The effect of thermal annealing on the interfacial properties and photoelectrochemical performance of Ti doped Fe₂O₃ nanowire arrays

Abstract

Hematite (α-Fe₂O₃) is considered a promising material for solar water splitting but still faces the challenges of poor conductivity and slow charge transfer kinetics. This issue is normally addressed by doping along with postsynthesis thermal treatments. In order to understand how the postsynthesis thermal treatment influences doped Fe₂O₃, Ti doped Fe₂O₃ nanorod arrays were synthesized and subjected to different post annealing processes. The TiCl₄ solution hydrolysis treatment for doping and the post annealing time were optimized. To understand the thermal annealing effect on the interfacial properties and photoelectrochemical performance, the structural and photoelectrochemical properties of Ti doped and undoped samples with different annealing conditions were investigated. The formation energies for Ti or Sn doping were estimated by theoretical calculation. Both our theoretical and experimental results showed that the dopant Ti was easier to incorporated into Fe₂O₃ than Sn and thus required less annealing time. Our results also demonstrated the significant influence of postsynthesis thermal treatment on the photoelectrochemical performance of doped Fe₂O₃ nanorod arrays.

---

Introduction

Solar energy is the richest renewable and environmentally safe alternative energy source on Earth, and hydrogen is one of the prime candidates to be a carrier for the storage and use of solar energy. Photoelectrochemical (PEC) water splitting for hydrogen production has been considered as a promising approach to solar energy conversion.¹ Since 1972,² a lot of semiconductor materials have been introduced into the PEC process, such as TiO₂,³ ZnO⁴ and WO₃.^5,6 Hematite (α-Fe₂O₃), a promising material due to its appropriate bandgap, non-toxicity and abundance, has drawn much attention in recent years.⁷ However, the application of hematite for solar water splitting is hindered by challenges of poor conductivity, high recombination rate, poor oxygen evolution reaction (OER) kinetics, short hole diffusion length (2–4 nm), and improper band position for unassisted water splitting.⁸

Ti doping has been considered as a feasible way to improve the performance of Hematite in PEC process.^9–11 Gongming Wang et al.¹² reported a facile method for preparing Ti-doped α-Fe₂O₃ films by a new deposition-annealing (DA) process. The method of APCVD was used to synthesize concentration controllable and photoelectrochemically active Ti doped hematite thin films.¹³ As reported by Omid Zandi et al.,¹⁴ a relatively large percentage of Ti doped hematite gave rise to a combination of improved bulk properties (hole collection length) and surface properties (resurrection of the dead surface to achieve a higher water oxidation efficiency) and resulted in an enhancement of PEC water oxidation performance. It has been demonstrated that Ti doping not only increases the free carrier concentration but also increases the motilities of electrons and holes.^14,15 Higher carriers' diffusion lengths were also observed in the doped samples than the intrinsic samples.¹⁶ Furthermore, Ti doping was also found capable of surface states passivation and resulting in the decrease of the electron–hole recombination rate.¹⁷ Cathodic shift of the onset potential was observed in Ti doped Fe₂O₃ and ascribed to suppressing of the back reaction.¹⁸ Improvement of charge-transfer rate coefficient at the surface was also observed by Glasscock et al.¹⁹ Zhong et al.²⁰ reported that oxygen vacancies in hematite mainly affected the performance by improving the donor density, while Ti-doping might affect the performance of hematite by surface catalytic effects or more active sites for water oxidation. The insight into the mechanisms of these properties change was investigated by state-of-art techniques. For example, soft X-ray absorption spectra at the oxygen K-edge show that Ti incorporation creates new oxygen 2p-hybridized states, which overlaps with and distorts the known unoccupied Fe 3d–O 2p band of α-Fe₂O₃.¹⁰

Although the Ti doped Fe₂O₃ have been thoroughly investigated, however, the effects of Ti incorporation on its electronic and photoelectrochemical properties appear to be highly synthesis-dependent.¹⁰ To increase the Ti doping concentration, in situ solid-state reaction method was adopted to improve Ti doping concentration from 2.2% to 19.7%.²¹ Co-doping of Ti with other elements such as Si,²² Sn^23,24 were also found increase their PEC performance.

