Single Crystalline Hematite Films for Solar Water Splitting: Ti-Dopingand Thickness EffectsMaxime Rioult, Helene 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 assistedmolecular beam epitaxy (AO-MBE) on Pt(111) substrates in the framework ofhydrogen harvesting from sunlight-induced water splitting. Such singlecrystalline samples are suitable model systems to study thickness and dopingeffects on the photoelectrochemical properties; we demonstrate that they alsoallow disentangling intrinsic transport properties from mingled overallproperties due to the usually unknown contributions from morphology orcrystalline structure defects. From their photoelectrochemical characteristics(I(V) curves, incident photon to current efficiency measurements, andelectrochemical impedance spectroscopy), we evidence the existence of anoptimum layer thickness, which is higher for Ti-doped samples (30 nm) as compared to undoped ones (20 nm). Our resultssuggest that this effect is due to an increase of the carrier concentration combined with higher carriers’ diffusion lengths in thedoped samples stressing intrinsic modifications of the hematite layer upon titanium doping that cannot be accounted for bysimple structural or electronic structure changes.

I. INTRODUCTIONSolving the concomitant worldwide increasing energy con-sumption problem and the need for greenhouse gas reductionto avoid or limit climate changes imposes a higher fraction ofrenewable energies in the total energy production. Sustainablegrowth cannot anymore be dissociated from fostering novelrenewable energies that become more and more a first-classissue. Within this framework, increasing attention is paid to thesunlight-assisted water splitting as a clean method of hydrogenproduction. As a matter of fact, hydrogen is an energy carrier ofchoice which does not lead to any greenhouse gas production.Although the idea of producing hydrogen using water splittingassisted by solar light is very seductive, it remains unfortunatelyalso very tricky, and many materials science issues have to besolved. During the process, electron−hole pairs are generatedin a semiconductor electrode, upon solar light absorption, andare subsequently used to promote the oxido-reduction reactionsof water leading to oxygen production at the photoanode andhydrogen production at the photocathode.1,2

Since the pioneering discovery of water-photoassistedelectrolysis using semiconducting TiO2 in 1972 by Fujushimaand Honda,1 several materials were investigated as photo-anodes3 where water oxidation occurs (2OH− + 2h+ → (1/2)O2 + H2O). Hematite, i.e., the α-Fe2O3 iron oxide, is one ofthe most promising materials regarding its characteristics. It isan n-type semiconductor with a quasi-ideal band gap (∼2.2 eV)for solar water splitting applications. Indeed this material is ableto absorb ca. 40% of the solar light spectrum, and its theoreticalsolar-to-hydrogen conversion yield reaches 13%.3 It is abundanton earth and very stable in aqueous environments, which makesit a serious candidate in the framework of green energyproduction.4 Moreover, the valence band edge of hematite is

located below the H2O/O2 redox potential which favors thewater oxidation reaction.5 Unfortunately, hematite has not onlyadvantages since it has been demonstrated to have weaktransport properties6,7 (low conductivity and low carrierlifetime) and poor surface kinetics.8 Also, the conductionband edge of α-Fe2O3 is not well positioned with respect to thepotential of the water reduction reaction (2H2O + 2e+ → H2 +2OH−), thus an external bias is necessary to promote watersplitting.9

Various strategies were proposed to overcome hematitedrawbacks3,4,10 including (i) doping with different elementssuch as Ti, Si, or Sn7,11−21 which can improve electricalproperties and also modify the electronic structure; (ii)nanostructuration (e.g., nanowires, mesoporous materials, ...)to compensate small carrier diffusion lengths;5,22,23 (iii) morecomplex semiconductor heterostructures24,25 to optimize lightabsorption 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 toimprove hematite properties and has been the subject ofnumerous studies in recent years.7,11−18 However, the improve-ment of the electrical conductivity induced by Ti doping hasnot been discussed in detail with respect to possible changes ofthe crystalline structure. As a matter of fact, previous studiesconcerning polycrystalline films or nanostructures include ahigh density of grain boundaries which may dominate theelectric conduction; they can either limit conduction (if theconduction occurs perpendicularly to them) or increase

