Single Crystalline Hematite Films for Solar Water Splitting: Ti-Doping

Jan 23, 2014 - Undoped and Ti-doped (2 at. %) epitaxial hematite thin films, in the thickness range 5–50 nm, were grown by atomic oxygen assisted mo...
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Single Crystalline Hematite Films for Solar Water Splitting: Ti-Doping and Thickness Effects Maxime Rioult, Hélène Magnan,* Dana Stanescu, and Antoine Barbier CEA-Saclay, DSM/IRAMIS/SPEC, F-91191 Gif-sur-Yvette Cedex, France ABSTRACT: Undoped and Ti-doped (2 at. %) epitaxial hematite thin films, in the thickness range 5−50 nm, were grown by atomic oxygen assisted molecular beam epitaxy (AO-MBE) on Pt(111) substrates in the framework of hydrogen harvesting from sunlight-induced water splitting. Such single crystalline samples are suitable model systems to study thickness and doping effects on the photoelectrochemical properties; we demonstrate that they also allow disentangling intrinsic transport properties from mingled overall properties due to the usually unknown contributions from morphology or crystalline structure defects. From their photoelectrochemical characteristics (I(V) curves, incident photon to current efficiency measurements, and electrochemical impedance spectroscopy), we evidence the existence of an optimum layer thickness, which is higher for Ti-doped samples (30 nm) as compared to undoped ones (20 nm). Our results suggest that this effect is due to an increase of the carrier concentration combined with higher carriers’ diffusion lengths in the doped samples stressing intrinsic modifications of the hematite layer upon titanium doping that cannot be accounted for by simple structural or electronic structure changes.

I. INTRODUCTION Solving the concomitant worldwide increasing energy consumption problem and the need for greenhouse gas reduction to avoid or limit climate changes imposes a higher fraction of renewable energies in the total energy production. Sustainable growth cannot anymore be dissociated from fostering novel renewable energies that become more and more a first-class issue. Within this framework, increasing attention is paid to the sunlight-assisted water splitting as a clean method of hydrogen production. As a matter of fact, hydrogen is an energy carrier of choice which does not lead to any greenhouse gas production. Although the idea of producing hydrogen using water splitting assisted by solar light is very seductive, it remains unfortunately also very tricky, and many materials science issues have to be solved. During the process, electron−hole pairs are generated in a semiconductor electrode, upon solar light absorption, and are subsequently used to promote the oxido-reduction reactions of water leading to oxygen production at the photoanode and hydrogen production at the photocathode.1,2 Since the pioneering discovery of water-photoassisted electrolysis using semiconducting TiO2 in 1972 by Fujushima and Honda,1 several materials were investigated as photoanodes3 where water oxidation occurs (2OH− + 2h+ → (1/ 2)O2 + H2O). Hematite, i.e., the α-Fe2O3 iron oxide, is one of the most promising materials regarding its characteristics. It is an n-type semiconductor with a quasi-ideal band gap (∼2.2 eV) for solar water splitting applications. Indeed this material is able to absorb ca. 40% of the solar light spectrum, and its theoretical solar-to-hydrogen conversion yield reaches 13%.3 It is abundant on earth and very stable in aqueous environments, which makes it a serious candidate in the framework of green energy production.4 Moreover, the valence band edge of hematite is © 2014 American Chemical Society

