1. Introduction

PEC water splitting is considered an attractive means of harvesting solar energy and producing storable chemical fuels such as hydrogen. A photoelectrode is critical to the light harvesting and conversion efficiency of a PEC cell, wherein incident photons are absorbed, electron–hole pairs are generated, and subsequent charge transfer and catalytic conversion occurs. Particularly, photoelectrodes with nanostructured morphology contribute greatly towards the performance of PEC solar energy conversion [1,2]. TiO₂ has been widely studied as a photoanode for PEC water splitting [3] because its band-edge positions are favorable for redox reactions in water, it is chemically stable, resistant to photocorrosion, non-toxic, earth abundant and inexpensive. However, the photo-conversion efficiency of TiO₂ remains limited due to the following reasons:- Ti being a transition metal, has empty d orbitals, and its valence band being strongly influenced by oxygen 2p orbitals, results in a large bandgap of 3.0 eV (rutile) ~3.2 eV (anatase), which only allows TiO₂ to absorb in the UV region of the solar spectrum [4]. For photocatalysis, anatase TiO₂ is more reactive than the rutile phase, due to the former's longer electron lifetime; while anatase TiO₂ does have a short minority carrier lifetime, which would increase the electron-hole recombination [5]. We can overcome some of these limitations by generating defects to improve photoconversion efficiency and increasing the surface area to enhance charge carrier diffusion pathways. Introduction of disorder in the lattice would enhance visible and infrared absorption with the additional benefit of carrier trapping which localises photogenerated charge carriers to prevent fast recombination. Introducing lattice disorder in TiO₂ would give rise to mid-gap states such as Ti³⁺ and oxygen vacancies. Higher mid-gap states (~3.0 eV) are derived from Ti 3d orbitals only, while the lower energy (~1.8 eV) are hybridized from both O 2p orbitals and Ti 3d orbitals, and mainly from the valence band states. Lower energy mid-gap states lie below the Fermi level and can account for a large red shift of the valence band edge. Optical transition between mid-gap states and the conduction band tail would produce a charge transfer from the O 2p orbital to the Ti 3d orbital [6]. The localization of both photoexcited electrons and holes prevents fast recombination and facilitate photogenerated carriers to the active sites to perform redox reactions [7]. There are various methods of generating defects such as plasma-treatment, high energy bombardment, thermal treatment in vacuum and thermal treatment in hydrogen environment, with last two methods being the most commonly used. Hydrogen treatment by exposing TiO₂ in hydrogen atmosphere is a widely used technique to introduce defects. Modifications of this technique by varying the temperature, pressure condition, and introducing hydrogen during the growth process by rf-sputtering has recently been explored for nanocrystals, nanostructures and nanotubes [8–11]. However, this technique requires the use of large amount of hydrogen, which makes the purpose of using TiO₂ to generate hydrogen for energy related applications counterproductive. Electrochemical reduction is a simple and cost-effective method to introduce defects into TiO₂. By applying an electric field, electrons would be driven into the TiO₂ lattice and reduces Ti⁴⁺ to Ti³⁺. Recently, this method has been employed to generate defects in nanostructures such as nanotubes, nanograss and nanorods [12–14].

A nanostructured electrode possesses several favorable intrinsic characteristics for efficient light harvesting, such as access to a greater surface area, direct charge transfer pathways for reduced charge recombination, and strong light-trapping ability that promotes light harvesting by confining light within the photoelectrode. Therefore, the research community has placed enormous efforts towards the development of nanostructured semiconductor photoelectrodes for PEC solar energy conversion applications [15,16]. A photocatalyst with ordered macroporous structure can enhance molecular diffusion kinetics through its interconnected pores, as well as increase the surface area and amount of adsorption and reactive sites. Multiple reflections and scatterings inside the materials’ pores may effectively extend the optical path length of incident photons within the pore network and improve light capture [17,18]. Among the various approaches to make such macroporous nanostructured materials, the template-assisted method offers a simple and versatile route to produce nanostructured materials with predefined morphology and optoelectronic properties. Due to their highly ordered and tunable porous structures, self-assembled opals have become widely popular as candidates for nanostructure fabrication [19–21].

