Charge-Transfer through Ultrathin Film TiO2 on n-Si(111

Article pubs.acs.org/JPCC

Charge-Transfer through Ultrathin Film TiO2 on n‑Si(111) Photoelectrodes: Experimental and Theoretical Investigation of Electric Field-Enhanced Transport with a Nonaqueous Redox Couple Hark Jin Kim,#,† Kara L. Kearney,#,‡,§ Luc H. Le,† Zachary J. Haber,‡ Angus A. Rockett,‡,§ and Michael J. Rose*,† †

Department of Chemistry, The University of Texas at Austin, Welch Hall, 105 E 24th Street, Austin, Texas 78712, United States Department of Materials Science and Engineering, University of Illinois, 1304 West Green Street, Urbana, Illinois 61801, United States § International Institute for Carbon Neutral Energy Research (WPI-I2CNER), Kyushu University, 744 Moto-oka, Nishi-ku, Fukuoka 819-0395, Japan ‡

S Supporting Information *

ABSTRACT: Ultrathin film amorphous (a-TiO2) and anatase crystalline (c-TiO2) titanium dioxide were investigated as corrosion passivation layers on n-type Si(111). Varying thicknesses of TiO2 (5−60 Å) were deposited on nSi(111)−CH3 substrates by atomic layer deposition (ALD) at 150 and 240 °C, thus yielding a-TiO2 and c-TiO2, respectively. The phase and morphology of the TiO2 films were determined using a combination of XRD, Raman spectroscopy, XPS, AFM and SEM. The electronic properties of a-TiO2 and c-TiO2 films were compared using 4-point sheet resistance and electrochemical impedance spectroscopy. Substrates functionalized with c-TiO2 exhibited higher conductivity, lower charge transfer resistance, and comparable anticorrosion behavior to a-TiO2. The photoelectrochemical response of n-Si(111)−CH3|TiO2 electrodes as a function of TiO2 thickness was characterized using cyclic voltammetry with ferrocene in acetonitrile. A thickness of 20 or 40 Å was required to block charge transport through a-TiO2 and c-TiO2, respectively. Lastly, the charge transport behaviors of both the amorphous and crystalline n-Si(111)−CH3|TiO2 constructs were enhanced via the deposition of platinum nanoparticles (ALD) on the TiO2 layer. Using a solid-state drift diffusion simulation package (wxAMPS), a theoretical basis for the charge-transport behavior was developed. The experimental thickness-dependence results were used as a basis of comparison to determine the charge transfer mechanism across the n-Si(111)−CH3|TiO2 electrodes. The simulations suggest that the charges conduct via field-assisted thermionic emission across the Si(111)−CH3|TiO2 interface, utilizing a defect band that is consistent with the “leaky dielectric” attribute of TiO2 films. In addition, the simulations suggest that the defects present in a-TiO2 behave as trap states, while the defects present in c-TiO2 behave as recombination centers; this is derived from the observed difference in photoelectrochemical behavior between the two films. p+-Si, n-GaP, and n-p+-GaAs.8 The np+-Si|a-TiO2+44 nm|Niisland photoanode revealed a stable operation of water oxidation for over 100 h in 1 M KOH(aq), while allowing the passage of a very high (>1 A cm−2) current density, which was ascribed to its “leaky dielectric” behavior. A very recent report by this same group explored the spectroscopic (XPS, UPS) and electrochemical properties of Si(100)|TiO2 heterojunctions.9 In related work, Bard and co-workers reported the surface modification of a BiVO4 photoanode by electrodeposition of a-TiO2 to avoid undesired back-reduction in the water oxidation reaction.10 By introducing 80−120 nm of a-