Received: January 10, 2014Revised: January 20, 2014Published: January 23, 2014

Article

pubs.acs.org/JPCC

© 2014 American Chemical Society 3007 dx.doi.org/10.1021/jp500290j | J. Phys. Chem. C 2014, 118, 3007−3014

conduction (if the conduction occurs along them). Theimportance of grain boundaries has been stressed by Glasscocket al.,7 for example, who proposed for polycrystallinemagnetron-sputtered samples that the passivation of hematitegrain boundaries and reduction of the crystalline quality bytitanium doping could explain the photoelectrochemicalproperty enhancement upon doping, whereas silicon dopingshowed only limited improvements supposedly because of adecrease of grain boundary size. For hematite films grown byhydrothermal methods, Miao et al.13 reported that Ti dopinginduces a change in nanostructure size and shape and Deng etal.14 a modification of the type of hematite nanostructure andcrystallographic orientation in agreement with the works ofTang et al.15,16 and Hahn and Mullins.18 However, because ofthe polycrystalline nature of these films, the contributions dueto structural, morphological, electronic, and crystallographiceffects and the intrinsic effects due to Ti doping are necessarilyentangled. It is thus not yet clear if the modifications of water-splitting properties upon Ti doping are linked to the change ofcrystallographic structure or not. The investigation of singlecrystalline samples is an elegant way to overcome this problem.For undoped hematite layers, the detailed behavior of thephotoelectrochemical properties as a function of the layerthickness, especially for well-crystallized layers, has also beenless considered. Previous studies used hematite photoanodesobtained by deposition techniques such as hydrothermalgrowth,13,14 chemical vapor deposition,5,8,9,26 or electro-deposition,20 patterned or not. In these approaches, it is verydifficult to distinguish which parameter, including preparationissues, morphology, or crystal quality, etc., influences thephotoelectrochemical properties. Moreover, other works usingtechniques suitable for model sample elaboration like atomiclayer deposition12,19,28,29 or sputtering,7,15,16 focused mainly ondopants and nanostructuration influence or on surface kineticsproperties.Therefore, the present work aims at studying thin epitaxial

undoped and Ti-doped hematite films of different thicknessesas photoanodes. To investigate single crystalline electrodes asmodel samples, the layers were deposited on Pt(111) substratesusing atomic oxygen plasma assisted molecular beam epitaxy(AO-MBE).6,30,31 These single crystalline samples allow us toinvestigate the modification of the intrinsic properties ofhematite by Ti doping, without any change of crystallographicstructure and/or morphology.

II. EXPERIMENTAL SECTIONSample Preparation. The hematite layers were deposited

on single crystalline Pt(111) substrates using AO-MBE, atechnique that makes possible the deposition of singlecrystalline layers of controlled morphology and thickness.High purity metals (99.99% grade) were evaporated fromdedicated Knudsen cells in the presence of an atomic oxygenplasma (350 W power) to obtain well-defined oxides undergood vacuum conditions (i.e., 10−7 mbar working pressure,10−10 mbar base pressure). During the deposition, the sampleswere rotated continuously around their normal to ensure ahomogeneous deposit and radiatively heated at a temperatureof ca. 450 °C with a filament located behind the sample. Theoxide deposition rate was about 0.15 nm/min. The thicknessesof our films were in the 5−50 nm range, and the Ti-dopinglevel was either 0 at. % (undoped) or ca. 2 at. % (doped).Crystallographic Structure and Chemistry. In situ

reflection high-energy electron diffraction (RHEED) patterns

were observed and acquired during film growth to monitor thecrystal quality and structure of the samples. In situ XPS spectrawere systematically recorded just after deposition to determinethe stoichiometry and the electronic structure of the films.More precisely, we recorded Fe 2p, Ti 2p, and O 1s core levelsand the valence band region using Al Kα radiation. The dopinglevel 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 intensitiescorrected by the corresponding Scofield factors (crosssections). For all samples studied, the measured doping levelis 2 at. % (±0.3 at. %). Additionally, X-ray reflectivitymeasurements were performed for thickness determination.