In addition to doping, introduction of a Fe₂O₃:Ti/ZnFe₂O₄ (ref. 25) or Fe₂TiO₅/Fe₂O₃ (ref. 26) heterojunction was used to further improve the electron–hole separation. TiO₂ over-layer was deposited on Fe₂O₃ nanosystem by ALD, which enhanced charge separation as well as the PEC performance.²⁷ Incorporation of g-C₃N₄ nanoparticles into hematite significantly improved the performance due to a synergy effect at the interface of the g-C₃N₄ and Ti–Fe₂O₃, including favorable contact, matched energy band levels and enhanced charge separation.²⁸ 6-fold enhancement of photocurrent density and negative shift of photocurrent onset by 500 mV were also observed when combining Ti-doped Fe₂O₃ and CdS nanoparticles.²⁹ Furthermore, surface modification³⁰ and corrosion³¹ were studied to shift the flat-band potential and the photocurrent onset potential of Ti-doped α-Fe₂O₃ films, respectively. Surface treatment with Al³⁺ was found passivate the surface states and decrease the accumulation of charge at the surface of Ti doped Fe₂O₃.³² Surface treatment with Zn however decreased the over-potential of the water oxidation reaction of hematite due to its interfacial catalytic activity for water oxidation.^33,34 Acid-treated hematite with decreased surface electron–hole recombination losses has also been reported.³⁵ With these results in mind, one can expect that postsynthesis thermal treatments would exert a great influence on its performance, by promoting crystallinity or reducing the defect density in solution-processed nanomaterials.^36,37 Here we investigated the post thermal treatment effect on Ti doped Fe₂O₃ nanostructures. We synthesized Fe₂O₃:Ti nanorod arrays for PEC water splitting by a TiCl₄ solution hydrolysis method. Ti element was incorporated into Fe₂O₃ nanowires by post annealing. And we also optimized the TiCl₄ solution hydrolysis treatment to obtain the Fe₂O₃:Ti structured nanorod arrays photoanode with the best PEC characteristic. As observed by Caroline Toussaint et al.,³⁸ Ti doped hematite films with its mesostructure preserved gave a better PEC performance than comparable dense and collapsed films. In order to keep the morphology of the nanorod arrays and conductivity of the FTO glass into Fe₂O₃ nanorods, we used a fast annealing process at high temperature for the post annealing treatment. The annealing condition was adjusted and the structural and photoelectrochemical properties of corresponding samples were investigated to understand the thermal annealing effect on the interfacial property and the photoelectrochemical performance.

Experimental

Sythesis of Fe₂O₃:Ti thin films

The hematite nanowire arrays were synthesized on a FTO coated glass substrate by the hydrothermal method.³⁹ Typically, 3.65 g (0.15 M) of ferric chloride (FeCl₃·6H₂O) and 7.65 g (1 M) of sodium nitrate (NaNO₃) were dissolved in 90 mL of deionized water and added with 100 μL of HCl (12 M). This precursor was then transferred to an autoclave, sealed and kept in an oven at 100 °C for 24 h. The resultant thin film was rinsed with deionized water and dried in a nitrogen stream, and then it was soaked in 0.2 M TiCl₄ solution with water bath at room temperature for a certain time (e.g. 24 h) to obtain the TiO₂ shell structure. Both pure hematite films (without TiCl₄ solution treatment, marked with “P-XX”) and Fe₂O₃:Ti nanowire films (with TiCl₄ solution treatment, marked with “T-XX”) were firstly annealed at 550 °C for 2 h with the rate of 4.5 °C min⁻¹ and then further annealed using a rapid thermal annealing system (RTA) in a MTI OTF-1200X tube furnace at 800 °C for different time with the rate of 13 °C s⁻¹.

Characterization

The phases of the samples were determined by X-ray diffraction (XRD) using an PANalytical X'pert MPD Pro X-ray diffractometer with Ni-filtered Cu-Kα irradiation (λ = 1.5406 Å). The morphologies of the thin films were observed using a JEOL JSM-7800FE scanning electron microscope (SEM). The transmission electron microscope (FEI Tecnai G2 F30 S-Twin, 300 kV) was used to determine the crystal structure of the samples and obtain the TEM-EDS mapping and line-scan data. The surface compositions of the samples were analyzed by X-ray photo-electron spectroscopy (Axis UltraDLD, Kratos, 150 W) with monochromatic an Al Kα irradiation. The controlled Ar⁺ etching was conducted at voltage of 4000 V and current of 4 μA cm⁻². A three-electrode cell with a 0.5 M Na₂SO₄ aqueous solution (pH = 6.5) as electrolyte, a Pt plate (1 cm × 2 cm) as the counter electrode and an Ag/AgCl electrode (saturated KCl solution) as the reference electrode were used for PEC measurements. A 350 W Xe lamp solar simulator (100 mW cm⁻²) with AM 1.5 G filter was used as the light source. The PEC results were recorded with a CHI 760D electrochemical workstation. Mott–Schottky curves were measured at a frequency of 1 kHz in the dark condition. Electrochemical impedance spectra (EIS) were measured at 0.5 V (vs. Ag/AgCl) with frequency range of 1 Hz to 100 kHz.

Computational details

A 2 × 2 × 1 Fe₂O₃ supercell with an iron atom replaced by a titanium (Ti) or tin (Sn) atom was used to mimic the Ti or Sn doping conducted in our experiment, with a dopant concentration of ∼2.1%. To estimate the doping capability of these foreign elements, the formation energy was calculated with the following formula,

E_f = E(X: Fe₂O₃) − E(Fe₂O₃) + μ_Fe − μ_X. (1)

In this formula E(X: Fe₂O₃) and E(Fe₂O₃) are the electronic energies of the Fe₂O₃ supercell with and without a dopant in the neutral state. μ_Fe and μ_X are the chemical potentials of Fe and dopant elements, of which the method to determine the chemical potential was given in detail in our previous paper.⁴⁰ The key point is to prevent dopant Ti and Sn in Fe₂O₃ forming the secondary phases of rutile TiO₂ and SnO₂.