1. INTRODUCTION Although silicon represents an attractive light-absorber for photoelectrochemical (PEC) applications because of its low material cost and high earth abundance, surface degradation in semiconductor|liquid junctions remains an important obstacle toward functional and stable devices. To prevent degradation, a number of organic and inorganic materials have been introduced as passivating layers on Si substrates: alkylfunctionalization,1 SiO2,2 Al2O3,3 TiO2,4 etc. Among them, TiO2 has attracted much attention owing to its electronic conductivity and high chemical stability at extreme pH values.5−7 Lewis et al. reported the use of amorphous TiO2 (a-TiO2) deposited by atomic layer deposition (ALD) as a passivating layer on a set of water oxidation photoanodes including n-Si, n© 2016 American Chemical Society

Received: August 13, 2016 Revised: October 17, 2016 Published: November 3, 2016 25697

DOI: 10.1021/acs.jpcc.6b08096 J. Phys. Chem. C 2016, 120, 25697−25708

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The Journal of Physical Chemistry C

on unterminated Si(111) surface sites, which could establish alternative charge transfer routes other than through the discrete interface. In the present work, we have studied the effect of a-TiO2 and c-TiO2 thickness on charge transfer across n-Si(111)−CH3| TiO2(5−60 Å) photoanodes with and without Pt-NPs using cyclic voltammetry. The experimental results are used as a basis for simulations of the charge transfer phenomena using wxAMPS, wherein the effects of both (i) tunneling and (ii) electric field-assisted thermionic emission are considered. The simulations imply that the charges conduct via a field-assisted thermionic emission model through a defect band in the “leakydielectric” TiO2. The effect of Pt-NPs in enhancing the thermionic emission via electric field effects is also investigated.

TiO2, the resulting device exhibited stability for at least 12 h, and the onset potential of water oxidation was shifted cathodically by ∼0.5 V. Additionally, the photocurrent of water oxidation was increased by roughly 6-fold. It was postulated that the observed high electronic conductivity of a-TiO2 was due to Ti3+ ions in the lattice arising from oxygen vacancies, thus resulting in a conductive electronic defect.10 Ager and Javey et al. reported the role of TiO2 deposited by ALD on p-InP photocathodes and provided a framework for understanding the important electronic differences based on the identity of the titanium(IV) ALD precursor. These researchers compared the PEC performance of p-InP|TiO2 as a function of the ALD-TiO2 precursor.11 Specifically, 10 nm of c-TiO2 was grown on p-InP at 250 °C using tetrakis(dimethylamido)titanium (TDMAT) and titanium tetraisopropoxide (TTIP) precursors. The InP|TiO2 photoelectrodes with N-TiO2 (grown from TDMAT) had a 0.23 V lower photovoltage compared to O-TiO2 (grown from TTIP). This difference was attributed to Ti3+ defects in the lattice, which were revealed by X-ray photoelectron spectroscopy to be present in N-TiO2 but not in O-TiO2. These authors postulated that the Ti3+ defects allow for hole conduction through TiO2, which enhances electron recombination at the surface and diminishes the onset potential for water reduction.11 Fermin et al. and Gooding et al. have thoroughly investigated the role of metal NPs in recovering the electrochemical kinetics through self-assembled monolayers (SAM) deposited on metal electrodes.12−14 Specifically, it was observed that the deposition of SAMs hinders electron transfer from the metal electrode to the redox couple; however, fast kinetics is restored when Pt or Au NPs are deposited on top of the metal|SAMs electrode. Chazalviel and Allongue suggested that this behavior could be understood by the fact that electron transfer from the underlying metal electrode to the metal NPs is orders of magnitude faster than electron transfer from the metal electrode to redox species in solution.15 Additionally, Fermin and co-workers have shown that the extent of enhancement in the NP’s charge transfer mediation is related to the overlap in the density of states of the nanoparticles and redox species in solution.16 Bard et al. extended this work by replacing the SAMs with TiO2 resulting in a Pt(UME)|TiO 2|Pt-NP assembly.17 In these electrodes, the TiO2 was sufficiently thin to facilitate electron conduction through the TiO2 via tunneling. Similarly, it was postulated that the Pt−NPs improved the electrochemical kinetics across the TiO2 layer because tunneling from the underlaying metal electrode to metal NPs is much more probable than tunneling from the metal electrode to redox species in solution.17 Si(111)−CH3 electrodes are ideal substrates for studying charge transfer across the Si|TiO2 interface due to the welldefined chemical and structural properties of the Si(111)−CH3 surface.18 The Si(111)−CH3 surface is achieved using a twostep chlorination/alkylation reaction on a Si(111)-H substrate.19 Unlike the Si(111)-H surface, Si(111)−CH3 resists oxidation in both air and under aqueous conditions due to the 100% termination of Si(111) surface sites with atop Si−C bonds. Thus, the Si(111)−CH3|TiO2 construct presents a different type of interface as compared with TiO2 films grown on alternate Si crystal faces (with hydride or oxide termination). While alternate crystal faces result in gradient junctions [Si → SiOx → TiO2], the methyl-terminated Si(111) results in a discrete, step junction [Si-CH3 → TiO2] with no intermingling SiOx layer. As such, TiO2 growth does not initiate

