungsten oxide (WO₃) has emerged as a promising candidate, due to its suitable physical and chemical properties [5,6]. The effective utilization of WO₃ in solar water splitting was first demonstrated in 1976 by Hodes et al. [7]. Since then, significant research contributions have been reported. It was found that WO₃ can absorb a maximum of ~12% of the incident solar light with a band gap of 2.6–2.8 eV. Ideally, based on its light absorption capability, as high as ~4.8% of solar energy to the hydrogen conversion efficiency could be extracted [8]. However, WO₃ faces numerous limitations that affect its charge generation, separation, and transportation properties, that decrease its ability to carry out water splitting. Therefore, the structural, optical, and electrochemical properties of WO₃ need to be tuned to increase its water splitting efficiency.

The tuning of the structural, optical, and electrochemical properties of WO₃ could be achieved by various modifications that improve its PEC water splitting efficiency. In particular, the crystal structure, band gap, carrier density, and charge transport at the electrolyte interface are the most significant factors influencing water splitting efficiency. Among the existing crystal structures, monoclinic WO₃ is known to be the most suitable crystal phase for PEC water splitting [9,10,11]. Therefore, it is essential to explore the possible modification through which the structural, optical, and electrical properties of monoclinic WO₃ could be tuned and optimized without disrupting the crystal nature and stability. Hence, an effectively tuned monoclinic WO₃ is expected to be beneficial for the water splitting reaction, compared to undoped WO₃.

A recent study indicates that the substitutional doping of metal ions is an effective strategy of tuning the physical and chemical properties of WO₃ [12,13,14,15]. Importantly, the type and extent of modification induced in doped WO₃ mainly depend on the nature of dopant metal ion and its concentration [13,16,17]. Therefore, exploring a suitable dopant material for WO₃ is essential for obtaining high efficiency in PEC water splitting [3]. Until now, various dopants, including Ti [16], Co [18], Mo [17], Zn [19], Bi [13], Ni [18], Nb [20], Cu [18], Al [21], Ta [22], Fe [23,24,25], Te [26], and Ba [27], have been employed to alter the structural, optical, electrical, and band edge properties of WO₃ and to enhance its photocatalytic properties. Previously, Sn has been actively used as an effective dopant material for metal oxides, such as Fe₂O₃, to enhance its PEC water splitting properties [28,29], revealing its positive impact as a dopant material. Moreover, a detailed literature survey revealed that despite a few reports, no in-depth study has been reported regarding the synthesis and PEC properties Sn-doped WO₃ thin films. Hence, it is essential to explore the effect of Sn doping on the structural, optical, electrical, band edge, and PEC water splitting properties of WO₃.

Given this, we here synthesize and explore the effect of Sn doping on the PEC water splitting efficiency of WO₃ thin films. To obtain the Sn-doped WO₃ thin films, a facile single-step hydrothermal approach is utilized, that is proceeded by introducing the Sn precursor into a peroxopolytungstic acid (PTA) solution, followed by reaction at 150 °C. The Sn-doped WO₃ thin films were characterized in detail to understand the effect on structural, crystallographic, optical, and electrical properties of WO₃. Furthermore, Sn ions were found to be uniformly distributed in the WO₃ crystals on the (fluorine doped tin oxide) FTO substrate, without affecting their adhesion. Using data from the absorbance, X-ray photoelectron spectroscopy, and Mott–Schottky (MS) analysis, an appropriate band edge positions of WO₃ was plotted with respect to RHE potentials before and after doping with Sn ions. The electrochemical measurements revealed that Sn-doped WO₃ thin films show increased photocurrent (at 1.23 V vs. RHE), incident photon-to-current efficiency (IPCE) and decreased charge transfer resistance across the electrode/electrolyte interface. The notable enhancement in the PEC properties of WO₃ upon Sn doping was ascribed to the proper crystal facet orientation, decreased band gap, and significant increase of carrier density.