The electronic energy of pristine and doped Fe₂O₃ and the chemical potential of O, Fe and dopants were calculated with the ab initio code VASP. The generalized gradient approximation in the Perdew–Burke–Ernzerhof formalism was used to treat the exchange–correlation interaction within the framework of density functional theory. The Coulomb correction (U = 5 eV, J = 1 eV) among Fe 3d electrons was introduced to describe the strongly correlated electronic nature of Fe₂O₃. The projector augmented wave potential was employed to treat the electron–ion interaction. The energy cutoff of 520 eV was adopted to expand the electronic wave function, and the k-point mesh of 3 × 3 × 3 for the Brillouin zone sampling. The valence electron configurations were O 2s²2p⁴, Fe 3d⁶4s², Ti 2p⁶3d³4s¹ and Sn 4d¹⁰5s²5p².

Results and discussion

Optimization of Fe₂O₃:Ti thin film synthesis

The as-prepared FeOOH nanowire arrays were soaked with TiCl₄ solution and followed with rapid annealing treatment. Two different TiCl₄ solution concentrations (0.2 M and 0.4 M) were used and the result showed that the sample treated with 0.2 M TiCl₄ solution for 12 h generated a higher photocurrent than that treated with 0.4 M TiCl₄ for 12 h as shown in Fig. S1.† Moreover, the sample exfoliated after being treated with 0.4 M TiCl₄ solution for 24 h. This could be attributed to the erosion by the high acidity of the 0.4 M TiCl₄ solution with an extended treatment time. It was also found that the amount of TiO₂ deposited on FeOOH nanowire was less with 0.4 M TiCl₄ solution due to the high acidity inhibiting the hydrolysis reaction of TiCl₄. The treatment time in TiCl₄ solution was then optimized in 0.2 M TiCl₄ solution.

Fig. S2† shows the photocurrents of Fe₂O₃:Ti structured thin films with different treatment time in TiCl₄ solution. The photocurrent density increased with treatment time and reached the maximum value (∼1.1 mA cm⁻² at 1.0 V vs. Ag/AgCl, ∼1.5 mA cm⁻² at 1.23 V vs. Ag/AgCl) when treatment time was 24 h. However, when the treatment time was increased to 48 h, the photocurrent density decreased significantly (∼0.7 mA cm⁻² at 1.0 V vs. Ag/AgCl, ∼1.0 mA cm⁻² at 1.23 V vs. Ag/AgCl). So a long treatment time is detrimental to the photocurrent and the optimal treatment time is determined to be 24 h.

High temperature treatment is usually used to improve the photoelectrochemical properties of hematite samples. It was reported that heating treatment at ∼800 °C significantly increased the conductivity of Fe₂O₃ by the distortion of C3v crystal⁴¹ and diffusion of Sn from FTO into Fe₂O₃ rods.^42,43 However the conductivity of FTO substrate usually decreased¹² and thermal deformation of substrate occurred⁴⁴ at such high temperature for relatively long annealing time.

To elucidate the effect of high temperature annealing time on the performance of Fe₂O₃:Ti samples, the as-prepared FeOOH nanowire arrays were treated with 0.2 M TiCl₄ solution for 24 h and then annealed at 800 °C for different time using a rapid thermal annealing system (RTA). As shown in Fig. S3,† annealing time affected the photocurrent performance of the Fe₂O₃:Ti samples annealed with different time. The photocurrent density increased with the annealing time when the time was no more than 120 s, but started decreasing with the time increasing from 3 min to 10 min. At the beginning of annealing at 800 °C, the crystallinity of Fe₂O₃ rods increased and showed improved photocurrent. However with a prolonged annealing time (longer than 3 min), the photocurrent decreased due to the decreasing of FTO conductivity during the annealing. The conductivity of the FTO substrate decreased significantly after long time high temperature annealing. Impedance measurement of Fe₂O₃:Ti structured thin films with different annealing time at 800 °C were also carried out and shown in Fig. S4.† The experimental data (dotted points) were fitted (solid curves) using an equivalent circuit and shown in the inset of Fig. S4.† In this equivalent circuit model, R_s represents the series resistance combining the electrolyte, FTO and external contact for the electrochemical device. It can be found that R_s increased gradually with increase of annealing time. As there should be no change of resistance of electrolyte and external contact, the increase of R_s should originate from resistance increase of FTO substrate. The optimum annealing time for the Fe₂O₃:Ti samples was then determined to be 90 s. The morphological and crystal structures of Fe₂O₃ and Fe₂O₃:Ti films were further investigated by using TEM. By comparing Fig. 1a and b, the Fe₂O₃ rods were covered by a uniform shell of TiO₂ with a thickness of about 2 nm after TiCl₄ treating. An interplanar spacing of 2.50 Å could be observed, which corresponds to the (110) planes of the rhombohedral Fe₂O₃. The HRTEM (inset of Fig. 1b) image of the interface area (marked by the box) in Fig. 4b reveals an interplanar lattice fringe of 2.12 Å, which could be attributed to the (121) plane of TiO₂.