Photoelectrochemistry. The photoelectrochemical re-sponse of our films was studied using a three-electrode cell.All electrochemical measurements were performed at roomtemperature using a NaOH 0.1 M (pH = 13) solution aselectrolyte, a platinum wire as counter electrode, and a Ag/AgCl electrode for the potential reference (VAg/AgCl = +0.197 Vvs SHE). The sample was mounted as an anode (workingelectrode) using a dedicated sample holder that allows thecontact only between the hematite surface and the electrolyte.The illumination source was a Newport 1000 W Xe Arc lampwith an infrared filter. The incident light flux is around 100mW/cm2 (measured with a Newport 1918-R Power Meter).Potential control and current acquisition between the threeelectrodes were done using a Princeton Applied Research(PAR) 263A potentiostat controlled by computer. Forphotocurrent−voltage curves, I(V), the potential was sweptfrom 0 to +0.8 V vs Ag/AgCl at a speed of 50 mV/s. Thephotocurrent is defined as the difference between the currentrecorded 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 Vvs Ag/AgCl. The IPCE can be calculated with18

λλ

λ

λ= ×hc J

ePIPCE( )

( )

( )(%)ph

(1)

where Jph(λ) is the photocurrent density, P(λ) the incidentpower density, λ the wavelength of the incident light, h thePlanck 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 thesemeasurements, we used a Cornerstone 130 model 74004monochromator (Newport). The wavelength was variedbetween 200 and 1000 nm, with 10 nm steps. Transientvariations of the photocurrent are expected to be due to (i)voltage application or to (ii) chopped incident light. Tominimize the transient signals due to the voltage application,the potential was held constant during 200 s prior to IPCEacquisition. Since the values of the photocurrents in somehematite films were very low (<μA range), we used a phase-sensitive detection technique which significantly increases thesignal-to-noise ratio. For this purpose we used a PAR 5210lock-in amplifier and an optical chopper at a referencefrequency of 20 Hz. The IPCE signal obtained from such atechnique (noted PSIPCE for phase-sensitive incident photonto current efficiency) includes also the transient photocurrentdue to the light intensity modulation by the chopper at 20 Hz.Therefore, these PSIPCE values cannot be directly comparedwith IPCE values. These measurements were performed on allfilms.

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Electrochemical impedance spectroscopy (EIS) was per-formed on two samples (undoped and 2 at. % Ti-doped 15 nmhematite films) in dark conditions. A dynamic signal of 20 mVamplitude with frequencies from 100 kHz down to 0.5 Hz wassuperposed to a DC bias ranging from −0.4 to 0.8 V vs Ag/AgCl, with 0.1 V steps, using the potentiostat and the lock-inamplifier to generate the alternative signal V(ω) and to measurethe generated current I(ω). The cell impedance was then Z =V(ω)/I(ω). Impedance data for frequencies for which theHelmholtz layer and the surface states can be neglected7 (>1kHz) were used to determine the capacitance C of the spacecharge layer at the surface−electrolyte interface. For this, wemodeled our system with an RS−RC circuit of equivalentimpedance Z given by

ω= + + − −Z R i C R( )S1 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 carrierconcentration (ND)

18

ε ε= × − −

⎛⎝⎜

⎞⎠⎟

AC eN

V Vk T

e22

2D 0 r

fbB

(3)

where A is the illuminated surface area, ε0 the dielectricconstant of vacuum (8.85 × 10−12 F/m), εr the dielectricconstant of hematite taken equal to 32,7,12 V the potentialapplied to the electrochemical cell, and kB the Boltzmannconstant (1.38 × 10−23 J/K). The flat band potentialcorresponds to the difference between the electrochemicalpotentials of the semiconductor (Fermi level) and of theelectrolyte (redox level).

III. RESULTS AND DISCUSSIONStructure and Chemistry. The growth of undoped

hematite thin films by AO-MBE on Pt(111) single crystalshas already been studied in detail in previous works.30,31 Herewe describe the growth of Ti-doped hematite films on Pt(111)single crystals. RHEED patterns as well as correspondingsurface reciprocal lattices in such films at different depositiontimes (different thicknesses) are presented on Figure 1,assuming the same epitaxial relationship between Pt(111) andγ-Fe2O3 or α-Fe2O3 as in ref 31 (i.e., cube on cube for γ-Fe2O3(111)/Pt(111) and √3 × √3R30° for α-Fe2O3(0001)/Pt(111)). We observe that the thin film growth is similar to thecase of undoped samples31 and follows a Stranski−Krastanovgrowth mechanism (layer-plus-islands). First, a 2D Ti-doped γ-Fe2O3 layer grows, evidenced by extra straight lines on thecorresponding RHEED patterns (step b in Figure 1). A phasetransition from γ-Fe2O3(111) to α-Fe2O3(0001) occursbetween 3 and 4.5 nm thickness (steps c and d), where the γand α phases coexist. After the γ → α transition, an epitaxialgrowth of Ti-doped hematite (α-Fe2O3) with some islands andextra roughness is observed (evidenced by the presence of spotsalong 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-planelattice parameter. Since different phases appear during thegrowth and since they give rise to lattices belonging to differentspace groups, it is however not straightforward to determinelattice parameter evolution. Therefore, we prefer, with the sameapproach as Barbier et al.,31 to represent a generic lattice