2. Results and Discussion

The hydrothermal synthesis method utilized in the present study allows for the precise tuning of Sn atom % in the WO₃ lattice by controlling the Sn precursor concentration in the hydrothermal synthesis solution. Table S1 shows the amount of precursor solution taken in the synthesis solution and the atom % and wt % of Sn detected via EDS. Based on the PEC studies, the WO₃ thin films doped with 0.8 atom % of Sn showed enhanced photocurrent compared to the samples doped with the other concentrations of Sn, and hence, were chosen for detailed characterization. Figure 1 shows the high- and low-resolution SEM images of undoped and Sn-doped WO₃ thin films. The morphology of undoped WO₃ (Figure 1a,b) appeared to be a mixture of nanoparticles, and nanorod structure assembled on the FTO substrate. After Sn doping, a significant change in morphology was observed, indicating the disappearance of nanorod morphology and the formation of smaller (5–10 nm) coagulated nanoparticles (Figure 1d,e) and the presence of larger (500–1000 nm) nanoparticles and plates at different places. This suggests that the substitution of Sn in crystal lattice affects the structure of WO₃. The cross-section SEM image of undoped WO₃ (Figure 1c) clearly shows the nanorod morphology assembled on the thin film of the nanoparticle layer. The thickness of the thin film was measured to be ~2 µm. Figure 1f shows the cross-section SEM image of Sn-doped WO₃ thin film, confirming the change in morphology and the decrease in film thickness to 1 µm. Furthermore, the low-resolution SEM image confirmed the uniform deposition of doped WO₃ throughout the FTO substrate with a reduced thickness, indicating that the doping of Sn affects the deposition amount of WO₃ on the substrate.

It is well known that the WO₃ thin films synthesized by hydrothermal treatment of PTA solution produce orthorhombic WO₃·0.33H₂O that is converted to anhydrous monoclinic WO₃ upon annealing at 500 °C. Figure 2a confirms that the synthesized undoped WO₃ thin films (annealed at 500 °C) showed monoclinic crystal structure (with lattice parameters of a = 0.729 nm, b = 0.753 nm, c = 0.767 nm, and β = 90.92°) by producing three characteristics peaks related to [002], [020], and [200] plane at 2θ degree value of 23.18, 23.63, and 24.41 (JCPDS no. 43−1035), respectively. In Figure 2, the peaks arising from the FTO substrate (SnO₂) are indicated with black dots, whereas the monoclinic peaks of WO₃ are marked with green star marks. Moreover, due to the dominant appearance of substrate SnO₂ peaks, the diffraction lines related to the very low concentration of doped SnO₂ peaks were expected to merge and are indistinguishable. Notably, the XRD peaks of the FTO substrate appeared comparatively intense in doped WO₃ compared to undoped WO₃, due to the reduced thickness of WO₃ film in the Sn-doped samples, as observed in Figure 1f. Careful observation of the XRD data of Sn-doped WO₃ thin films (Figure 2b) revealed few critical changes in the diffraction patterns. Firstly, the peak positions of [002], [020], and [200] planes of monoclinic WO₃ (with lattice parameters of a = 0.729 nm, b = 0.753 nm, c = 0.767 nm, and β = 90.92°) show a significant shift towards higher 2θ degree values. That is, the peak positions of [002], [020], and [200] planes were found to be shifted to 23.39, 23.80, and 24.56 of 2θ degree values, respectively, after Sn doping (Figure S1). This indicates that the doping of Sn takes places by displacing a few of the W ions from their lattice positions.

The most significant change observed in the XRD pattern of WO₃ after Sn doping was in the peak intensities related to [002], [020], and [200] planes. The peak height associated with the different planes in the XRD pattern suggests the crystal growth direction. Interestingly, the PEC water splitting efficiency of WO₃ depends on the crystal plane orientation [30]. As reported previously, a [002] plane-oriented WO₃ crystal is beneficial, and performs an efficient water oxidation reaction at the surface and eventually shows higher PEC water splitting efficiency compared to the [020] and [200] plane-oriented crystals. The XRD pattern of undoped WO₃ reveals the dominant and most intense peak of the [200] plane, which is reported to be a non-suitable orientation for the PEC water splitting reaction. However, upon Sn doping, a significant decrease in [200] peak and an increase in [002] was observed. This indicates that doping of Sn promotes the preferred growth direction of WO₃ crystals towards the [002] orientation, from the [200] plane, and could influence its PEC water splitting efficiency. In addition to the above changes, the doping of Sn tends to induce multiple phases in WO₃ as noted from the appearance of new XRD peak at a low diffraction angle, which is ascribed to the hexagonal phase of WO₃. The presence of a peak related to the hexagonal phase of WO₃ was due to the reconstructive transformation process that occurs during the annealing step.