Fig. 1 TEM micrographs of pure Fe₂O₃ films (a) and Fe₂O₃:Ti films (b) annealed at 800 °C for 90 s. (c) TEM EDS elemental mapping data and (d) linear scanning results of Fe₂O₃:Ti films annealed at 800 °C for 90 s.

The TEM EDX elemental mapping data (Fig. 1c) indicated that the spatially distribution of Ti element was a little broader than that of Fe element, especially at the tip area. Fig. 1d displays the TEM EDX linear scanning results which show that the concentration of Ti element at cave area and the edge of the rod was much higher than that in other areas. The enrichment of Ti element in the holes of the rods could be a result of chemical bath method used for TiO₂ deposition.

Thermal annealing effects on the performance of Fe₂O₃:Ti nanowire arrays

To gain insight into the mechanism of trade-off between crystallinity improvement of Fe₂O₃ rods, doping effect and FTO conductivity decreasing, detailed investigation of interfacial property and photoelectrochemical performance of pristine and TiCl₄ treated Fe₂O₃ nanowires annealed with different time were conducted and discussed.

Fig. 2 shows the SEM images of pure Fe₂O₃ and Fe₂O₃:Ti electrodes annealed with 90 s or 10 min. It indicated that with a fast annealing of 90 s, both Fe₂O₃ and Fe₂O₃:Ti nanorods showed a porous structure with 10–50 nm holes in the rods. As can be seen in Fig. 2b, the TiCl₄ solution treatment did not lead to a significant change of morphology. While for both Fe₂O₃:Ti sample and pure Fe₂O₃ sample annealed for 10 min, the surface became smoother to lower the surface free energy.

Fig. 2 FESEM images of thin films annealed at 800 °C for different time. (a) Pure Fe₂O₃ films for 90 s and (b) Fe₂O₃:Ti films for 90 s. (c) Pure Fe₂O₃ films for 10 min and (d) Fe₂O₃:Ti films for 10 min.

X-ray diffraction pattern analysis was performed to study the crystal structure of the nanowire samples with different annealing process, as shown in Fig. 3. It showed that there was little difference between pure Fe₂O₃ films and Fe₂O₃:Ti films. The strongest diffraction peak of [110] provided evidence that these nanowires grew along the [110] direction. No peak of TiO₂ was observed for the Fe₂O₃:Ti samples which could be due to the thin thickness of the TiO₂ shell that was not detectable for XRD measurements. For the 10 min annealed samples, the intensities of SnO₂ peaks increased while intensities of Fe₂O₃ peaks decreased compared to those of the 90 s annealed samples. This could be due to the annealing effect which converted the nanowires into debris by recrystallization, thereby decreasing the X-ray diffraction of (110) crystal plane.

Fig. 3 X-ray diffraction patterns of pure Fe₂O₃ films and Fe₂O₃:Ti films annealed at 800 °C for 90 s and 10 min (SnO₂: JCPDS #46-1088, Fe₂O₃: JPCDS #13-0534).

To confirm the successful doping of Ti element into the hematite, XPS spectra of Ti doped sample after Ar⁺ etching (4000 V, 4 μA cm⁻²) for different time were recorded and shown in Fig. S5.† It was found that before etching, only peak of Ti 2p_3/2 for TiO₂ (458.50 eV) was detected, while after etching for 940 s, the peak of Ti 2p_3/2 for Ti dopants (458.00 eV) emerged, which could be due to the exposed Ti dopants inside the hematite rods after etching.¹⁸ For sample etched for 1540 s, the peak intensity of Ti 2p_3/2 for Ti dopants further increased. This result indicated that Ti was successfully incorporated into hematite nanorod by diffusion.