In this work, we report the generation of defects by electrochemical reduction (ER) of TiO₂ inverse opal (IO) photoelectrodes to increase photocatalytic activity and improve PEC water splitting performance. TiO₂-IO photoelectrodes were electrochemically reduced from 300 to 500 seconds, with the photoelectrode reduced for 400 seconds (ER400s) showing the highest current density compared to pristine TiO₂-IO. The defects generated have been shown to be in the form of oxygen vacancies, which shifts the Fermi level towards the conduction band edge, thus improving the conductivity of the TiO₂-IO electrode.

2. Experimental

TiO₂-IOs were fabricated by using the evaporation-assisted self-assembly via vertical deposition method with polystyrene (PS) particles as a sacrificial template on fluorine-doped SnO₂ (FTO) coated glass substrates. FTO substrates were first ultrasonically cleaned in ethanol for 1 hour before being dipped in a mixed suspension of Thermo Scientific carboxylate modified polystyrene particles of 530 nm in diameter and deionized water in the proportion of 0.15 wt%. FTO substrates dipped in the mixture were placed in an oven for approximately 22 hours, allowing lateral capillary forces to deposit PS particles on the substrate as water evaporated. The infiltration of sacrificial PS template with TiO₂ was carried by atomic layer deposition (ALD), using titanium tetrachloride and deionized water as sources for Ti and O respectively in a Picosun R200 ALD system at 85°C. Annealing the infiltrated opal templates at 450° C for 30 mins in air converted them from amorphous phase to crystalline phase and remove the PS sacrificial template, leaving behind anatase TiO₂-IO. Samco ICP-RIE (RIE-4100iP) was used to etch the top layer of TiO₂-IO to allow electrolyte to infiltrate deep into the porous structure.

Electrochemical reduction of TiO₂ inverse opals was carried by an anodic reduction process using a two-electrode configuration and a potentiostat (Autolab, PGSTAT302N) setup, with TiO₂-IO as the working electrode and Pt wire as the counter electrode. With the two photoelectrodes placed approximately 2.5 cm apart, electrochemical reduction was carried at a voltage of − 3V from 300 to 600 seconds in an electrolyte bath of 0.5 M Na₂SO₄. After the anodic reduction, the samples were rinsed in deionized water and dried with nitrogen.

The surface morphology and structure were characterized by scanning electron microscopy (SEM, FEI Helios Nanolab 600) operating at a voltage of 3 kV. Phase identification of TiO₂ was carried out by grazing incidence X-ray diffraction (GIXRD, PANanalytical), using Cu-Kα radiation (λ = 1.54 angstrom) with a voltage of 45 kV from 20° to 80°. Raman spectroscopy (Renishaw) using a 532 nm laser with Raman shift from 100 to 800 cm⁻¹, was carried out to confirm the defects generated from the electrochemical reduction process. X-ray photoelectron spectroscopy (XPS) using monochromated Al K alpha x-ray source (energy 1486.68 eV) and ultra-violet photoelectron spectroscopy (UPS) using with a He(I) beam (21.1 eV) was performed with Thermo Scientific ESCALAD 250Xi. Diffuse reflectance and transmittance were measured using a Perkin Elmer UV-Vis spectrophotometer equipped with an integrating sphere.

PEC performance, electrochemical impedance spectroscopy (EIS) and Mott-Schottky analyses of the photoanodes were evaluated in a three-electrode PEC cell configuration using a potentiostat (Autolab, PGSTAT302N). All potentials were reported against the reversible hydrogen electrode (RHE) using Nernst’s equation E (V_RHE) = E (V_Ag/AgCl) + 0.059 pH + 0.197. Ag/AgCl in saturated KCl and a Pt wire were used as reference and counter electrodes, respectively in 1 M NaOH (pH = 13) as the electrolyte. Photoresponse characteristics were measured using a solar simulator (Newport) equipped with a 150 W Xe lamp and an AM 1.5G filter, calibrated with a standard Si solar cell. Transient photocurrent measurements were performed under chopped light irradiation (on-off light cycles of 50 s) at an applied bias of 1.23 V_RHE in 1 M NaOH. Photocurrent stability measurement was performed at an applied bias of 1.23 V_RHE in 1 M NaOH for a duration of 3.5 hours. Open circuit voltages for all samples were measured before EIS was performed from 10⁻² to 10⁵ Hz at 10 mV amplitude. Mott-Schottky analyses were performed at a frequency of 500 Hz and an applied voltage of 10 mV. ZView from Scribner Associates was used to analyse EIS data. The incident-photon-to-current conversion efficiency (IPCE) measurements were taken in a three-electrode configuration as a function of wavelength from 300 to 600 nm using a 1000 W Xe lamp coupled to a monochromator. The incident light intensity was calibrated using a Thorlabs PM100D power meter and a standard silicon photodiode power sensor of wavelength range 200 – 1100 nm. IPCE was calculated according to the following equation:

3. Results and discussion

Figure 1(a) shows the highly ordered PS sacrificial template consisting of a face-centered cubic array structure with PS spheres stacked in an orderly manner. By using a stop-flow ALD method, we can effectively infiltrate the PS sacrificial template to achieve uniform and high filling fraction of the macroporous structure with an average pore size of approximately 500 ± 10 nm as seen in Figs. 1(b) and 1(c). The magnified cross-section of TiO₂-IO in Fig. 1(d) shows a well-connected macroporous structure with unclogged and well-defined pore openings. This would allow infiltration of the electrolyte deep into the structure ensuring shorter pathways for photogenerated carriers. Photogenerated holes can easily reach reactive sites to perform the oxidation process, while photogenerated electrons can travel quickly to the FTO ohmic contact, through the cell and to the counter electrode to perform the reduction process. Recombination of electron-hole pairs is reduced since photogenerated carriers can access their respective sites quickly. After undergoing electrochemical reduction and stability test for 3.5 hours, the TiO₂-IO maintains its structure and the pores remained unclogged, as shown in Figs. 1(e) and 1(f).

 

Fig. 1 SEM images of (a) PS opal template (b) top view of TiO₂-IO (c) cross section of TiO₂-IO (d) magnified cross sectional image TiO₂-IO showing the pore structure (e) ER400s-TiO₂-IO and (f) ER400s TiO₂-IO after 3.5 hrs of photoelectrochemical stability test.

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The crystalline structure of both pristine and ER TiO₂-IO was characterized by GIXRD as shown in Fig. 2(a). The major XRD peaks at 25.23°, 36.84°, 48°, 53.65°, 62.53°, 70.37°, and 75° with d_hkl 3.51, 2.44, 1.89, 1.70, 1.48, 1.34 and 1.26 $\dot{A}$ can be indexed to the tetragonal anatase phase of TiO₂ (PDF card no. 00-044-0477). These peaks correspond to the diffraction planes (101), (103), (200), (105), (204), (220) and (215) respectively. There is no change in the phase after the electrochemical reduction process and after 3.5 hours of photoelectrochemical stability test. The Raman spectrum as seen in Fig. 2(b) for pristine TiO₂-IO shows six active modes with three E_g modes at 144.1 cm⁻¹, 196.8 cm⁻¹ and 637.2 cm⁻¹; two B_1g modes at 395 cm⁻¹ and 516.5 cm⁻¹; and one A_1g mode at 516.5 cm⁻¹. These modes are typical of anatase TiO₂ [22]. The active mode E_g is caused by symmetric stretching vibration, B_1g by the symmetric bending and A_1g by the antisymmetric bending of O-Ti-O in TiO₂ [23]. Both the E_g(1) and E_g(2) modes at 144.1 cm⁻¹ and 196.8 cm⁻¹ from ER400s TiO₂-IO are blue-shifted, while the rest of the Raman modes remain unchanged. These two peak shifts have been attributed to lattice disorder resulting from a phonon confinement effect or a stoichiometric change [24]. From the XRD pattern and using Scherrer’s equation, the crystallite size of both pristine and reduced TiO₂ was estimated to be 22 nm. Phonon confinement effect can be ruled out since there was no change in grain size of the pristine and ER400s TiO₂-IO; and only the E_g modes were shifted. Removing an oxygen atom from the TiO₂ lattice forces the three Ti atoms to shorten their overlapping dangling bonds with each other and strengthen their bonds with the rest of the lattice, thus decreasing the Ti-O bond length [25]. This reduction in the bond length is reflected in the shift of E_g modes to a higher wave number, since the E_g modes are mainly caused by the symmetric stretching of the Ti-O bond and is sensitive to change in bond length. All the peaks of the ER400s broadened, as compared to the pristine sample, with full width half maximum (FWHM), shown in Table 1. Peak broadening can be attributed to the excess oxygen vacancies leading to non-stoichiometry of TiO₂ [26]. Thus, the peak shift, broadening of the peaks and decrease in intensity of surface structure of TiO₂ are due to the defects, in this case oxygen vacancies generated from the reduction process [27].

