Various types of iron oxide and their physical properties (ρ: electrical resistivity, TN: Néel temperature, and Eg: bandgap energy).
Epitaxial films of Rh-substituted α-Fe2O3 were fabricated by a pulsed laser deposition technique, and their photoelectrochemical characteristics were investigated for the development of visible light-responsive photoanodes for water splitting. The photocurrent in the films upon irradiation in the visible region was significantly enhanced after Rh substitution. Moreover, a near-infrared photocurrent was clearly observed for Rh:Fe2O3 photoanodes, whereas no photoresponse could be detected for the α-Fe2O3 films. These improved photoelectrochemical properties are attributed to the increased light absorption due to the hybridization of Rh-4d states and O-2p states at the valence band maximum. Moreover, Rh substitution also strongly influences the photocarrier transport properties of the films. The electrical conductivity of Rh:Fe2O3 is higher than that for α-Fe2O3 by two orders of magnitude, which is possibly due to the extended 4d orbitals of the Rh3+ ions. Thus, the improved electrical properties may lead to an increased photocurrent by lowering the recombination rate of photogenerated carriers.
- solar water splitting
- pulsed laser deposition
- photoelectrochemical cell
- iron oxides
- bandgap engineering
Iron oxides are well known to have various physical properties depending on their composition and crystal structures (see Table 1). They have been the subject of extensive investigation over the past decades from both fundamental and practical perspectives. For example, magnetite (Fe3O4) has been one of the most widely investigated oxides in various research fields owing to its high magnetic transition temperature (~585°C) and high spin polarization of carriers [1, 2, 3]. Numerous Fe3O4-based ferromagnetic semiconductors and related spintronics devices have been reported. Another simple iron oxide, wüstite (FeO) has attracted much attention in various fields such as Earth sciences, oxide electronics, spintronics, and chemical engineering [4, 5, 6]. Moreover, multifunctional bismuth ferrite (BiFeO3, BFO) has been of great interest owing to its potential applications in numerous room temperature multiferroic devices [7, 8, 9]. BFO is also considered to be a good candidate for use in solar energy conversion systems because of its electrical polarization-induced photovoltaic effects . The triangular antiferromagnet RFe2O4 (R = Ho, Y, Yb, Lu, and In) is a multilayered oxide and was discovered in the 1970s by Kimizuka et al. . RFe2O4 is composed of alternating hexagonal Fe─O and R─O layers stacked along the c-axis, and Fe2+/Fe3+ charge order occurs in the Fe─O layers below 320 K, which is followed by magnetic ordering below ~240 K . Recently, a number of studies on RFe2O4 have been stimulated by the discovery of the giant magnetoelectric response in LuFe2O4 and its application to multiferroic devices is currently the subject of extensive investigations [13, 14]. A great number of investigations on the magneto-optical (MO) properties of garnet-type ferrites (R3Fe5O12) have been carried out for applications in the field of optical communications. They are currently recognized as the most promising materials in magnonics and related areas. Especially, they are widely used in ferromagnetic resonance experiments and magnon-based Bose-Einstein-condensates owing to their extremely low damping [15, 16, 17, 18]. Furthermore, there has been much interest in hexaferrites, MFe12O19 (M = Ba and Sr) [19, 20], which are commonly applied in a wide variety of data storage and recording devices. One of the most favorable characteristics of the above iron oxides is their chemical stability, and they are also nontoxic. Moreover, iron and oxygen are abundant in the Earth. These features imply that iron oxides are favorable materials for applications in environmentally friendly electronics, spintronics, and magnonics. The author focuses on α-Fe2O3 commonly referred as a hematite, which is known as a promising candidate for semiconductor photoanodes for photoelectrochemical (PEC) water splitting [21, 22, 23]. A schematic of a PEC cell is shown in Figure 1. They consist of a photoactive electrode and a metal counter electrode immersed in a suitable electrolyte solution. The photogenerated electron-hole pairs are split by the electric field in the space-charge region at the surface of photoelectrodes. Since Honda and Fujishima’s pioneered work on PEC water splitting with a TiO2 photoelectrode , there has been worldwide research focused on the solar generation of hydrogen as a renewable and clean energy source. Many