Fig. 4 displays the photocurrents of pure Fe₂O₃ and Fe₂O₃:Ti thin films annealed at 800 °C for 90 s and 10 min. For pure samples, the photocurrent density of the sample annealed for 90 s was ∼0.05 mA cm⁻² at 1.0 V vs. Ag/AgCl and didn't increase much with bias potential increasing. When the annealing time was prolonged to 10 min, the photocurrent density increased to ∼0.3 mA cm⁻² at 1.0 V vs. Ag/AgCl. This improvement could be ascribed to Sn doping¹ from the FTO glass into the Fe₂O₃ rods⁴² during the relatively longer time annealing at high temperature. While for the case of Fe₂O₃:Ti thin films, different results were obtained. The photocurrent density of Fe₂O₃:Ti sample annealed for 90 s reached ∼1.1 mA cm⁻² at 1.0 V vs. Ag/AgCl (almost 22 times higher than pure sample). However, with an extended annealing time of 10 min, its photocurrent density decreased sharply (to ∼0.4 mA cm⁻² at 1.0 V vs. Ag/AgCl) but it was still higher than that of pure sample annealed for 10 min. When compare the photocurrent density of the two samples annealed for 90 s, one can draw the conclusion that the TiO₂ shell plays an important role in the improvement of hematite's PEC performance. The layer of TiO₂ coating could reduce the negative influence by surface states of hematite.⁴⁵ Such surface treatment was also used with Fe_xSn_1−xO₄ as a surface passivation layer on hematite rods to improve PEC performance.⁴⁶ As for the decline in photocurrent of Fe₂O₃:Ti sample with prolonged annealing time, there are two possible explanations for this phenomenon. One is that the conductivity of the FTO substrate decreases (the resistance of the FTO substrate increases from 22 to 167 Ω after 4 h of annealing at 700 °C)¹² and the other one is that the Sn diffusing from the FTO glass into the Fe₂O₃ rods. We measured the resistances of bare FTO substrate before and after the high temperature annealing treatment. It was found that the resistance of the FTO increased from 15 Ω □⁻¹ (4 cm × 1.4 cm, the long side) to 20.4 Ω □⁻¹ after annealing at 800 °C for 90 s (with an additional heat up section for 60 s) and 78 Ω □⁻¹ after 10 min. The increase of resistivity of FTO substrate can be due to loss of Sn by diffusion and distortion of the glass substrate upon high temperature annealing.⁴⁷ In addition to the resistance increase, both Fe₂O₃:Ti sample and pure Fe₂O₃ sample showed a great change in surface morphology after 10 min annealing at 800 °C, as shown in Fig. 2. This change in nanostructure led to a decrease in specific area of the samples, which could had a marked impact on the absorption of light and charge collection pathways.⁷

Fig. 4 The photocurrents of pure Fe₂O₃ and Fe₂O₃:Ti thin films with different annealing time at 800 °C.

The Mott–schottky of pure Fe₂O₃ and Fe₂O₃:Ti structured thin films with different annealing time at 800 °C were also measured and shown in Fig. 5. After doping with Ti element, a positive shift of flat-band potential of 0.09 V was found for the samples and a nearly two-order increase of charge carrier density was observed, with the N_d values listed in Table 1. The charge transfer process at the semiconductor–electrolyte interface was also analyzed with electrochemical impedance spectroscopy (EIS) as shown in Fig. 6, and the equivalent circuit parameters were summarized in Table 1. The series resistance (R_S) increased significantly with thermal treatment time increased for both pure Fe₂O₃ and Fe₂O₃:Ti samples due to the conductivity loss of the FTO substrate. R_CT1 and CPE1 are considered as internal resistance and capacitance of depletion region in the films, respectively. R_CT2 and CPE2 are referred to charge transfer resistance and double layer capacitance at the semiconductor–electrolyte interface, respectively. For pure Fe₂O₃ samples, both R_CT1 and R_CT2 decreased significantly when the annealing time increased from 90 s to 10 min which could be a result of Sn diffusion from the FTO substrate. However, as for the Fe₂O₃:Ti samples, it was unexpected that both R_CT1 and R_CT2 increased with annealing time prolonged from 90 s to 10 min. This increase should be a result of TiO₂ shell affecting the charge transfer inside the Fe₂O₃ nanowire and at the semiconductor–electrolyte interface. To look into the insight of annealing effects on pure Fe₂O₃ and Fe₂O₃:Ti samples, X-ray photoelectron spectroscopy (XPS) spectra were recorded to determine the interaction of Ti shell and Fe₂O₃ nanowire upon the thermal treatment (Table 2).

Fig. 5 Mott–schottky plot of pure Fe₂O₃ and Fe₂O₃:Ti structured thin films with different annealing time at 800 °C.

Table 1 Parameters of Mott–Schottky and the equivalent circuit model based on EIS data

Sample	Nd (cm^−3)	RS (Ω cm^−2)	CPE1 (F cm^−2)	RCT1 (Ω cm^−2)	CPE2 (F cm^−2)	RCT2 (Ω cm^−2)
P-800 for 90 s	1.06 ± 10^17	78.35	2.73 ± 10^−5	1768	4.20 ± 10^−6	879.7
P-800 for 10 min	4.45 ± 10^17	122.6	8.16 ± 10^−6	181.4	1.46 ± 10^−4	250.9
T-800 for 90 s	8.19 ± 10^18	84.78	5.64 ± 10^−5	105.8	2.09 ± 10^−4	161.1
T-800 for 10 min	6.63 ± 10^18	117.2	5.72 ± 10^−5	130.5	1.12 ± 10^−4	306.2

---

Fig. 6 Nyquist plots of the EIS measurements on the pure Fe₂O₃ and Fe₂O₃:Ti films (in 0.5 M Na₂SO₄ aqueous electrolyte).