 

Fig. 2 (a) XRD patterns of FTO substrate, pristine, ER400s and ER400s (after stability test) TiO₂-IO (b) Raman spectra of pristine and ER400s TiO₂-IOs.

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Table 1. FWHM of Raman spectra for pristine and ER400s samples

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XPS analysis was carried out to determine the chemical environment and type of surface defects induced by the reduction process. For the XPS core-level spectra of Ti 2p in Fig. 3(a), the peaks at binding energy (BE) 459 eV for pristine and 458.5 eV for ER TiO₂-IO have been assigned to Ti⁴⁺ species. No peaks for Ti³⁺ species have been observed hence suggesting there was no reduction of Ti⁴⁺ to Ti³⁺ [28,29]. The spectra of O 1s in Fig. 3(b) has one main peak at BE 530.2 eV (pristine) and 529.8 eV (ER TiO₂-IO) characteristic of Ti-O. The two smaller shoulder peaks are assigned to the surface oxygen vacancies, with BE 531.6 eV (pristine) and 531.1 eV (ER) TiO₂-IO; and the surface OH group with BE 532.6 eV (pristine) and 532.2 eV (ER) TiO₂-IO as shown in Fig. 3(c). With the reduction process, there is a decrease in atomic percentages in Ti⁴⁺ and Ti-O peaks while an increase in the peaks is attributed to oxygen vacancies and the OH group, shown in the inset of Fig. 3(c). A shift to lower binding energies of both O 1s and Ti 2p spectra of ER-TiO₂ is observed across all the peaks. A lack of Ti³⁺ in ER TiO₂-IO with a shift to lower BE in the spectra indicates that electrons bound to the oxygen and titanium ions migrated to the oxygen vacancies; with the latter serving as electron traps. The shift in O 1s spectra can be attributed to the transfer of electrons to the neighbouring oxygen vacancies [30–32].

 

Fig. 3 (a) XPS core level spectra of Ti 2p (b) XPS core level spectra of O 1s (c) Fitting of XPS core level spectra (d) UPS spectra of pristine and ER400s TiO₂-IOs.

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UPS measurements were performed to determine the valence band maximum and the work function of TiO₂-IOs before and after the reduction process. The reduction of TiO₂-IOs did not change the valence band maximum which is 3.57 eV below the Fermi level as indicated in Fig. 3(d). The work function of a semiconductor surface is given by ϕ_s = hυ - E_sc where hυ is the energy of photons emitted by helium gas of 21.2 eV and E_sc is the secondary electron cut-off. The work function for both samples is 5.12 eV, which agrees with reported values for polycrystalline anatase TiO₂ [33]. However, there is an increase in the peak intensity due to the increase in oxygen vacancies from the reduction process.

Figure 4 shows the linear sweep voltammetry (LSV) of pristine and ER TiO₂-IO photoelectrodes illuminated through the front and back side. Pristine TiO₂-IO photoelectrode has a current density of 0.09 mA/cm² (front side illumination) and 0.18 mA/cm² (back side illumination) with an applied potential of 1.23 V_RHE, while electrochemically reduced for 400s (ER400s) TiO₂-IO photoelectrode has the highest current density of 0.35 mA/cm² (front side illumination) with an applied potential of 1.23 V_RHE. When pristine TiO₂-IO photoelectrode was illuminated from the front, the Fermi level downshift towards the conduction band and surface state assisted recombination decreases leading to an initial photoanodic current increase. However, as the surface states are filled with photogenerated holes, the photoanodic current starts to plateau which corresponds to the maximum achievable current from the surface states. The photocurrent saturates once the surface states assisted recombination have been eliminated by electrons. When ER400s TiO₂-IO photoelectrode was illuminated from front and back side, we see only an increasing photoanodic current until saturation is reached.

 

Fig. 4 Linear sweep voltammograms obtained for pristine and electrochemically reduced TiO₂-IO photoelectrodes for duration of 300 s, 400 s and 500 s in 1 M NaOH (pH = 13) electrolyte, in darkness and under AM1.5 G simulated sunlight (100 mW cm⁻²), and at a scan rate of 10 mV s⁻¹.