located below the H2O/O2 redox potential which favors the water oxidation reaction.5 Unfortunately, hematite has not only advantages since it has been demonstrated to have weak transport properties6,7 (low conductivity and low carrier lifetime) and poor surface kinetics.8 Also, the conduction band edge of α-Fe2O3 is not well positioned with respect to the potential of the water reduction reaction (2H2O + 2e+ → H2 + 2OH−), thus an external bias is necessary to promote water splitting.9 Various strategies were proposed to overcome hematite drawbacks3,4,10 including (i) doping with different elements such as Ti, Si, or Sn7,11−21 which can improve electrical properties and also modify the electronic structure; (ii) nanostructuration (e.g., nanowires, mesoporous materials, ...) to compensate small carrier diffusion lengths;5,22,23 (iii) more complex semiconductor heterostructures24,25 to optimize light absorption and photogenerated carrier separation; and finally (iv) surface engineering with overlayers or catalysts.8,26−28 The case of hematite doped with Ti is a seductive idea to improve hematite properties and has been the subject of numerous studies in recent years.7,11−18 However, the improvement of the electrical conductivity induced by Ti doping has not been discussed in detail with respect to possible changes of the crystalline structure. As a matter of fact, previous studies concerning polycrystalline films or nanostructures include a high density of grain boundaries which may dominate the electric conduction; they can either limit conduction (if the conduction occurs perpendicularly to them) or increase Received: January 10, 2014 Revised: January 20, 2014 Published: January 23, 2014 3007

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were observed and acquired during film growth to monitor the crystal quality and structure of the samples. In situ XPS spectra were systematically recorded just after deposition to determine the stoichiometry and the electronic structure of the films. More precisely, we recorded Fe 2p, Ti 2p, and O 1s core levels and the valence band region using Al Kα radiation. The doping level in atomic percentage of Ti (at. % Ti) is defined as follows: at. % Ti = I(Ti)/[I(Ti) + I(Fe)], where I(Ti) and I(Fe) are, respectively, the Ti 2p and Fe 2p integrated intensities corrected by the corresponding Scofield factors (cross sections). For all samples studied, the measured doping level is 2 at. % (±0.3 at. %). Additionally, X-ray reflectivity measurements were performed for thickness determination. Photoelectrochemistry. The photoelectrochemical response of our films was studied using a three-electrode cell. All electrochemical measurements were performed at room temperature using a NaOH 0.1 M (pH = 13) solution as electrolyte, a platinum wire as counter electrode, and a Ag/ AgCl electrode for the potential reference (VAg/AgCl = +0.197 V vs SHE). The sample was mounted as an anode (working electrode) using a dedicated sample holder that allows the contact only between the hematite surface and the electrolyte. The illumination source was a Newport 1000 W Xe Arc lamp with an infrared filter. The incident light flux is around 100 mW/cm2 (measured with a Newport 1918-R Power Meter). Potential control and current acquisition between the three electrodes were done using a Princeton Applied Research (PAR) 263A potentiostat controlled by computer. For photocurrent−voltage curves, I(V), the potential was swept from 0 to +0.8 V vs Ag/AgCl at a speed of 50 mV/s. The photocurrent is defined as the difference between the current recorded under light and the one without (dark). Incident photon to current efficiency (IPCE) measurements were carried out under monochromatic light, at a bias of 0.6 V vs Ag/AgCl. The IPCE can be calculated with18

conduction (if the conduction occurs along them). The importance of grain boundaries has been stressed by Glasscock et al.,7 for example, who proposed for polycrystalline magnetron-sputtered samples that the passivation of hematite grain boundaries and reduction of the crystalline quality by titanium doping could explain the photoelectrochemical property enhancement upon doping, whereas silicon doping showed only limited improvements supposedly because of a decrease of grain boundary size. For hematite films grown by hydrothermal methods, Miao et al.13 reported that Ti doping induces a change in nanostructure size and shape and Deng et al.14 a modification of the type of hematite nanostructure and crystallographic orientation in agreement with the works of Tang et al.15,16 and Hahn and Mullins.18 However, because of the polycrystalline nature of these films, the contributions due to structural, morphological, electronic, and crystallographic effects and the intrinsic effects due to Ti doping are necessarily entangled. It is thus not yet clear if the modifications of watersplitting properties upon Ti doping are linked to the change of crystallographic structure or not. The investigation of single crystalline samples is an elegant way to overcome this problem. For undoped hematite layers, the detailed behavior of the photoelectrochemical properties as a function of the layer thickness, especially for well-crystallized layers, has also been less considered. Previous studies used hematite photoanodes obtained by deposition techniques such as hydrothermal growth,13,14 chemical vapor deposition,5,8,9,26 or electrodeposition,20 patterned or not. In these approaches, it is very difficult to distinguish which parameter, including preparation issues, morphology, or crystal quality, etc., influences the photoelectrochemical properties. Moreover, other works using techniques suitable for model sample elaboration like atomic layer deposition12,19,28,29 or sputtering,7,15,16 focused mainly on dopants and nanostructuration influence or on surface kinetics properties. Therefore, the present work aims at studying thin epitaxial undoped and Ti-doped hematite films of different thicknesses as photoanodes. To investigate single crystalline electrodes as model samples, the layers were deposited on Pt(111) substrates using atomic oxygen plasma assisted molecular beam epitaxy (AO-MBE).6,30,31 These single crystalline samples allow us to investigate the modification of the intrinsic properties of hematite by Ti doping, without any change of crystallographic structure and/or morphology.