Table 2 Binding energies, valence band potential and Sn/Fe, Ti/Fe atomic ratios of pure Fe₂O₃ and Fe₂O₃:Ti thin films with different annealing time at 800 °C obtained from XPS results

Sample	Fe 2p3/2	O 1s	Sn 3d	Ti 2p	VB	Sn/Fe	Ti/Fe
P-800 for 90 s	710.80	529.89	486.71	—	1.52	0.058	—
P-800 for 10 min	710.90	529.91	486.72	—	1.62	0.107	—
T-800 for 90 s	711.00	529.79	486.66	458.40	1.70	0.049	1.75
T-800 for 10 min	711.13	529.84	486.70	458.44	1.77	0.176	1.69

---

The stoichiometry and the electronic structure of pure Fe₂O₃ and Fe₂O₃:Ti thin films with different annealing time at 800 °C were determined by means of in situ XPS measurements. Fe 2p, O 1s, Sn 3d, Ti 2p core levels and valence band spectra are shown in Fig. S6.† Similarly, Fe 2p_3/2 and the Fe 2p_1/2 core levels for all samples were observed. More precisely, the binding energies for the Fe 2p_3/2 core level and the Fe 2p_1/2 showed an increase of 0.1 eV when annealing increased from 90 s to 10 min, and the Ti treatment also increased by 0.1 eV when comparing to the pure one under same annealing condition (as shown in Fig. S6a–d†). These values are consistent with typical values observed for Fe₂O₃ as other article reported.¹² O 1s XPS spectra of pure Fe₂O₃ and Fe₂O₃:Ti films are presented in Fig. S6e–h.† The binding energy of O 1s line (corresponding to Fe–O bonds) is 529.89 eV for pure Fe₂O₃ sample and 529.79 eV for Fe₂O₃:Ti sample with annealing time of 90 s, this shift could be a result of Ti doping which is consistent with the reported results.^12,48 The minor peaks at binding energies of 531.10 and 532.20 eV for both samples correspond to C=O and C–O bonds respectively.⁴⁹ By carefully examination of binding energy of Ti 2p_3/2 of sample Fe₂O₃:Ti, two binding energy peaks were obtained by fitting the Ti 2p peak. As shown in Fig. S6o,† these two peaks are at 458.39 eV and 459.26 eV. The 458.39 eV is consistent with typical values reported for Ti 2p_3/2 of Ti doped hematite, while 459.26 eV can be corresponding to Ti 2p_3/2 of TiO₂.^12,20,50 This result confirmed the successful incorporation of Ti element into Fe₂O₃ nanorod from the TiO₂ layer. For both pure Fe₂O₃ and Fe₂O₃:Ti thin films, longer annealing time increased Sn/Fe ratio which should due to Sn diffusion into Fe₂O₃ nanorod from the FTO substrate.^42,43 The valence band of the Ti doped samples is shifted by 0.18 eV and 0.16 eV toward higher binding energies with respect to the undoped sample annealed for 90 s and 10 min, respectively, which is inconsistent with observation by Magnan et al.⁵¹

Based on the results above, one can infer that Ti is an effective doping element. With a short time of 90 s, Ti was successfully incorporated into Fe₂O₃, as proved by XPS and Mott–schottky results. With annealing time increasing to 10 min, the photocurrent dropped, which could be due to the deterioration of FTO conductivity. For Ti doping, only a short annealing time is needed. While Sn doping by thermal diffusion seems like requiring a longer time. With annealing time prolonged from 90 s to 10 min at 800 °C, the photocurrent of pure Fe₂O₃ sample further enhanced even with the deterioration of FTO conductivity, which should be ascribed to increase of Sn doping from thermal diffusion. To compare the difficulties of Ti and Sn doping, the energies of formation for the neutral Ti and Sn doping in Fe₂O₃ were calculated according to formula (1) and plotted as a function of oxygen chemical potential in Fig. 7. The formation energy decreased with oxygen chemical potential from −4.52 eV at O-rich condition to −7.18 eV at O-poor condition, suggesting that low oxygen pressure is preferred for incorporation of foreign elements into Fe₂O₃. This result agrees with the conclusion we reported previously.⁴⁰ More importantly, the formation energy for Ti doping is always smaller than that for Sn doping throughout the allowed oxygen chemical potential, indicating that Ti dopant is easier to be incorporated into Fe₂O₃ than Sn dopant. This result well supports different behavior of thermal activation by Ti doping and Sn doping observed in this experiment.

Fig. 7 Energy of formation for Ti and Sn doping in Fe₂O₃ as a function of allowed chemical potential of oxygen. Oxygen chemical potential at O-rich condition was taken as half the energy of molecular O₂, while it at O-poor condition was determined when Fe₂O₃ transforming to Fe₃O₄.