2. EXPERIMENTAL SECTION 2.1. Methyl-Terminated Si(111) Substrate Preparation. Si(111)−CH3 substrates were prepared analogous to the method in our previous report.20 A single-side-polished, phosphorus-doped, Cz grown n-Si(111) wafer (350 ± 25 μm thick, Virginia Semiconductor, Inc.) with 1−10 Ω·cm resistivity was used for all experiments. First, the Si wafer was subsequently cleaned using sonication for 10 min in acetone, ethanol, and water. Immediately following sonication the Si wafer was immersed in a Piranha solution of 3:1 (v/v) ratio of H2SO4 (96.6%, Fisher Scientific) and H2O2(aq) (30%, Fisher Scientific) at 90 °C for 20 min. After the Piranha wash, the wafer was rinsed thoroughly with high purity water (18 MΩ). Next, the Si wafer was etched in a HF(aq) solution (semiconductor grade, Transene Company, Inc.) for 20 s at room temperature. The atomically flat Si(111)-H surface was then prepared by immersing the wafer in a degassed NH4F(aq) solution (semiconductor grade, Transene Company, Inc.) for 20 min. After washing with high purity water and drying under a stream of N2 gas, the wafer was transferred to a glovebox under inert atmosphere (N2). The Si(111)-H surface was chlorinated in a solution of chlorobenzene (99.8%, SigmaAldrich) saturated with PCl5 (metal basis, 99.998%, Alfa Aesar) at 90 °C for 45 min; to initiate the chlorination, one grain of benzoyl peroxide (reagent grade, 97%, Sigma-Aldrich) was added prior to heating. After the chlorination reaction, the Si(111)−Cl was washed sequentially with chlorobenzene and tetrahydrofuran (THF). Finally, the Si wafer was methylated with a commercially available Grignard reagent prepared in a 2:1 (v/v) ratio of THF and methylmagnesium chloride (3.0 M in THF) at 60 °C for 30 min. 2.2. Atomic Layer Deposition of TiO2 and Pt. The TiO2 films and Pt-NPs were deposited on Si(111)−CH3 wafers by ALD using a Savannah S100 apparatus (Cambridge Nanotechnology, Inc.) analogous to our previous report.12 The deposition of TiO 2 is a result of cycling tetrakis(dimethylamido)titanium (TDMAT, Sigma-Aldrich) with chromatography-grade H2O. The reaction chamber was heated at 150 and 240 °C to generate a-TiO2 and c-TiO2, respectively. The Ti precursor container was held at 75 °C. Each pulse length of the Ti precursor and H2O was 0.1 and 0.015 s, respectively. Purging time between pulses by N2 gas was 20 and 5 s for a-TiO2 and c-TiO2, respectively. Pt-NPs were deposited on TiO2 passivated Si(111)−CH3 wafers using the same Savannah S100 apparatus. The deposition of Pt results from the reaction of trimethyl(methylcyclopentadienyl)-platinum ([(MeCp)Pt(Me)3], Strem) with high-purity O2 (99.999%). The temperature of the substrate was maintained at 240 °C 25698