As described in earlier reports, the hydrothermal treatment of PTA solution produces WO₃·0.33H₂O (orthorhombic phase), that undergoes a reconstructive transformation during the calcination step [31]. In particular, the reconstructive transformation involves temperature dependence, breaking of bonds, elimination of a water molecules, and a change in crystal structure with the formation of new bonds [31]. Briefly, annealing of WO₃·0.33H₂O at 500 °C transforms the orthorhombic phase to hexagonal phase (at 400 °C), and then to stable monoclinic phase (at 500 °C). However, the doping of Sn ions in WO₃·0.33H₂O (orthorhombic structure) will result in obstruction of the reconstructive transformation during calcination. Hence, the partial transformation from hexagonal to monoclinic phase takes place, resulting in multiple structures in the product. In fact, the interference of doped metal ions during reconstructive transformation, causing multiple phases of WO₃, has already been observed before [16].

The low-resolution TEM images shown in Figure 3a,c indicate the nanorod and particle morphology of undoped and Sn-doped WO₃ thin films, respectively. In the high-resolution TEM image of undoped WO₃, a d-spacing value of 0.3645 was calculated (Figure 3b) which was ascribed to the [200] plane of monoclinic WO₃. The d-spacing value of Sn-doped WO₃ was estimated to be 0.3634 nm (Figure 3d), which was noted to be slightly less than undoped WO₃. The slight decrease in d-spacing value agrees with the shift in diffraction angles to the higher degree observed in the XRD patterns after Sn doping, and hence confirms the insertion of Sn ions into the WO₃ lattice. The distribution of Sn ions in WO₃ lattice was observed by collecting elemental mapping data attached to the TEM. In the selected area (Figure 4a) of the Sn-doped WO₃ samples, the elemental mapping showed a uniform distribution of W (Figure 4b), O (Figure 4c), and Sn (Figure 4d), and the EDS measurement (Figure 4e) confirmed atom % and wt % of individual elements and the absence of impurities. The UV–vis absorption spectra of undoped WO₃ shows the absorption onset of 470 nm, as shown in Figure 5a. Upon Sn doping, a very slight redshift in absorption onset to 500 nm was observed (Figure 5a). In addition, the doping of Sn tends to cause tailing of absorbance onset into the visible region, indicating an increase in oxygen vacancies. The Tauc plot shown in Figure 5b confirms the reduction in band gap from 2.62 to 2.46 eV for undoped WO₃ and Sn-doped WO₃ thin films, respectively.

The XPS survey spectrum of Sn-doped WO₃ thin films is shown in Figure 6a, indicating the presence of W, O, and Sn without any impurities. The doping of metal ions is reported to alter the concentration of oxygen vacancies in WO₃. Hence, the high-resolution W 4f core-level XPS spectra were deconvoluted as shown in Figure 6b,c. In the high-resolution XPS W 4f core-level spectra, two major peaks were detected, and were assigned to W⁶⁺ state (at 35.25 and 37.40 eV of binding energy) of 4 f_7/2 and 4 f_5/2. In addition to the two major peaks mentioned above, two low-intensity peaks were obtained that are ascribed to the presence of W⁵⁺ state (at binding energy values of 33.95 and 36.10 eV), and hence, the presence of oxygen vacancies (i.e., causing a partial reduction of W⁶⁺) were confirmed. Similarly, the deconvolution of W 4f core-level XPS spectra of Sn-doped WO₃ revealed two doublets each for W⁶⁺ and W⁵⁺ state. However, careful observation revealed that the peak intensity of W⁵⁺ state shows a significant increase after Sn doping. Precisely, the peak area of W⁵⁺ state in undoped WO₃ was calculated to be 0.82% and 1.12%, corresponding to W 4f_5/2 and W 4f_7/2 states, respectively, and was found to be increased to 6.22% and 4.68% with respect to W 4f_5/2 and W 4f_7/2 state, respectively, for the Sn-doped WO₃. This indicates a drastic increase in the oxygen vacancies in WO₃ after Sn doping. This could be due to the change in the valence state of dopant Sn ion and the host W ions. Furthermore, in the high-resolution XPS spectra of W 4f, a significant shift in the peak positions of W⁶⁺ and W⁵⁺ states towards higher binding energy was observed. This indicates that the doping of Sn takes place by replacing W atoms via substitutional doping [32]. The high-resolution XPS spectra of Sn shows (Figure 6d) two peaks at binding energy values of 485.5 and 494.5eV, with spin–orbit components of 8.4 eV, corresponding to 3 d_3/2 and 3 d_5/2 core levels. Furthermore, the peak fitting deconvolution produced two doublets corresponding to the Sn(II) and Sn(IV) states. However, the presence of Sn(II) was observed to be dominant compared to the Sn(IV) state [33].