Conclusions

In conclusion, thermal annealing effect on the interfacial property of Fe₂O₃:Ti thin films, which were synthesized by hydrothermal and TiCl₄ precursor soaking treatment, were investigated to look into the behavior of Ti doping upon postsynthesis thermal treatment to search for optimal doping strategy and procedure. The doping procedure was optimized by TiCl₄ precursor concentration and soaking time. The postsynthesis annealing time were also optimized regarding their photocurrents. By carefully comparing the structural, optical and photoelectrochemical properties of pure Fe₂O₃ and Fe₂O₃:Ti thin films under different annealing time, the trade-off between crystallinity improvement of Fe₂O₃ nanowires, doping effect and FTO conductivity decreasing upon the annealing process was discussed. The thermal annealing facilitates the Ti incorporation through diffusion from the surface of Fe₂O₃ nanowires and Sn incorporation through diffusion from FTO substrate. The Ti incorporation was found require much shorter annealing time than Sn incorporation. The theoretical calculation shows that the formation energy for Ti doping is always smaller than that for Sn doping throughout the allowed oxygen chemical potential.

Acknowledgements

This work was financially supported by the National Natural Science Foundation of China (no. 51202186, 51236007) and the Fundamental Research Funds for the Central University (xjj2016039).

References

1. K. Sivula, F. Le Formal and M. Gratzel, ChemSusChem, 2011, 4, 432–449.
2. A. Fujishima and K. Honda, Nature, 1972, 238, 37–38.
3. B. Liu and E. S. Aydil, J. Am. Chem. Soc., 2009, 131, 3985–3990.
4. L. Vayssieres, Adv. Mater., 2003, 15, 464–466.
5. J. Z. Su, L. J. Guo, N. Z. Bao and C. A. Grimes, Nano Lett., 2011, 11, 1928–1933.
6. J. Z. Su, X. J. Feng, J. D. Sloppy, L. J. Guo and C. A. Grimes, Nano Lett., 2011, 11, 203–208.
7. T. W. Hamann, Dalton Trans., 2012, 41, 7830–7834.
8. K. Sivula, F. Le Formal and M. Grätzel, ChemSusChem, 2011, 4, 432–449.
9. R. Franking, L. Li, M. A. Lukowski, F. Meng, Y. Tan, R. J. Hamers and S. Jin, Energy Environ. Sci., 2013, 6, 500–512.
10. C. X. Kronawitter, I. Zegkinoglou, S. H. Shen, P. Liao, I. S. Cho, O. Zandi, Y. S. Liu, K. Lashgari, G. Westin, J. H. Guo, F. J. Himpsel, E. A. Carter, X. L. Zheng, T. W. Hamann, B. E. Koel, S. S. Mao and L. Vayssieres, Energy Environ. Sci., 2014, 7, 3100–3121.
11. J. H. Kennedy and K. W. Frese, J. Electrochem. Soc., 1978, 125, 709–714.
12. G. Wang, Y. Ling, D. A. Wheeler, K. E. N. George, K. Horsley, C. Heske, J. Z. Zhang and Y. Li, Nano Lett., 2011, 11, 3503–3509.
13. P. Zhang, A. Kleiman-Shwarsctein, Y.-S. Hu, J. Lefton, S. Sharma, A. J. Forman and E. McFarland, Energy Environ. Sci., 2011, 4, 1020.
14. O. Zandi, B. M. Klahr and T. W. Hamann, Energy Environ. Sci., 2013, 6, 634–642.
15. M. Rioult, R. Belkhou, H. Magnan, D. Stanescu, S. Stanescu, F. Maccherozzi, C. Rountree and A. Barbier, Surf. Sci., 2015, 641, 310–313.
16. M. Rioult, H. Magnan, D. Stanescu and A. Barbier, J. Phys. Chem. C, 2014, 118, 3007–3014.
17. R. Franking, L. Li, M. A. Lukowski, F. Meng, Y. Tan, R. J. Hamers and S. Jin, Energy Environ. Sci., 2013, 6, 500–512.
18. D. Cao, W. Luo, J. Feng, X. Zhao, Z. Li and Z. Zou, Energy Environ. Sci., 2014, 7, 752–759.
19. J. A. Glasscock, P. R. F. Barnes, I. C. Plumb and N. Savvides, J. Phys. Chem. C, 2007, 111, 16477–16488.
20. A. Pu, J. Deng, M. Li, J. Gao, H. Zhang, Y. Hao, J. Zhong and X. Sun, J. Mater. Chem. A, 2014, 2, 2491–2497.
21. C. Miao, T. Shi, G. Xu, S. Ji and C. Ye, ACS Appl. Mater. Interfaces, 2013, 5, 1310–1316.