kinds of materials including TiO2 have been investigated for their application as photoelectrodes. However, most of them are wide-gap semiconductors, and only a small fraction of the solar spectrum can be utilized by the PEC cells based on these materials. A high PEC responsivity to visible (VIS) and near-infrared (near-IR) light is required to harvest the lower energy region of the solar spectrum. From this viewpoint, α-Fe2O3 has attracted much attention because of its promising properties for application as a photoanode in a solar water splitting cell. It possesses a narrow bandgap energy (Eg) of 2.1 eV that allows for the absorption of up to 40% of solar spectrum. However, the reported efficiencies for PEC water splitting using α-Fe2O3-based photoelectrodes are significantly low. This poor PEC property of α–Fe2O3can be attributed to the short diffusion length of the photogenerated holes. For α–Fe2O3-based PEC cells, only the holes generated near the electrolyte/photoanode interface can oxidize water [25, 26]. That is, most of the photogenerated electron-hole pairs recombine before reaching the photoelectrode surface. The hematite lattice is composed of an alternating stack of Fe bilayers and O layers along c-axis as illustrated in Figure 2. Spins of Fe3+ ions within each bilayer are parallel, whereas adjacent Fe bilayers have opposite spins. 3d electrons of Fe can move by hopping via the change in the Fe2+/Fe3+ valence within the Fe bilayers, whereas the exchange of electrons between neighboring Fe bilayers is spin forbidden [27, 28, 29]. Therefore, the orientation of a highly conducting (001) plane vertical to the substrate will facilitate the collection of photogenerated carriers and suppress their recombination. The author employed a Ta-doped SnO2 (TTO) layer grown on α–Al2O3 (110) single-crystal substrates for the epitaxial growth of α–Fe2O3 films along the  direction. As shown in Figure 3(a), the SnO2 (101) plane matches the α–Fe2O3 (110) plane with a lattice mismatch of approximately 1.3%, which is favorable for the epitaxial growth of hematite along the  direction on the SnO2(101)/α–Al2O3(101) substrate . Another issue regarding α-Fe2O3 concerns its low responsivity to near-IR light. It is well known that the photocurrent in α-Fe2O3 is maximized at a wavelength (λ) of ~350 nm, exhibits a significant decrease with increasing λ, and approaches zero at approximately 600 nm, corresponding to its bandgap . An improvement of the PEC responsivity in VIS and near-IR regions by controlling the bandgap would be useful for solar energy harvesting. Unfortunately, there exist few reports on such bandgap engineering in α-Fe2O3. From this viewpoint, the author focused on Rh-substituted α-Fe2O3 (FRO). α-Rh2O3 has a bandgap Eg of 1.2–1.4 eV  and the same corundum-type crystal structure as α-Fe2O3. Therefore, the bandgap of α-Fe2O3 could be narrowed by Rh substitution in the films [33, 34]. Figure 3(b) shows a schematic of the band alignment of FRO [35, 36]. α-Fe2O3 is a charge transfer-type insulator with a bandgap between the Fe 3d state (upper Hubbard band) and the fully occupied O 2p state. In contrast, the bandgap of α-Rh2O3 originates from the ligand field splitting of the Rh 4d orbitals. The Rh 4d (t2g) band in α-Rh2O3 lies near the O 2p band, and they effectively hybridize at the valence band maximum (VBM) [32, 37, 38]. In this chapter, the PEC characteristics of FRO photoanodes fabricated using pulsed laser deposition (PLD) are discussed in association with their electronic structures.
2. Experimental procedures
The FRO films were grown using a PLD technique with an argon fluoride (ArF) excimer laser (λ = 193 nm). The laser pulse frequency was 5 Hz. The fluence remained constant at 1.1 J/cm2. The typical growth rate of the films was 0.5 nm/min. After deposition, the FRO films were annealed in air at 700°C to improve their crystallinity. The author employed two types of bottom electrodes, viz., TTO deposited onto an α-Al2O3 (110) substrate and polycrystalline fluorine-doped SnO2 (FTO) formed on a soda-lime glass substrate. An Fe2-xRhxO3 (x = 0.0–2.0) pellet prepared by a solid-state reaction was used as a target for PLD. The growth temperature was kept at 700 and 800°C for the FRO and TTO films, respectively. The crystal structures of the samples were confirmed using X-ray diffraction (XRD). In the PEC measurement, the I-V properties were measured using an electrochemical analyzer under the illumination of Xe lamp (500 W). Optical measurements were conducted using a Vis-UV spectrometer. X-ray photoemission spectroscopy (XPS) was performed to evaluate the structure of the valence band (VB) in the FRO films.