The Fermi-level position of WO₃ is sensitive to its crystal environment and constituent atoms. Hence, doping WO₃ tends to alter its Fermi energy level with respect to valence and conduction band edge positions. Using XPS valence band edge analysis, the energy difference between the Fermi level and the valence band edge was determined for both the undoped and Sn-doped WO₃. Figure 7a,b shows the XPS valence band edge plot. The undoped WO₃ show 2.07 eV energy difference between the valence band and Fermi level that increased to 2.21 eV upon Sn doping. The significant increase in energy difference tends to push the Fermi level towards and near the conduction band edge position. Hence, Sn-doped WO₃ is expected to behave more n-type compared to undoped WO₃.

The Mott–Schottky (MS) plot was used to understand the band edge positions of WO₃ with respect to RHE values and water splitting potentials. Also, the carrier density could be determined using MS plots. The MS experiments were carried out in the dark in 0.5 M Na₂SO₄ at 1000 Hz of frequency using the equation

1/C² = [2/(N_Deε₀ε)] [(V_S−V_FB) – (k_BT/e)],

(1)

where C, N_D, e_, ε₀, ε, V_S, V_FB, k_B, and T are the space charge capacitance, carrier density, elementary charge, permittivity of the vacuum, the relative permittivity of WO₃, applied potential, flat band potential, Boltzmann constant, and temperature (299 K), respectively. Using the slope and intercept of the plot of 1/C² vs. applied potential (Figure 7c), N_D and V_FB were calculated. The V_FB and N_D values of undoped and Sn-doped WO₃ were obtained as 0.54 V (vs. RHE), 6.35 × 10¹⁷ cm⁻³ and 1.21 V (vs. RHE), 8.38 × 10¹⁷ cm⁻³, respectively.

The flat band potential values obtained using MS plots suggest the position of conduction band edge position of WO₃. Precisely, for the n-type semiconductors, the conduction band edge could be positioned at slightly lower potential with respect to its flat band potential [34]. Therefore, the conduction band edge of WO₃ is expected to be placed at less than 0.1 V of its flat band potentials [35]. Based on the flat band edge potentials, the conduction band edge of undoped and Sn-doped WO₃ was determined to be at 0.44 and 1.11 V vs. RHE, respectively. Subsequently, the positions of valence band edge position with respect to RHE were calculated using the band gap values (Figure 5) and conduction band edge positions. Figure 7d shows the band edge diagram of undoped and Sn-doped WO₃ plotted against RHE values and with reference to the overall water splitting potentials. Note that the positions of Fermi levels were also included, based on the XPS results. Overall, the band edge plot indicates that the doping of Sn pushes both the conduction and valence band edge of WO₃ towards higher potentials. At the same time, the Fermi level shifts upward towards the conduction band edge position, making it more n-type compared to the undoped WO₃. Such an observation was evident as the doping of post-transition metal ions into WO₃ shift its band edge positions towards higher potentials as previously observed in case of Bi dopant [13]. In general, it can be concluded that the doping of post-transition metal atoms into WO₃ causes a downward shift in band edge positions towards higher potentials and increases the carrier density in WO₃.