22. M. Rahman, N. Wadnerkar, N. J. English and J. MacElroy, Chem. Phys. Lett., 2014, 592, 242–246.
23. A. J. Abel, I. Garcia-Torregrosa, A. M. Patel, B. Opasanont and J. B. Baxter, J. Phys. Chem. C, 2015, 119, 4454–4465.
24. N. T. Hahn and C. B. Mullins, Chem. Mater., 2010, 22, 6474–6482.
25. C. Miao, S. Ji, G. Xu, G. Liu, L. Zhang and C. Ye, ACS Appl. Mater. Interfaces, 2012, 4, 4428–4433.
26. J. J. Deng, X. X. Lv, J. Y. Liu, H. Zhang, K. Q. Nie, C. H. Hong, J. O. Wang, X. H. Sun, J. Zhong and S. T. Lee, ACS Nano, 2015, 9, 5348–5356.
27. D. Barreca, G. Carraro, A. Gasparotto, C. Maccato, M. E. A. Warwick, K. Kaunisto, C. Sada, S. Turner, Y. Gonullu, T. P. Ruoko, L. Borgese, E. Bontempi, G. Van Tendeloo, H. Lemmetyinen and S. Mathur, Adv. Mater. Interfaces, 2015, 2, 1500313.
28. Y. Liu, F. Y. Su, Y. X. Yu and W. D. Zhang, Int. J. Hydrogen Energy, 2016, 41, 7270–7279.
29. X. Xie, K. Li and W. D. Zhang, RSC Adv., 2016, 6, 74234–74240.
30. Y.-S. Hu, A. Kleiman-Shwarsctein, G. D. Stucky and E. W. McFarland, Chem. Commun., 2009, 2652,  10.1039/b901135h.
31. D. P. Cao, W. J. Luo, J. Y. Feng, X. Zhao, Z. S. Li and Z. G. Zou, Energy Environ. Sci., 2014, 7, 752–759.
32. Z. Fu, T. Jiang, L. Zhang, B. Liu, D. Wang, L. Wang and T. Xie, J. Mater. Chem. A, 2014, 2, 13705–13712.
33. N. Mirbagheri, D. Wang, C. Peng, J. Wang, Q. Huang, C. Fan and E. E. Ferapontova, ACS Catal., 2014, 4, 2006–2015.
34. C. J. Sartoretti, B. D. Alexander, R. Solarska, I. A. Rutkowska, J. Augustynski and R. Cerny, J. Phys. Chem. B, 2005, 109, 13685–13692.
35. Y. Yang, M. Forster, Y. Ling, G. Wang, T. Zhai, Y. Tong, A. J. Cowan and Y. Li, Angew. Chem., Int. Ed., 2016, 55, 3403–3407.
36. J. Fan, Y. Hao, C. Munuera, M. Garcia-Hernandez, F. Güell, E. M. Johansson, G. Boschloo, A. Hagfeldt and A. Cabot, J. Phys. Chem. C, 2013, 117, 16349–16356.
37. C. Ng, Y. H. Ng, A. Iwase and R. Amal, ACS Appl. Mater. Interfaces, 2013, 5, 5269–5275.
38. C. Toussaint, H. L. L. Tran, P. Colson, J. Dewalque, B. Vertruyen, B. Gilbert, N. D. Nguyen, R. Cloots and C. Henrist, J. Phys. Chem. C, 2015, 119, 1642–1650.
39. L. Vayssieres, N. Beermann, S. E. Lindquist and A. Hagfeldt, Chem. Mater., 2001, 13, 233–235.
40. Z. Zhou, P. Huo, L. Guo and O. V. Prezhdo, J. Phys. Chem. C, 2015, 119, 26303–26310.
41. A. Annamalai, A. Subramanian, U. Kang, H. Park, S. H. Choi and J. S. Jang, J. Phys. Chem. C, 2015, 119, 3810–3817.
42. L. Wang, C.-Y. Lee and P. Schmuki, Electrochem. Commun., 2013, 30, 21–25.
43. C. X. Kronawitter, I. Zegkinoglou, C. Rogero, J. H. Guo, S. S. Mao, F. J. Himpsel and L. Vayssieres, J. Phys. Chem. C, 2012, 116, 22780–22785.
44. K. Sivula, R. Zboril, F. Le Formal, R. Robert, A. Weidenkaff, J. Tucek, J. Frydrych and M. Gratzel, J. Am. Chem. Soc., 2010, 132, 7436–7444.
45. X. G. Yang, R. Liu, C. Du, P. C. Dai, Z. Zheng and D. W. Wang, ACS Appl. Mater. Interfaces, 2014, 6, 12005–12011.
46. L. F. Xi, S. Y. Chiam, W. F. Mak, P. D. Tran, J. Barber, S. C. J. Loo and L. H. Wong, Chem. Sci., 2013, 4, 164–169.
47. Y. C. Ling, G. M. Wang, D. A. Wheeler, J. Z. Zhang and Y. Li, Nano Lett., 2011, 11, 2119–2125.
48. P. Mills and J. L. Sullivan, J. Phys. D: Appl. Phys., 1983, 16, 723–732.
49. T. A. Loh and D. H. Chua, ECS J. Solid State Sci. Technol., 2014, 3, M11–M17.
50. B. Wang, Y. Song, W. Ren, W. Xu and H. Cui, J. Sol-Gel Sci. Technol., 2009, 51, 119–123.
51. H. Magnan, D. Stanescu, M. Rioult, E. Fonda and A. Barbier, Appl. Phys. Lett., 2012, 101, 133908.

---

Footnote

† Electronic supplementary information (ESI) available: Photocurrent plots, EIS result and XPS spectra. See DOI: 10.1039/c6ra19699c

---