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voltammetry (CV) was performed to evaluate the electrochemical properties of n-Si(111)|CH3|TiO2 electrodes at varying thicknesses of TiO2 both with and without Pt using a 2 mM ferrocene (98%, Acros Organics) in 0.3 M LiClO4 acetonitrile solution. Before performing each CV scan, the electrolyte solution was degassed by N2 bubbling for 30 min to remove dioxygen from the solution. The reported potentials are shown as a raw potential scale (vs Ag) and the redox potential of Fc0/+ was +0.284 V vs Ag. All PEC measurements were performed at 100 mV s−1. A detailed experimental setup is available in our previous report.25

while the Pt precursor container and gas lines were held at 70 and 150 °C, respectively, to provide a steady state flux of Pt precursor from source to reactor. Each pulse length of the Pt precursor and the oxygen source was 1 and 0.015 s, respectively, and the purging time by N2 gas between pulses was 5 s. 2.3. Physical Characterization. X-ray diffraction (XRD) patterns of the TiO2 deposited on n-Si(111)−CH3 were obtained using a Bruker D8 Advance diffractometer at 40 kV and 30 mA with Cu Kα radiation (λ = 1.5406 Å). Raman spectra were acquired in ambient atmosphere using a Witec Micro-Raman spectrometer Alpha 300 with a 488 nm laser. Atomic force microscopy (AFM) images were obtained using a Park Scientific instrument in contact mode. To reduce friction between the cantilever and the surface, dodecylamine was vapor transferred to the surfaces before analysis by placing the samples in a chamber with solid dodecylamine for 10 min. A 500 nm × 500 nm analysis window was used with a force constant of zero. WSxM was used to analyze the images acquired.21 The surface morphology of the TiO2 films was determined by a Quanta 650 SEM (FEI). The SEMs images were obtained at 30 kV and a 10 mm working distance with a secondary electron detector. XPS measurements were performed using an X-ray photoelectron spectrometer (Kratos Axis Ultra) with a monochromated Al Kα X-ray source (hv = 1486.5 eV). The photoelectron takeoff angle was normal to the sample surface and 45° with respect to the X-ray beam. The obtained spectra were analyzed by Casa XPS software (version 2.3.15, Casa Software Ltd.) using 70% Gaussian and 30% Lorentzian function after subtraction of Shirley background. From the XPS signal, the TiO2 film thickness (tox) was calculated using the following equation: tox = λ sin(θ ) ln(R /R ∞ + 1)

3. THEORETICAL METHODS A drift-diffusion solid-state device software package, wxAMPS, was utilized in this work to model carrier transport through nSi(111)−CH3|TiO2 photoelectrodes.26−28 wxAMPS calculates the concentration of charge carriers throughout the device as a function of position and applied potential by solving the three coupled nonlinear differential equations using the method of finite differences: −

dψ ⎞ d ⎛ ⎜ε(x ) ⎟ = qρ dx ⎝ dx ⎠

(2)

1 ⎛ dJn ⎞ ⎜ ⎟ = −Gop(x) + R(x) q ⎝ dx ⎠

(3)

1 ⎛ dJp ⎞ ⎜⎜ ⎟⎟ = Gop(x) − R(x) q ⎝ dx ⎠

(4)