IPCE(λ) =

Jph (λ) hc × (%) λ eP(λ)

(1)

where Jph(λ) is the photocurrent density, P(λ) the incident power density, λ the wavelength of the incident light, h the Planck constant (6.62 × 10−34 J.s), c the light velocity (3 × 108 m/s), and e the elementary charge (1.6 × 10−19 C). For these measurements, we used a Cornerstone 130 model 74004 monochromator (Newport). The wavelength was varied between 200 and 1000 nm, with 10 nm steps. Transient variations of the photocurrent are expected to be due to (i) voltage application or to (ii) chopped incident light. To minimize the transient signals due to the voltage application, the potential was held constant during 200 s prior to IPCE acquisition. Since the values of the photocurrents in some hematite films were very low (1 kHz) were used to determine the capacitance C of the space charge layer at the surface−electrolyte interface. For this, we modeled our system with an RS−RC circuit of equivalent impedance Z given by Z = R S + (iωC + R−1)−1

(2)

where RS and R are the series and space charge layer resistances, respectively. The Mott−Schottky plots A2/C2 = f(V) were fitted using eq 3 to determine the flat band potential (Vfb) and the carrier concentration (ND)18 ⎛ k T⎞ A2 2 = × ⎜V − Vfb − B ⎟ 2 ⎝ eNDε0εr e ⎠ C

(3)

where A is the illuminated surface area, ε0 the dielectric constant of vacuum (8.85 × 10−12 F/m), εr the dielectric constant of hematite taken equal to 32,7,12 V the potential applied to the electrochemical cell, and kB the Boltzmann constant (1.38 × 10−23 J/K). The flat band potential corresponds to the difference between the electrochemical potentials of the semiconductor (Fermi level) and of the electrolyte (redox level).

Figure 1. (Left panel) RHEED patterns of a 50 nm thick 2 at. % Tidoped α-Fe2O3 film grown by AO-MBE at different deposition times/ thicknesses (steps a to e) over the two lowest index surface diffraction directions called, for simplicity, D1 and D2. (Right panel) The corresponding surface reciprocal lattices are represented, making explicit the diffraction directions. (Blue box), (red circle), and (green circle) represent the Pt(111), γ-Fe2O3(111), and α-Fe2O3(0001) surface reciprocal lattices, respectively, assuming a cube on cube and √3 × √3R30° epitaxial relationship between Pt(111) and γ-Fe2O3 and α-Fe2O3, respectively.