parameter p which is calculated using the RHEED patternsobtained over the particular diffraction direction D1 (madeexplicit 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 in-plane lattice parameter of γ-Fe2O3(111), assuming acube-on-cube epitaxial relationship with Pt(111));

• aα‑Fe2O3/√3 in the case of α-Fe2O3 (where aα‑Fe2O3 is thein-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 samelattice parameter for all crystallographic phases, its evolution asa function of the film thickness remains relevant for the overalllattice relaxation and is presented in Figure 2. The latticeparameter relaxation during the epitaxial growth process withrespect to increasing film thickness is consistent with previousworks.31 The parameter p converges to a value of 2.92 ± 0.02Å, which corresponds to an in-plane lattice parameter of 5.06 ±0.02 Å for hematite. This is very close, or even equal within theexperimental error bar, to the bulk hematite lattice parametervalue (5.04 Å). Figure 2 shows also that the film relaxes withinthe first ca. 20 nm.The stoichiometry and the electronic structure of our layers

were determined by means of in situ XPS measurements. Fe 2p

Figure 1. (Left panel) RHEED patterns of a 50 nm thick 2 at. % Ti-doped α-Fe2O3 film grown by AO-MBE at different deposition times/thicknesses (steps a to e) over the two lowest index surface diffractiondirections called, for simplicity, D1 and D2. (Right panel) Thecorresponding surface reciprocal lattices are represented, makingexplicit the diffraction directions. (Blue box), (red circle), and (greencircle) 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-Fe2O3and α-Fe2O3, respectively.

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

doping levels. More precisely, we measured binding energies of711.2 eV for the Fe 2p3/2 core level and of 719.2 eV for thetypical Fe3+ shakeup satellite (dash-dotted gray lines on Figure3a). This confirms the Fe3+ ionic state of iron in all films andthe absence of detected Fe2+ species, excluding polaronhopping (Fe3+ + Ti4+ → Fe2+ + Ti3+)32 in our Ti-dopedhematite samples. For Ti-doped samples, the Ti 2p3/2 lineshows a narrow single feature at a binding energy of 458 eV(dash-dotted gray line on Figure 3b). This value is between thetabulated binding energies of Ti4+ and Ti3+ in TiO2 (458.8 and457.1 eV, respectively). The crystallographic structure of 40 nmfilms has been studied in detail by EXAFS,11 and we haveshown that Ti substitutes Fe in the hematite structure with aslight deformation of the oxygen octahedron. Since we do notobserve any change of RHEED patterns or XPS spectra withthickness, we can reasonably assume that this substitution isalso present in the thinner samples with a lower doping levelstudied here. Therefore, as in references 6 and 11, we canconclude that this binding energy corresponds to Ti4+ speciesincluded in a hematite host matrix. Valence band spectra(Figure 3c) did not show changes over thickness or Ti dopinglevel variations. We observe neither a shift of the valence bandlike in ref 11 nor states in the gap like in ref 6. In a previousstudy11 we have shown that the valence band shifts towardhigher binding energies (0.2 eV shift) when hematite is dopedwith higher Ti doping level (>5 at. %). The present resultsshow that a lower doping level, at least up to 2 at. %, does notinduce a detectable modification in the electronic structure.

Photoelectrochemical Properties. Figure 4 presents thephotocurrent density vs voltage curves for undoped and Ti-doped hematite films of various thicknesses. We see that thephotocurrent density obtained for Ti-doped samples is higherthan for undoped ones by 2 orders of magnitude, which isconsistent with previous results.7,15,16 Let us note that theincrease of photocurrent with Ti doping in the present study is

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

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-dopedfilms 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 solidline) and (c) valence band XPS spectra (same legend as (a)).

Figure 4. Photocurrent density vs voltage curves: (a) undopedhematite 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 (greenplus).