where ψ is the electrostatic potential referenced to the position of the local vacuum level, ε is the permittivity of the material, q is the magnitude of the charge of an electron, and ρ is the summed charge density (free carriers, trapped carriers, and ionized impurities). A more detailed description of the software is included in a previous report.29 In the present work, the organic layer and electrolyte contact are modeled using a previously established method for adapting wxAMPS to simulate carrier transport across functionalized Si(111) photoelectrodes in contact with electrolyte.29 The semiconductor−liquid junction is treated as a solid-state Schottky junction with the contact potential set as the formal potential of the redox couple. The dipole introduced at the surface by the organic group is modeled as a varied electron affinity in the bulk of the semiconductor.29 This simulation model was previously validated on the basis of an experimental study by Lewis et al. in which the Voc values for Si(111)-H versus Si(111)−CH3 using both p- and n-type electrodes were measured in contact with a series of redox couples in CH3CN containing 1.0 M LiClO4.30 The electronic parameters used as inputs into wxAMPS were developed as follows. The bulk silicon values are those tabulated by the Ioffee Institute: permittivity = 11.7, energy gap (Eg) = 1.12 eV, conduction band density of states (Nc) = 3.2 × 1019 cm−3, valence band density of states (Nv) = 1.8 × 1019 cm−3, electron mobility (μn) = 1400 cm2 V−1 s−1, and hole mobility (μp) = 450 cm2 V−1 s−1.31 The electron affinity of the bulk Si layer was set to 3.8 eV, which is the value established in previous work for a CH3-terminated Si(111) surface.29 The electrolyte contact potential was set as the formal energy of the ferrocene redox couple. The electronic parameters for TiO2 were based on values from Chorkendorff et al. for ALD-grown

(1)

where λ is the attenuation length of the photoelectrons, θ is the photoelectron takeoff angle (90° in this work), R is the ratio Iox/ISi in the sample, and R∞ is the ratio Iox/ISi in a sample where the oxide film is greater than 10 nm thick.22 It has been shown that the measured attenuation length and R∞ values for thermally grown silicon oxide films can be applied to TiO2 deposited via ALD in ultrathin oxide films (less than 4.5 nm) as such thin films only elicit minimal elastic photoelectron scattering. Thus, from the literature, 2.96 nm and 0.91 were used for λ and R∞, respectively.23,24 2.4. Electronic and Photoelectrochemical Characterization. The sheet resistance was obtained using a four-point probe conductivity measurement system comprising a Lucas Laboratories SP4 four-point probe head combined with a Keithley 2400 source meter. Electrochemical impedance spectroscopy (EIS) was carried out using an Interface 1000 (Gamry Instruments) at open circuit potential (Voc) under dark conditions with an AC amplitude of 10 mV over a frequency range of 0.1−105 Hz. The PEC properties of n-Si(111)−CH3| TiO2 electrodes were investigated using a WaveNow (Pine Research Instrumentation) potentiostat. A three-electrode configuration was utilized, consisting of the n-Si(111)−CH3| TiO2 substrate as the working electrode, a Pt-wire (99.95%, Strem) counter electrode, and an Ag-wire quasi-reference electrode purchased from CHI112 (CH Instruments). The light source used to simulate sunlight at 100 mW cm−2 was a 150 W Xe lamp (Newport, Co., USA) equipped with an AM1.5G solar filter (model# 81094, Newport, Co., USA). Cyclic 25699


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The Journal of Physical Chemistry C TiO2: permittivity = 55, Nc = 1 × 1020 cm−3, Nv = 1 × 1020 cm−3, and Nd = 1 × 1017 cm−3.32 The electron affinity was set to 4.1 eV and μn was approximated as 0.001 cm2 V−1 s−1. The value of Eg and μp were adjusted simulation parameters to reflect the charge transfer through the TiO2, which is discussed in detail in the results section. A detailed description of the defect model incorporated into wxAMPS is included in a previous work.29 Charge carriers may conduct across a potential-energy barrier by a thermally excited charge emission process known as thermionic emission. The thermionic emission current density (JTE) is described by eq 533 ⎛ qΦ ⎞ JTE = AT 2 exp⎜ − b ⎟ ⎝ kT ⎠