3. Crystal structures
The XRD 2 theta-omega scan of the FeRhO3 films is shown in Figure 4(a). For the as-deposited sample, broad peaks are observed at 35 and 75°, which are ascribed to the (110) and (220) reflections of corundum-type FRO, respectively. This indicates that the films grown along  despite their low crystalline quality. Sharp peaks appear after thermal annealing, suggesting an improvement in the crystallinity. The in-plane epitaxial relationship was evaluated to be TTO //FRO  by in-plane XRD measurements. This result agrees with the atomic configurations in Figure 3(a) . The lattice constants obey Vegard’s law, implying that Fe had been appropriately substituted with Rh. In contrast to the films deposited onto the sapphire substrates, the films deposited on the glass substrates are polycrystalline in nature, as shown in Figure 4(b) and (c).
4. Optical properties
Figure 5(a) shows the light absorption spectra of the films. The fundamental absorption edge of α-Fe2O3 is related to charge transfer from O 2p states to the upper Hubbard band promoted by photons [denoted by TCT in Figure 3(b)]. For films with a higher Rh content, a broadband appears at 1.5–4.5 eV that is possibly related to α-Rh2O3. The optical transition in α-Rh2O3 is unclear; its absorption edge is considered to be associated with the d-d transition of Rh3+, judging from the bandgap structure [31, 39]. Figure 5(b) shows the values of an indirect bandgap (Eg), which were derived from the Tauc relation, αhν ∝ (hν − Eg)2 (α: optical absorption coefficient and hν: photon energy). Eg decreases as the content of Rh in the films increases according to the above discussion. Value of Eg of 2.1 and 1.2 eV were obtained for α-Fe2O3 and α-Rh2O3, respectively. These values are almost identical to those reported for polycrystalline films .
5. XPS spectroscopy
The results of XPS are presented in Figure 6. In the spectra of Fe 2p core level (Figure 6(a)), main peaks are at around 710 and 723 eV and are assigned to Fe 2p2/3 and 2p1/3 orbitals of α-Fe2O3, respectively [41, 42, 43]. These core level peaks become weaker as the Rh content increases. In turn, new distinct peaks appears at approximately 310 and 315 eV, which are assigned to Rh 3d3/2 and 3d5/2 orbitals, respectively [44, 45]. As seen in Figure 6(c), the VBM of α-Fe2O3 is estimated to be 0.65 eV. In contrast, the VBM of α-Rh2O3 is located near the Fermi level (~0.0 eV). Three distinct peaks are observed in the VB spectrum of the films. The bands centered at 1, 2, and 3 eV in the VB spectrum of α-Fe2O3 are assigned to the Fe 3eg, 2tg, and 2eg orbitals, respectively . The crystal field splitting energy between the Fe 3eg and 2tg orbitals was estimated to be 2.5 eV in a previous study , which agrees with the experimental value (2.4 eV) well. In the VB spectrum of α-Rh2O3, three distinct peaks similarly appear. Unfortunately, there are hardly any reports of VB spectra for α-Rh2O3. However, by comparing with results obtained by ultraviolet photoemission spectroscopy (UPS) for ZnRh2O4 [48, 49], the peak observed at 1 eV is attributed to the t2g orbitals of the RhO6 octahedra in α-Rh2O3. The peaks at 2 and 3 eV are assigned to the Rh 4d, 5 s, and 5p mixed states . The VBM and Eg exhibit a similar dependences on the Rh content, as shown in Figure 6(d). In addition, the change in the VBM (0.7 eV) for x = 1.0 is close to the bandgap decrease (0.8 eV) for x = 1.0. From these results, we can conclude that the bandgap decrease by Rh substitution is due to the hybridization of the O 2p valence band with the Rh t2g band at the VBM.