III. RESULTS AND DISCUSSION Structure and Chemistry. The growth of undoped hematite thin films by AO-MBE on Pt(111) single crystals has already been studied in detail in previous works.30,31 Here we describe the growth of Ti-doped hematite films on Pt(111) single crystals. RHEED patterns as well as corresponding surface reciprocal lattices in such films at different deposition times (different thicknesses) are presented on Figure 1, assuming the same epitaxial relationship between Pt(111) and γ-Fe2O3 or α-Fe2O3 as in ref 31 (i.e., cube on cube for γFe2O3(111)/Pt(111) and √3 × √3R30° for α-Fe2O3(0001)/ Pt(111)). We observe that the thin film growth is similar to the case of undoped samples31 and follows a Stranski−Krastanov growth mechanism (layer-plus-islands). First, a 2D Ti-doped γFe2O3 layer grows, evidenced by extra straight lines on the corresponding RHEED patterns (step b in Figure 1). A phase transition from γ-Fe2O3(111) to α-Fe2O3(0001) occurs between 3 and 4.5 nm thickness (steps c and d), where the γ and α phases coexist. After the γ → α transition, an epitaxial growth of Ti-doped hematite (α-Fe2O3) with some islands and extra roughness is observed (evidenced by the presence of spots along lines on RHEED patterns, step e in Figure 1). In addition to the structural phase identification, we can also derive from the RHEED patterns the relaxation of the in-plane lattice parameter. Since different phases appear during the growth and since they give rise to lattices belonging to different space groups, it is however not straightforward to determine lattice parameter evolution. Therefore, we prefer, with the same approach as Barbier et al.,31 to represent a generic lattice

parameter p which is calculated using the RHEED patterns obtained over the particular diffraction direction D1 (made explicit on Figure 1). This parameter p is equal to: • the in-plane lattice parameter in the case of Pt(111) (2.77 Å); • aγ‑Fe2O3/2 in the case of γ-Fe2O3 (where aγ‑Fe2O3 is the inplane lattice parameter of γ-Fe2O3(111), assuming a cube-on-cube epitaxial relationship with Pt(111)); • aα‑Fe2O3/√3 in the case of α-Fe2O3 (where aα‑Fe2O3 is the in-plane lattice parameter of α-Fe2O3(0001), assuming a √3 × √3R30° epitaxial relationship with Pt(111)). Although this parameter p does not correspond to the same lattice parameter for all crystallographic phases, its evolution as a function of the film thickness remains relevant for the overall lattice relaxation and is presented in Figure 2. The lattice parameter relaxation during the epitaxial growth process with respect to increasing film thickness is consistent with previous works.31 The parameter p converges to a value of 2.92 ± 0.02 Å, which corresponds to an in-plane lattice parameter of 5.06 ± 0.02 Å for hematite. This is very close, or even equal within the experimental error bar, to the bulk hematite lattice parameter value (5.04 Å). Figure 2 shows also that the film relaxes within the first ca. 20 nm. The stoichiometry and the electronic structure of our layers were determined by means of in situ XPS measurements. Fe 2p 3009

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doping levels. More precisely, we measured binding energies of 711.2 eV for the Fe 2p3/2 core level and of 719.2 eV for the typical Fe3+ shakeup satellite (dash-dotted gray lines on Figure 3a). This confirms the Fe3+ ionic state of iron in all films and the absence of detected Fe2+ species, excluding polaron hopping (Fe3+ + Ti4+ → Fe2+ + Ti3+)32 in our Ti-doped hematite samples. For Ti-doped samples, the Ti 2p3/2 line shows a narrow single feature at a binding energy of 458 eV (dash-dotted gray line on Figure 3b). This value is between the tabulated binding energies of Ti4+ and Ti3+ in TiO2 (458.8 and 457.1 eV, respectively). The crystallographic structure of 40 nm films has been studied in detail by EXAFS,11 and we have shown that Ti substitutes Fe in the hematite structure with a slight deformation of the oxygen octahedron. Since we do not observe any change of RHEED patterns or XPS spectra with thickness, we can reasonably assume that this substitution is also present in the thinner samples with a lower doping level studied here. Therefore, as in references 6 and 11, we can conclude that this binding energy corresponds to Ti4+ species included in a hematite host matrix. Valence band spectra (Figure 3c) did not show changes over thickness or Ti doping level variations. We observe neither a shift of the valence band like in ref 11 nor states in the gap like in ref 6. In a previous study11 we have shown that the valence band shifts toward higher binding energies (0.2 eV shift) when hematite is doped with higher Ti doping level (>5 at. %). The present results show that a lower doping level, at least up to 2 at. %, does not induce a detectable modification in the electronic structure. Photoelectrochemical Properties. Figure 4 presents the photocurrent density vs voltage curves for undoped and Tidoped hematite films of various thicknesses. We see that the photocurrent density obtained for Ti-doped samples is higher than for undoped ones by 2 orders of magnitude, which is consistent with previous results.7,15,16 Let us note that the increase of photocurrent with Ti doping in the present study is