of anatase TiO2 (JCPDS #75-1537). Although the peak intensity is not high due to the ultrathin thickness (∼15 nm), it indicates the crystalline phase of TiO2 (c-TiO2). Furthermore, the Raman spectrum of the TiO2 film deposited at 240 °C (Supporting Information (SI), Figure S1) exhibits a peak at 145 cm−1, which is also characteristic of anatase TiO2; the Raman spectrum of the amorphous film (ALD at 150 °C; SI Figure S1) exhibits no discernible feature in the same region. To understand the phase formation and transition properties depending on temperature, a-TiO2 was heated post-deposition at 240 °C for 10 min (the c-TiO2 deposition temperature and time). The post-deposition heat treatment did not convert aTiO2 to the crystalline phase at 240 °C as indicated by the lack of the anatase diffraction peak at 25.62°. Also, if a phase change to c-TiO2 had occurred, the TiO2 should have exhibited a photocurrent similar to the authentic Si(111)−CH3|c-TiO2 sample in the cyclic voltammetry experiment (vide infra). As shown in Figure S2, the post-heating of n-Si(111)|CH3|a-TiO260 (∼21 Å) at 240 °C shows blocked hole transfer behavior similar to the authentic a-TiO2 substrate; the cyclic voltammogram shows no anodic peak indicative of ferrocene oxidation. This result is not surprising because the phase transition of TiO2 from amorphous to anatase is generally known to be above 500 °C.34 The point demonstrated by this result is that anatase c-TiO2 can be produced at a considerably lower temperature by ALD.11 This is essential in this work because Si substrates modified with organic moieties generally decompose at temperatures above ∼250 °C. Characterization of the films by topography and friction AFM reveals different morphologies for a-TiO2 (SI Figure S3a,b) and c-TiO2 (SI Figure S3c,d). The a-TiO2 images reveal a surface consisting of a micro-textured morphology that appears malleable during consecutive AFM scans. The a-TiO2 “mound-like” features presumably result from dominant nucleation points during ALD. On the other hand, the cTiO2 film exhibits a lower drag coefficient and is composed of a uniform sheet of nanocrystallites. Furthermore, the SEM images of c-TiO2 (SI Figure S4a) and a-TiO2 (SI Figure S4b) reveal a greater extent of particle aggregation for a-TiO2 versus the nanocrystalline sheet observed for c-TiO2. To precisely determine the thickness of the TiO2 films, we performed X-ray photoelectron spectroscopy (XPS) using the Si 2p underlayer signal as a benchmark. Figure 2 shows the calculated thicknesses of TiO2 deposited on n-Si(111)−CH3 as a function of ALD cycles. The raw XPS spectra used for

(5)

where A* is the effective Richardson’s constant, T is temperature, q is the charge of an electron, k is Boltzmann’s constant, and Φb is the effective Schottky barrier height. The presence of a strong electric field at the potential-energy barrier may lower the effective Schottky barrier height by an amount ΔΦb resulting in a field-enhanced thermionic emission current (JFETE) as described by ⎛ q(Φ b − ΔΦ b) ⎞ JFETE = AT 2 exp⎜ − ⎟ ⎝ ⎠ kT

ΔΦ b =

q3F 4πεo

(6)

(7)

where F is the magnitude of the electric field and εo is the vacuum permittivity.

4. RESULTS AND DISCUSSION 4.1. Physical Characterization of TiO2 Films. We determined the phase of TiO2 (amorphous vs crystalline) based on the deposition temperature. Figure 1a shows the XRD

Figure 1. (a) XRD patterns for 400-cycle TiO2 on n-Si(111)|CH3 substrates depending on heating temperature and (b) a magnified inset: (A) bare n-Si(111)|CH3, (B) n-Si(111)|CH3|a-TiO2 (ALD at 150 °C), (C) additional post heat treatment of sample B at 240 °C for 10 min, (D) n-Si(111)−CH3|c-TiO2 (ALD at 240 °C).