6. Photoelectrochemical properties
The current-potential curves of the films are shown in Figure 7(a) and (b). For α-Fe2O3, the photocurrent is 2.87 μA/cm2 at 0.5 V under irradiation with VIS light (λ = 400–700 nm). As shown in Figure 7(a), the VIS photocurrent is remarkably increased after Rh substitution (17.3 μA/cm2 at 0.5 V for x = 0.2). The effect of Rh substitution becomes more obvious with near-IR irradiation (λ = 700–900 nm). As shown in Figure 7(b), for x = 0.2, a near-IR photocurrent is evidently observed, whereas a photocurrent is hardly detected for the α-Fe2O3 film. These improved PEC properties of the FRO films may be caused by the increased light absorption in these wavelength regions. Furthermore, the electrical properties of the films also affect their PEC performance. The electrical conductivity for x = 0.2 (σ = 3.8 × 10−4 Ω−1 cm−1 at 300 K) is significantly larger than that for x = 0.0 (σ = 2.6 × 10−6 Ω−1 cm−1 at 300 K). This is possibly due to the extended profile of the Rh 4d state . Thus, the improved electrical conductivity possibly causes an increased photocurrent by lowering the recombination rate of the photocarriers as in the case for Ti- or Si-doped α-Fe2O3 [22, 38, 40]. In Figure 7(c), the spectra of the incident photon-to-current efficiency (IPCE) are shown. The IPCE was estimated using the following relationship: IPCE (%) = 100 × [hc/e] × I/[P × λ], where h, c, e, I(mA/cm2), and P(mW/cm2) denote the Planck constant, the light velocity, the elementary charge, the photocurrent, and the power of the illumination per unit area, respectively . The IPCE for x = 0.1 and 0.2 is much higher than that of α-Fe2O3 in the 340–850 nm wavelength region. For α-Fe2O3, the IPCE decreases to zero when the wavelength exceeds 610 nm, corresponding to its bandgap. On the other hand, for x = 0.2, the IPCE is 2.35% at 610 nm and gradually decreased to 0.11% at 850 nm as shown in the inset of Figure 7(c). The IPCE decreases drastically when x exceeds 0.2 as shown in Figure 7(c), and a photocurrent is hardly detected for x ≥ 0.75 in the 340–900 nm wavelength region. These results possibly reflect the drastic change in the optical transition process caused by Rh substitution. On the one hand, the photogenerated carriers in α-Fe2O3 diffuse through the bands related to the Fe 3d and O 2p states . On the other hand, the recombination probability of the carriers generated in α-Rh2O3 following the d-d transition in Rh3+ is significantly high [32, 51]. This nature impairs the PEC performance for a higher Rh content. We note that that the rate of decrease in Eg is drastically decreased when x exceeds 0.2. This result suggests that the band-edge electronic structure is not strongly influenced by Rh substitution; therefore, the d-d transition predominantly occurs for x > 0.2. The IPCE peak wavelength shifts from 350 to 430 nm as Rh content increases from x = 0.0 to 0.2. This is a desirable feature for energy harvesting, because the peak of the solar spectrum is at ~475 nm. The PEC properties of the polycrystalline film (x = 0.2) are shown in Figure 7(d). The photocurrent of the single-crystal FRO grown on a TTO/Al2O3 (110) substrate is higher than that of the polycrystalline FRO grown on FTO glass. This result can be explained by the electronic transport properties of the films. The conductivity of the (110)-oriented single-crystalline film along the out-of-plane direction (σ = 3.4 × 10−4 Ω−1 cm−1) is much larger than that of the polycrystalline film (σ = 8.8 × 10−6 Ω−1 cm−1). The improvement in the electrical conductivity in the single-crystal films may accelerate the collection of photocarriers, resulting in their enhanced PEC properties. In Figure 7(e), the Mott-Schottky plots are shown. The density of donors N is expressed as follows:
where e0 represents an electron charge, and ε0 and ε are the vacuum and relative electric permittivities, respectively. By employing the reported value of ε = 80 for hematite, the donor densities are calculated to be 4.2 × 1017 (x = 0.2) and 5.3 × 1017 cm−3 for x = 0.0 (hematite) and 0.2, respectively. Thus, the donor density does not significantly change after Rh substitution, in contrast to the case for Ti- or Si-doped α-Fe2O3. It is considered that the number of Fe2+ ions, which acts as electron donors, increases when Fe3+ is substituted with Ti4+ or Si4+ owing to the charge neutrality. In contrast, in the FRO films, the valence of the Rh ions is +3, and hence, the content of Fe2+ does not increase after Rh substitution. Nevertheless, the conductivity for x = 0.2 is two orders of magnitude larger than that for x = 0.0. Therefore, it is considered that the carrier mobility is remarkably increased after Rh substitution probably owing to the extended nature of the Rh 4d states.
The Rh-substituted α-Fe2O3 photoelectrodes were successfully fabricated using a pulsed laser deposition method. Their bandgap narrowed as the Rh content increased. XPS analyses revealed that the bandgap narrowing is brought by the hybridization of the Rh 4d state with the O2p–Fe 3d states at the VBM of α-Fe2O3. The photocurrent was significantly enhanced for lower Rh contents. Moreover, the PEC properties were improved by the control of crystal growth condition. These results will be utilized in the development of high-efficiency solar energy conversion devices based on iron oxides.
This work was supported by JSPS Core-to-Core Program, A. Advanced Research Networks, and JSPS KAKENHI grant numbers JP16K14226 and JP15H03563. The author would like to thank Prof. H. Tabata, Prof, H. Matsui, and Dr. H. Yamahara for their support and helpful discussion.