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For an n-type semiconductor such as TiO₂, photogenerated holes migrate to the semiconductor-electrolyte interface (SEI), while photogenerated electrons travel to the FTO ohmic contact. Absorption of photons and generation of carriers occur at the surface where they are illuminated. Direction and distance travelled by photogenerated carriers are the same regardless of direction of illumination. However, the photocurrent is dependent on the direction of illumination and can indicate the limitations of electron or hole transport. The fact that photocurrent generated is significantly different for pristine and ER400s TiO₂-IO photoelectrodes confirms that the charge recombination is significant in pristine TiO₂-IO. This is addressed by introducing oxygen vacancies in TiO₂-IO. Oxygen vacancies not only act as electron donors thus lowering the Fermi level towards the conduction band, they also serve as electron traps to prevent charge recombination. Since the depletion layer width is dependent on the donor density, oxygen vacancies present in the TiO₂ lattice can affect the depletion layer width [34]. With increased oxygen vacancies generated by the electrochemical reduction process, the depletion layer width is shortened, recombination is reduced, thus improving charge carrier migration to their respective sites. However, increasing the reduction time from 400 s to 500 s leads to a decrease in photocurrent density, which can be attributed to the excessive generation of oxygen vacancies. With excess oxygen vacancies, more electrons are trapped thus lowering density of electrons collected at the FTO ohmic contact.

Figure 5(a) shows the transient photocurrent responses for pristine and ER400s TiO₂-IO photoelectrodes at an applied potential of 1.23 V_RHE with 50 s on-off light cycles. The photocurrents for both pristine and ER400s-IO photoelectrodes reflect the measured values from LSV, and is maintained for over a period of 600 s. The transient features indicate whether recombination is limiting the charge transfer process and the type of recombination that occurs at the SEI [35]. An anodic transient would correspond to holes accumulated, while a cathodic transient corresponds to electrons being dissipated at the SEI. For the ER400s TiO₂-IO photoelectrode, an anodic photocurrent spike is observed due to the sudden generation of charge carriers, which then decays to a steady-state photocurrent due to recombination processes or a decrease in band bending [36]. For pristine TiO₂-IO photoelectrode, photogenerated holes travel to the surface at a slow but sufficient rate to react with the electrolyte, thereby avoiding accumulation. By creating oxygen vacancies, electrons are trapped in the mid-gap states as well as absorption sites where hole transfer to the absorbed ionic species can prevent further recombination [37]. Long-term stability of ER400s TiO₂-IO photoelectrode was subjected to continuous AM1.5 simulated sunlight illumination for 3.5 hours. Under an applied potential of 1.23 V_RHE, the initial photocurrent density of 0.35 mA/cm² decreased by 14% to 0.29 mA/cm², as shown in Fig. 5(b). To quantify light absorption and photoconversion efficiencies of the TiO₂-IO photoelectrodes, we measure the incident photon-to-current efficiency (IPCE) where the photocurrents are obtained as a function of wavelength. Figure 5(c) shows the IPCE increasing almost 3 times in the absorption region of TiO₂ after electrochemical reduction, as well as increase slightly in the visible region. The increase in photoconversion efficiency of TiO₂-IO in the UV as well as in the visible regions can be attributed to the enhanced photocurrent density in ER400s TiO₂-IO photoelectrodes. The UV-Vis spectra in Fig. 5(d) shows the absorption band edges for both TiO₂-IO photoelectrodes at the wavelength of 400 nm, which is at the absorption bandgap of anatase TiO₂. The ER400s TiO₂-IO photoelectrode exhibits a strong sub-bandgap absorption from 400 nm to the infrared region due to an extended Urbach tail, which has been attributed to lattice disorder or vacancies [38].

 

Fig. 5 (a) Transient photocurrent responses of pristine and ER400s TiO₂-IO photoelectrodes measured under 50 s on-off light cycles (b) Photocurrent-time measurement of ER400s TiO₂-IO photoelectrode measured for 3.5 hours (c) IPCE of pristine and ER400s TiO₂-IO photoelectrodes measured at incident wavelength from 300 to 500 nm and (d) absorption spectra of pristine and ER400s TiO₂-IO photoelectrodes as a function of wavelength. Transient photocurrent, photocurrent-time and IPCE measurements were carried out in 1 M NaOH at an applied potential of 1.23 V_RHE.