Figure 2. In-plane lattice parameter p derived from the D1 direction in RHEED patterns as a function of the film thickness for a 2 at. % Tidoped α-Fe2O3(0001) film grown by AO-MBE. Dashed lines stand for the lattice parameter p of bulk γ-Fe2O3, α-Fe2O3, and Pt(111).

and Ti 2p core levels and valence band spectra obtained on our hematite films are shown in Figure 3. By convention, we considered the O 1s line at 530.1 eV. We observe similar electronic structures for all thicknesses and for 0 and 2 at. %

Figure 3. (a) Fe 2p core level XPS spectra for undoped films of 11 nm (red dotted line) and 45 nm (cyan solid line) and for 2 at. % Ti-doped films of 12 nm (black dashed line) and 50 nm (green solid line). (b) Ti 2p core level XPS spectra for 2 at. % Ti-doped films of 21 nm (black dashed line), 29 nm (blue dotted line), and 50 nm (green solid line) and (c) valence band XPS spectra (same legend as (a)).

Figure 4. Photocurrent density vs voltage curves: (a) undoped hematite films of 11 nm (red empty box), 20 nm (black filled circle), and 40 nm (blue x) and (b) 2 at. % Ti-doped hematite films of 21 nm (black open diamond), 29 nm (blue down triangle), and 50 nm (green plus). 3010

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much higher than values reported in the literature for samples with identical doping levels12,14,17,18 where increases of a factor between 2 and 10 only are reported. Because of the different nature of the layers, these differences can be attributed to morphology and structure differences. The undoped samples have poor photoelectrochemical properties (Figure 4a). Below 0.2 V vs Ag/AgCl the photocurrent is negative, which indicates a cathode behavior, whereas for larger anodic potentials a positive photocurrent standing for an anode behavior appears. For the Ti-doped samples (Figure 4b), only an anodic photocurrent appears at positive potential vs Ag/AgCl. Let us now consider the photocurrent at 0.6 V vs Ag/AgCl as a function of the film thickness (Figure 5). The dependences of

dealing with the variation of water splitting properties with the electrode thickness. This is mainly due to deposition technique limitations making very difficult the production of samples having well-defined thicknesses without changing additional important parameters like the crystallographic structure or the surface morphology. Well-controlled techniques like atomic layer deposition (ALD),12,29 pulsed laser deposition,33 or molecular beam epitaxy (present case) are desirable to make a study on variable thickness. On one hand, Zandi et al.12 observed on hematite films made by ALD an increase of the photocurrent followed by a saturation when the film thickness increases, which can be explained as follows: when the thickness is larger than the depletion region width, no additional photocurrent is produced since any excess of charge carriers created in the bulk recombines. On the other hand, Dotan et al.33 showed that with a reflective substrate like platinum the variation of photocurrent with thickness is due to resonant light trapping in the film and to minority carriers (holes) recombination. Interestingly, our measurements on Tidoped hematite compare well with their results obtained on Tidoped hematite films with a platinum substrate where the photocurrent was simulated with a minority carrier (hole) collection length of 20 nm. Multiple reflections and hole recombinations explain thus well the photocurrent vs thickness curve shape in our Ti-doped samples but not in undoped ones where the amplitude of the photocurrent variation is larger than 75%. In this case, for thicknesses higher than the space charge layer width, we conclude that the photocurrent is limited by the diffusion length of both holes and electrons. In a simple picture this limitation can be illustrated as follows: in undoped samples the holes photogenerated far from the surface recombine before reaching the surface, and the electrons created near the surface cannot reach the substrate because of recombination with holes.21 This limitation by electron diffusion length is maybe due to the well-known weak electron mobility in αFe2O3(0001).6 Our study of photocurrent vs thickness allows unraveling the influence of Ti doping in photoelectrochemical properties; indeed, we clearly show that the improvement observed in Ti-doped hematite is mainly due to an increase of both electrons and holes diffusion lengths. IPCE and PSIPCE measurements at 0.6 V vs Ag/AgCl for samples of different thickness are reported on Figure 6. The very weak photoresponse of our undoped samples unfortunately leads to unfavorable signal-to-noise ratios and made impossible IPCE measurements for these films; therefore, only PSIPCEs are shown for these films, as a function of the