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much higher than values reported in the literature for sampleswith identical doping levels12,14,17,18 where increases of a factorbetween 2 and 10 only are reported. Because of the differentnature of the layers, these differences can be attributed tomorphology and structure differences. The undoped sampleshave poor photoelectrochemical properties (Figure 4a). Below0.2 V vs Ag/AgCl the photocurrent is negative, which indicatesa cathode behavior, whereas for larger anodic potentials apositive photocurrent standing for an anode behavior appears.For the Ti-doped samples (Figure 4b), only an anodicphotocurrent 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

photocurrent with thickness for undoped and doped films aredifferent. For the undoped films we can observe a well-definedsharp photocurrent maximum around 20 nm, whereas for theTi-doped films the curve is smoother and the photocurrentmaximum shifted to a thickness of 30 nm. The shape of thecurves at low thicknesses can be easily understood. First, beforethe peak, the photocurrent increases when the thicknessincreases, thanks to the progressive photon absorption. Themuch higher photocurrent for Ti-doped hematite compared toundoped hematite is due to a higher diffusion length of holes(lower recombination rate) in the depletion layer upon Tidoping. Second, at higher thicknesses, the behavior ofphotocurrent is highly dependent on doping level: the variationof the photocurrent is larger in undoped films (>75%) than indoped ones (∼50%) when the thickness becomes larger than20 and 30 nm, respectively. There are very few studies12,29,33

dealing with the variation of water splitting properties with theelectrode thickness. This is mainly due to deposition techniquelimitations making very difficult the production of sampleshaving well-defined thicknesses without changing additionalimportant parameters like the crystallographic structure or thesurface morphology. Well-controlled techniques like atomiclayer deposition (ALD),12,29 pulsed laser deposition,33 ormolecular beam epitaxy (present case) are desirable to makea study on variable thickness. On one hand, Zandi et al.12

observed on hematite films made by ALD an increase of thephotocurrent followed by a saturation when the film thicknessincreases, which can be explained as follows: when thethickness is larger than the depletion region width, noadditional photocurrent is produced since any excess of chargecarriers created in the bulk recombines. On the other hand,Dotan et al.33 showed that with a reflective substrate likeplatinum the variation of photocurrent with thickness is due toresonant light trapping in the film and to minority carriers(holes) recombination. Interestingly, our measurements on Ti-doped hematite compare well with their results obtained on Ti-doped hematite films with a platinum substrate where thephotocurrent was simulated with a minority carrier (hole)collection length of 20 nm. Multiple reflections and holerecombinations explain thus well the photocurrent vs thicknesscurve shape in our Ti-doped samples but not in undoped oneswhere the amplitude of the photocurrent variation is larger than75%. In this case, for thicknesses higher than the space chargelayer width, we conclude that the photocurrent is limited by thediffusion length of both holes and electrons. In a simple picturethis limitation can be illustrated as follows: in undoped samplesthe holes photogenerated far from the surface recombine beforereaching the surface, and the electrons created near the surfacecannot reach the substrate because of recombination withholes.21 This limitation by electron diffusion length is maybedue to the well-known weak electron mobility in α-Fe2O3(0001).

6 Our study of photocurrent vs thickness allowsunraveling the influence of Ti doping in photoelectrochemicalproperties; indeed, we clearly show that the improvementobserved in Ti-doped hematite is mainly due to an increase ofboth 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. Thevery weak photoresponse of our undoped samples unfortu-nately leads to unfavorable signal-to-noise ratios and madeimpossible IPCE measurements for these films; therefore, onlyPSIPCEs are shown for these films, as a function of the

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

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), and50 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.

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wavelength. Interestingly, in Figure 6a, we see that whatever thewavelength a ca. 30 nm thickness for Ti-doped hematite filmsgives the best IPCE values, which is consistent with the whitelight response (photocurrent−voltage curve Figure 4b andphotocurrent−thickness curve Figure 5). Our efficiency valuesare of the same order of magnitude as values reported in theliterature for the same kind of samples (thicknesses and dopinglevel).7,15,16 However, other studies reported higher IPCEvalues12,14,18 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 seethat these curves can be used for the study of wavelengthdependency. We observe the same behavior for all samples (allthicknesses and the two doping levels): a decrease of efficiencybetween 400 and 600 nm and a cutoff at 600 nm (equivalent toa band gap of ca. 2.1 eV). These results show that the band gapdoes 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 undopedand 2 at. % Ti-doped hematite thin films). For each sample, thecharge capacitance was determined by a fit using eq 2 at eachpotential. The results are reported in Figure 7, in the form of