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EIS was used to evaluate charge transfer kinetics at the SEI by measuring the AC response of pristine and ER400s TiO₂-IO photoelectrodes. Nyquist and Bode plots of pristine and ER400s TiO₂-IO photoelectrodes obtained from EIS was carried out at open-circuit potential, under dark and illuminated conditions. The internal resistances corresponding to the overall charge transfer process under illumination were significantly lower than those in dark conditions due to the generation of electron-hole pairs. In Nyquist plots, the real and negative imaginary components of the complex impedance Z are plotted on the horizontal and vertical axis, respectively. The presence of one semicircle indicates that an electrochemical interface is present in the system. In Fig. 6(a), all TiO₂-IO photoelectrodes under illumination have lower impedances than those performed in the dark indicating that photogenerated electron-hole pairs lower the resistance of the photoelectrodes. The ER400s TiO₂-IO photoelectrode has a lower impedance compared to the pristine electrode which shows that oxygen vacancies improved charge transfer kinetics at the SEI. A Randles circuit as shown in the inset of Fig. 6(b), consists of a series resistance (R_s) element followed by a parallel RC, which is a combination of parallel resistance (R_p) and a constant phase element (CPE), and is used for modeling the response of the photoelectrodes. The series resistance represents the ohmic resistance between FTO and TiO₂-IO, while the parallel resistance is associated with charge transfer from TiO₂-IO to the electrolyte, and the CPE is the impedance originating from the charge in the space charge layer in TiO₂-IO. The CPE is a fitting element with impedance of Z_CPE = 1 / Q(jω)^α, and when α equals 1, 0, or − 1 a CPE behaves as a capacitor, a resistor, or an inductor. By fitting the Nyquist plots using the equivalent Randles circuit, we can ascertain the appropriate behavior of system to AC response. Through the creation of oxygen vacancies, Ti-O bonds are broken, producing excess valence electrons that are bonded to oxygen atoms in the valence band, and in migrating from a lower to higher energy level, conductivity of the electrode is improved [39]. From the Bode magnitude plots shown in Fig. 6(c), the impedance measured corresponds to R_s and R_p at low frequencies, while at higher frequencies, we see a drop in the overall impedance indicating that the charge transfer resistance at the interface has dropped since the resistance between the ohmic contact and TiO₂-IO did not change. From the frequency peaks in the Bode phase plots in Fig. 6(d), lifetime of photogenerated electrons can be calculated using, τ_n = 1/ (2πƒ_max), where ƒ_max is the maximum oscillation frequency of the impedance. Photogenerated electron lifetimes for pristine and ER400s TiO₂-IO photoelectrodes were calculated to be 75 and 234 ms respectively. By generating oxygen vacancies, electron lifetimes have been extended leading to the increase of electron diffusion length and inhibition of electron-hole recombination.

 

Fig. 6 (a) Nyquist plot of pristine, ER300s, ER400s and ER500s TiO₂-IO photoelectrodes (b) Nyquist plots of pristine and ER400s TiO₂-IO photoelectrodes fitted to the equivalent circuit. (c) Bode magnitude and (d) Bode phase plots of pristine, ER300s, ER400s and ER500s TiO₂-IO photoelectrodes.

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Mott-Schottky analysis for both pristine and ER400s TiO₂-IO photoelectrodes was carried to determine the flatband potential and charge carrier density, as shown in Figs. 7(a) and 7(b). The Mott-Schottky equation can be written as:

 

Fig. 7 (a) Mott-Schottky plots collected for (a) pristine (b) ER400s TiO₂-IO photoelectrodes at 500Hz in dark condition.

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4. Conclusions

Defects in TiO₂ can be generated easily and quickly by electrochemical reduction process. Electrochemically reduced TiO₂ macroporous nanostructures such as TiO₂-IO photoelectrodes show an improvement in photon absorption and PEC water splitting performance. The defects generated in these TiO₂-IO photoelectrodes are oxygen vacancies. The performance of the electrochemically reduced photoelectrodes is dependent on the reduction times, with 400 s being the optimal reduction time. Longer reduction time produces excessive oxygen vacancies that act as electron traps, thereby limiting the density of electrons collected at the FTO ohmic contact. There is an increase in the photon absorption of the ER400s TiO₂-IO photoelectrode in the visible region and photoconversion efficiency in the UV region increasing by almost 3 folds. This improved performance ostensibly results from the creation of oxygen vacancies in the TiO₂-IO photoelectrodes due to shift in Fermi level towards the conduction band, which in turn leads to lower electron-hole recombination rates and improved charge transport. However, in the experiment, the valence band did not change and the reduction of the bandgap of TiO₂-IO was not achieved. This work has shown that by introducing oxygen vacancies, we can increase photogenerated carrier lifetime and thus improve the conductivity of TiO₂ to enhance PEC water splitting performance.