Figure 5. Photocurrent density at 0.6 V vs Ag/AgCl as a function of the film thickness for undoped (black open circle, multiplied by 75) and 2 at. % Ti-doped (green filled diamond) hematite. Continuous lines are only an eye guide.

photocurrent with thickness for undoped and doped films are different. For the undoped films we can observe a well-defined sharp photocurrent maximum around 20 nm, whereas for the Ti-doped films the curve is smoother and the photocurrent maximum shifted to a thickness of 30 nm. The shape of the curves at low thicknesses can be easily understood. First, before the peak, the photocurrent increases when the thickness increases, thanks to the progressive photon absorption. The much higher photocurrent for Ti-doped hematite compared to undoped hematite is due to a higher diffusion length of holes (lower recombination rate) in the depletion layer upon Ti doping. Second, at higher thicknesses, the behavior of photocurrent is highly dependent on doping level: the variation of the photocurrent is larger in undoped films (>75%) than in doped ones (∼50%) when the thickness becomes larger than 20 and 30 nm, respectively. There are very few studies12,29,33

Figure 6. (a) IPCE curves at 0.6 V vs Ag/AgCl for 2 at. % Ti-doped hematite films of 21 nm (black open triangle), 29 nm (blue down triangle), and 50 nm (green plus). (b) Phase-sensitive IPCE (PSIPCE) curves at 0.6 V vs Ag/AgCl for 2 at. % Ti-doped hematite films of 29 nm (blue up triangle) and 50 nm (green asterisk) and for a 45 nm thick undoped hematite film (red open circle). Experimental data were acquired with 10 nm steps. 3011

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in contact with the same electrolyte (same redox level), which is expressed by the relation

wavelength. Interestingly, in Figure 6a, we see that whatever the wavelength a ca. 30 nm thickness for Ti-doped hematite films gives the best IPCE values, which is consistent with the white light response (photocurrent−voltage curve Figure 4b and photocurrent−thickness curve Figure 5). Our efficiency values are of the same order of magnitude as values reported in the literature for the same kind of samples (thicknesses and doping level).7,15,16 However, other studies reported higher IPCE values12,14,18 which may be linked to different morphology and/ or structure. PSIPCE is not a quantitative measure of efficiency; however, by comparing Figure 6a and Figure 6b one can see that these curves can be used for the study of wavelength dependency. We observe the same behavior for all samples (all thicknesses and the two doping levels): a decrease of efficiency between 400 and 600 nm and a cutoff at 600 nm (equivalent to a band gap of ca. 2.1 eV). These results show that the band gap does not vary either with the thickness or with the doping level. Electrochemical impedance spectroscopy was performed on two different samples of the same thickness (15 nm undoped and 2 at. % Ti-doped hematite thin films). For each sample, the charge capacitance was determined by a fit using eq 2 at each potential. The results are reported in Figure 7, in the form of

V fb0%Ti − Vfb2%Ti = E F2%Ti − E F0%Ti

(4)