The enhancement in PEC properties of Sn-doped WO₃ was evaluated by measuring linear sweep voltammetry (LSV) of doped WO₃ photoanode by exposing the 1.5 G illumination in 0.5 M Na₂SO₄ electrolyte. Figure 8a depicts the LSV plots of undoped and Sn-doped WO₃ thin films. The substitutional doping of Sn in WO₃ was found to increase the photocurrent to 0.427 mA cm⁻² from 0.253 mA cm⁻² (at 1.23 V vs. RHE) of undoped WO₃. As shown in Figure S2, the photocurrent of WO₃ increases with the increase in atom % of doped Sn. The highest photocurrent was recorded at 0.80 atom % of Sn-doped WO₃. The IPCE measurements were recorded to understand the PEC behavior with respect to incident light wavelength using the below equation.

IPCE = (1240I)/(λP),

(2)

where I, λ, and P refers to the current density (mA cm⁻²) recorded at 1.23 V vs. RHE, the wavelength of incident light (nm), and power density at a given wavelength (mW cm⁻²), respectively. The IPCE plots shown in Figure 8b confirm that doping of Sn increases the IPCE values to 44% (at 300 nm) along with the slight redshift of onset absorbance. This suggests that Sn-doped WO₃ yields significantly higher currents than the undoped WO₃ at higher wavelengths. The calculated IPCE data were integrated between the wavelength range of 300 to 600 nm using standard values of AM 1.5 G data. The integrated plots, along with the IPCE data in Figure 8b, validate the photocurrent values obtained using the solar simulator. That is, the integrated photocurrents values were in a good agreement with the values recorded in the LSV plots. Figure 8c shows Nyquist plots of the undoped and Sn-doped WO₃ samples. The plots indicate that the 0.80 atom % Sn-doped WO₃ shows smaller semicircle and offers lower charge transfer resistance compared to the undoped WO₃ at the interface of electrode and electrolyte.

Furthermore, the stability and gas evolution amount of undoped and Sn-doped WO₃ thin films were measured, and the results are shown in Figure 9. Both photoanodes exhibited good stability under continuous illumination of AM 1.5 G sunlight at a constant bias potential of 1.23 V vs. RHE (Figure 9a). This indicates that both the undoped and Sn-doped WO₃ thin films show significant stability during water splitting. Moreover, the evolution of O₂ gas during the water splitting were measured using GC. As shown in the Figure 9b, the Sn-doped WO₃ show significant increase in the O₂ evolution rate compared to the undoped WO₃. Note that the valence band edge position of both undoped and Sn-doped WO₃ allow the spontaneous evolution O₂ gas only, and the H₂ gas was not detected due to its unsuitable conduction band edge position and low level of produced photocurrents. The faradaic efficiencies for the O₂ production was calculated using the following equation:

Faradaic efficiency = (4F × n_O2)/Q,

(3)

where, F, n_O2, and Q are defined as the Faraday constant, amount of O₂ (mol), and amount of charge (C), respectively. The faradaic efficiencies of undoped and Sn-doped WO₃ were calculated to be 85% and 89%, respectively.

The enhancement in PEC properties of WO₃ after Sn doping was ascribed to several factors, based on the characterization results. Specifically, the orientation of crystal facet-induced band gap reduction, increase in carrier density, and the reduced charge transfer resistance were the main influencing factors for PEC activity. That is, the change in crystal facet orientation from [200] to [002] direction significantly increases water oxidation reaction kinetics at the electrode interface. This was notably evident from the impedance measurements exhibiting reduced charge transfer resistance across the interface. The optical data suggest that Sn-doped WO₃ is capable of absorbing significantly more incident light compared to the undoped electrode, which also affects its photocatalytic properties. Notably, the increase in oxygen vacancies in WO₃ after Sn doping extends the light-harvesting ability to the visible and near-infrared region [11]. Consequently, the increase in oxygen vacancies tends to elevate the donor density levels in WO₃ as evident from the MS studies, and also improves the conductivity and enhances the PEC water splitting activities [3].