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

Interestingly, the Ti doping induces an important increase ofthe carrier concentration (1 order of magnitude) without asignificant change in the flat band potential (within theexperimental error bar). This result is consistent with valuesreported in the literature7,12,18 where it is shown that a variationof the flat band potential occurs only for doping levels higherthan 5 at. %.15,16 The difference of flat band potentialscorresponds to the difference of Fermi levels when materials are

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

− = −V V E Efb0%Ti

fb2%Ti

F2%Ti

F0%Ti

(4)

Therefore, we can conclude that the position of the Fermi leveldoes not vary with the Ti doping level used in the present work.As expected, Ti acts as a negative charge carrier provider in thehematite layer, increasing the sample free carrier concentration.As a matter of fact, in a covalent bond scheme Ti atoms canfurnish up to 4 electrons to become Ti4+, whereas Fe atomsonly give 3 electrons to become Fe3+ ions (see XPS results). Anincrease in carrier concentration increases the conductivity ofthe samples since the conductivity σ of hematite can beexpressed by σ = ND × e × μ where μ is the electron mobility inthe semiconductor.RHEED patterns showed a constant crystallographic

structure of hematite upon Ti-doping, excluding any morphol-ogy or crystallographic structure modifications. XPS, IPCE, andEIS measurements showed that an ca. 2 at. % Ti doping ofhematite, within the experimental uncertainties, does notmodify the electronic structure (energy bands, gap, Fermilevel). Thus it is legitimate to consider that the improvedphotoelectrochemical properties upon Ti doping in oursamples are not linked to any crystallographic or electronicstructure modification but are induced by intrinsic propertieschanges only. The incorporation of titanium ions in iron sitesdramatically increases the carriers’ diffusion lengths in hematiteand also increases the conductivity of the samples. This lowersthe recombination losses in Ti-doped hematite. The carriers’behavior in undoped and Ti-doped hematite films (thin andthick) is summarized and illustrated in Scheme 1.

IV. CONCLUSIONSWe have studied the growth and the photoelectrochemicalproperties of single crystalline Ti-doped and undoped hematitelayers of various thicknesses (5−50 nm range) in theframework of solar-assisted water splitting. We used AO-MBEthat 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 dashedlines indicate the linear fit of each curve according to eq 3.

Table 1. Flat Band Potentials and Carrier Concentrations for15 nm Undoped and 2 at. % Ti-Doped Hematite

α-Fe2O3 film Vfb (V vs Ag/AgCl)a ND (cm−3)a

undoped −0.41 ± 0.06 (2.3 ± 0.2) × 1018

2 at. % Ti-doped −0.45 ± 0.05 (1.4 ± 0.1) × 1019

aValues obtained from the fit of Mott−Schottky plots (Figure 7) usingeq 3.

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

aPrecisions of symbol meanings: Simple arrows and cycling arrowsstand 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 thecorresponding phenomena.

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layers. Undoped and Ti-doped hematite layers were found tohave the same crystallographic structure and to adopt similarStransky−Krastanov growth mechanisms. XPS measurements,wavelength dependency, and EIS revealed undistinguishableelectronic structures (band edge position, gap, Fermi level) forboth types of layers (undoped and Ti-doped) and whatever thethickness, whereas photoelectrochemical measurementsshowed a strong improvement for low Ti-doped films.Comparative photoelectrochemical measurements allowedevidencing the intrinsic pertinent parameters that contributeto the water splitting performance improvement and showedthat Ti doping increases (i) the carrier concentration (higherconductivity) and (ii) the carriers’ diffusion lengths as revealedby photocurrent−thickness curves.

■ AUTHOR INFORMATION

Corresponding Author*E-mail: helene.magnan@cea.fr. Phone: +33 (0)1 69 08 94 04.

Author ContributionsThe manuscript was written through contributions of allauthors. All authors have given approval to the final version ofthe manuscript.

NotesThe authors declare no competing financial interest.

■ ACKNOWLEDGMENTS

This work was funded by the CEA project DSM-EnergieHemaphoto. We acknowledge C. Mocuta and D. Thiaudiere formaking possible access to DIFFABS beamline for X-Rayreflectivity measurements.

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