Therefore, we can conclude that the position of the Fermi level does not vary with the Ti doping level used in the present work. As expected, Ti acts as a negative charge carrier provider in the hematite layer, increasing the sample free carrier concentration. As a matter of fact, in a covalent bond scheme Ti atoms can furnish up to 4 electrons to become Ti4+, whereas Fe atoms only give 3 electrons to become Fe3+ ions (see XPS results). An increase in carrier concentration increases the conductivity of the samples since the conductivity σ of hematite can be expressed by σ = ND × e × μ where μ is the electron mobility in the semiconductor. RHEED patterns showed a constant crystallographic structure of hematite upon Ti-doping, excluding any morphology or crystallographic structure modifications. XPS, IPCE, and EIS measurements showed that an ca. 2 at. % Ti doping of hematite, within the experimental uncertainties, does not modify the electronic structure (energy bands, gap, Fermi level). Thus it is legitimate to consider that the improved photoelectrochemical properties upon Ti doping in our samples are not linked to any crystallographic or electronic structure modification but are induced by intrinsic properties changes only. The incorporation of titanium ions in iron sites dramatically increases the carriers’ diffusion lengths in hematite and also increases the conductivity of the samples. This lowers the recombination losses in Ti-doped hematite. The carriers’ behavior in undoped and Ti-doped hematite films (thin and thick) is summarized and illustrated in Scheme 1.

IV. CONCLUSIONS We have studied the growth and the photoelectrochemical properties of single crystalline Ti-doped and undoped hematite layers of various thicknesses (5−50 nm range) in the framework of solar-assisted water splitting. We used AO-MBE that allows the deposition of very high quality flat epitaxial

Figure 7. Mott−Schottky plots for 15 nm undoped (black open circle) and 2 at. % Ti-doped (green plus) hematite thin films. Blue dashed lines indicate the linear fit of each curve according to eq 3.

Mott−Schottky plots (A2/C2 = f(V)). From these data we can deduce the flat-band potential Vfb (the difference between the semiconductor Fermi level and the electrolyte redox level) and the carrier concentration (ND) by a fit of the Mott−Schottky curve using eq 3. The fits are indicated on Figure 7 (blue dashed lines), and the values are reported in Table 1.

Scheme 1. Illustration with a Band Diagram of the Carriers’ Behavior in Undoped and Ti-Doped Films (Thin and Thick)a

Table 1. Flat Band Potentials and Carrier Concentrations for 15 nm Undoped and 2 at. % Ti-Doped Hematite α-Fe2O3 film

Vfb (V vs Ag/AgCl)a

ND (cm−3)a

undoped 2 at. % Ti-doped

−0.41 ± 0.06 −0.45 ± 0.05

(2.3 ± 0.2) × 1018 (1.4 ± 0.1) × 1019

Values obtained from the fit of Mott−Schottky plots (Figure 7) using eq 3.

a

Interestingly, the Ti doping induces an important increase of the carrier concentration (1 order of magnitude) without a significant change in the flat band potential (within the experimental error bar). This result is consistent with values reported in the literature7,12,18 where it is shown that a variation of the flat band potential occurs only for doping levels higher than 5 at. %.15,16 The difference of flat band potentials corresponds to the difference of Fermi levels when materials are

a

Precisions of symbol meanings: Simple arrows and cycling arrows stand for charge movement (current) and charge recombination, respectively. Carriers concerned are electrons or holes when the line is, respectively, red or blue. Line thicknesses convey the strength of the corresponding phenomena.

3012

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layers. Undoped and Ti-doped hematite layers were found to have the same crystallographic structure and to adopt similar Stransky−Krastanov growth mechanisms. XPS measurements, wavelength dependency, and EIS revealed undistinguishable electronic structures (band edge position, gap, Fermi level) for both types of layers (undoped and Ti-doped) and whatever the thickness, whereas photoelectrochemical measurements showed a strong improvement for low Ti-doped films. Comparative photoelectrochemical measurements allowed evidencing the intrinsic pertinent parameters that contribute to the water splitting performance improvement and showed that Ti doping increases (i) the carrier concentration (higher conductivity) and (ii) the carriers’ diffusion lengths as revealed by photocurrent−thickness curves.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Phone: +33 (0)1 69 08 94 04. Author Contributions

The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was funded by the CEA project DSM-Energie Hémaphoto. We acknowledge C. Mocuta and D. Thiaudière for making possible access to DIFFABS beamline for X-Ray reflectivity measurements.

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REFERENCES

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