3. Materials and Methods

The thin films of Sn-doped WO₃ on FTO were synthesized using an earlier reported hydrothermal method, with minor modifications [27]. The tin(II) chloride dihydrate (Sigma-Aldrich ≥99.0%, St. Louis, MI, USA) precursor was employed as the source of Sn dopant in WO₃. The doping of Sn was achieved during the hydrothermal reaction of PTA solution at 150 °C (for 3 h) in the presence of Sn ions of desired concentrations. The amount of Sn dopant in WO₃ was varied by controlling the concentration of tin(II) chloride dihydrate in PTA solution as described in Table S1. During the hydrothermal reaction, the FTO substrate was placed vertically in the Teflon container with its conducting side facing downwards. The synthesized undoped and doped thin films were annealed at 500 °C for 2 h. For the detailed characterization, the Sn-doped WO₃ (0.80 atom %) thin film that showed the highest PEC activity was chosen.

A MiniFlex desktop XRD instrument from Rigaku (Tokyo, Japan) was used to analyze the crystalline properties of synthesized thin films between the 10° and 90° diffraction angle (2θ). The morphology and distribution of dopants in WO₃ were investigated using a Hitachi S4800 (Tokyo, Japan) scanning electron microscope (SEM) and a JEOL JEM-2100F (Peabody, MA, USA) transmission electron microscope (TEM). The optical properties were explored using a Varian Cary 5000 spectrophotometer (Sydney, Australia). A Thermo Fisher Scientific (Waltham, MA, USA) instrument was utilized to analyze the valency and valence band edge properties. Linear sweep, transient and impedance measurements were carried out using an IviumStat (Eindhoven, Netherlands) electrochemical interface. The ABET Technologies Sun 2000 solar simulator (ABET Technologies, USA) calibrated to AM 1.5G was used as a light source. For IPCE measurements, a xenon arc lamp (1000 W) connected to a Thermo Oriel 66902 monochromator (Thermo Oriel Cornerstone 7400 1/8, USA) was used. The O₂ gas evolved during the water splitting reaction was measured by using a YOUNG LIN YL 6500 gas chromatograph (GC) (Anyang, South Korea). Helium gas was used as a carrier medium in the capillary column, and a pulsed discharge ionization detector (PDD) was used to detect evolved O₂ gas. For the PEC measurements, a three-electrode system setup was selected in a quartz cell with fabricated WO₃ thin film as photoanode, Pt counter electrode, and Ag/AgCl reference electrode. All electrochemical measurements were carried out in 0.5 M Na₂SO₄ electrolyte. The reaction cell was thoroughly sealed, and the evolved O₂ gas was fed to GC using a pump with carrier gas. For the impedance measurements under illumination, a frequency range of 0.1−10⁵ Hz with AC voltage perturbation of 10 mV was used. The MS analysis was carried out in the dark at a frequency of 1000 Hz. The Nernst equation was used to convert the measured potential to RHE values.

4. Conclusions

Sn-doped WO₃ thin films were hydrothermally synthesized, and their PEC water splitting properties were explored. The uniform and substitutional doping of Sn in the WO₃ lattice was found to influence its morphology and crystal orientation. A slight reduction in the band gap was observed, along with the significant increase in the oxygen vacancies. The presence of Sn with a valency of (II) and (IV) in the WO₃ lattice was confirmed via XPS studies. The MS plots revealed an anodic shift in the flat band potential and increase in the carrier densities. The band edge diagram of WO₃ before and after Sn doping was plotted using XPS valence band edge, MS, and absorbance data. Based on this, a downward shift of both the valence and conduction band edge was revealed with a more n-type character of Sn-doped WO₃. The PEC studies carried out in 0.5 M Na₂SO₄ indicated that Sn-doped WO₃ performs significantly better as a photoanode than the undoped WO₃. The synergic effect of Sn doping yields increased photocurrent and incident photon-to-current efficiency values. The results of this work conclude that Sn-doped WO₃ performs as superior photoanode and could be utilized in performing efficient PEC water splitting. Furthermore, the scheme proposed in this article will benefit future research in fabricating more efficient WO₃ photoanodes for PEC water splitting reactions.