patterns for 400 cycles of ALD-grown TiO2 on n-Si(111)−CH3 and Figure 1b shows a magnified inset. Compared to a bare nSi(111)−CH3 sample, the TiO2 film deposited at 150 °C exhibits a single diffraction peak at ∼28.5°, which is indicative of the (111)-orientation of Si (JCPDS #71-4631). In contrast, the TiO2 film deposited at 240 °C exhibits an additional diffraction peak at 25.62°, which is identified as the (101) peak

Figure 2. Thickness of TiO2 films grown on n-Si(111)−CH3 versus number of ALD cycles as derived from the X-ray photoelectron spectra of the Si 2p and Ti 2p region. 25700


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or ordering properties). In general, a large number of defects or disordered Ti and O atoms will scatter mobile carriers resulting in a lower conductivity/higher resistivity.36 Therefore, the high crystallinity (well-ordered Ti and O atoms) in c-TiO2 compared to a-TiO2 affords an advantage in carrier transport. As mentioned in the introduction, the electronic defect Ti3+ serves as a defect site for recombination, but paradoxically it is also a driving force for the hole conductivity of TiO2. To investigate the possibility of Ti3+ defects in c-TiO2 and a-TiO2 on n-Si(111)−CH3, XPS was performed and the assigned XPS results are shown in Figure 4a and Figure 4b, respectively. The Ti 2p peaks at 458.7 and 464.4 eV are assigned to 2p3/2 and 2p1/2 binding energies, respectively, and the distance between two peaks was 5.7 eV corresponding to Ti4+.32 In addition, a noticeable XPS peak with lower binding energy of 457.1 eV is assigned to Ti3+ caused by a Ti−N bond; therefore, this feature is also indicative of the extent of N doping.37 As shown in Figure 4 panels a and b, both a-TiO2 and c-TiO2 show the presence of Ti3+, which appears inevitable due to the Ti−N bonds inherent in the ALD precursor (TDMAT).11 The 2p3/2 proportion of Ti3+/Ti4+ for c-TiO2 was 3.93 ± 0.40%, which is only slightly higher than 3.06 ± 0.26% of a-TiO2. Such a small difference (0.87%) indicates that the extent of N-doping does not vary between a-TiO2 and c-TiO2 and therefore is not responsible for the observed difference in conductivity observed between the two phases. For more quantitative insight into the interfacial properties at the semiconductor|TiO2|liquid junction, electrochemical impedance spectroscopy (EIS) was utilized. Figure 5 panels a−d shows the EIS Bode plot measured in the dark with an applied open circuit voltage for bare n-Si(111)−CH3 wafer, n-Si(111)− CH3|a-TiO2, and n-Si(111)−CH3|c-TiO2. In the Bode modulus of Figure 5a, both TiO2 functionalized wafers exhibit higher |Z| (ohm) values, indicating higher corrosion resistance.38 Additionally, in the Bode phase plot of Figure 5b two different time constant shapes are shown: one for the bare n-Si(111)−CH3 substrate and one for the n-Si(111)−CH3 substrate covered with a-TiO2 or c-TiO2. The increasing time constant is correlated with the formation of an additional layer that is resistive and capacitive to interfacial reactions. The observed phenomenon can be analyzed quantitatively by fitting with an equivalent circuit model. SI Figure S6 shows the equivalent circuit for an ultrathin metal oxide passivated electrode.39,40 The equivalent circuit consists of several elements:40−42 a series resistance (Rs) for the electrolyte and/ or electrical lead, a resistance (R1) and capacitance (CPE1) for the substrate/oxide interface, a charge transfer resistance (R2) for the redox reaction, an electrical double layer capacitance

calculating the thickness are shown in SI Figure S5a−d. The thickness for both a-TiO2 and c-TiO2 was found to linearly increase with the number of ALD cycles (20−80 cycles). However, the x-intercept of the trend-line is not zero, thus indicating the nucleation lag phase of the ALD process (20 cycles) can be defined by d(Å) = 4.15(x=20) + 0.41(x − 20) and d(Å) = 2.86(x=20) + 0.45(x − 20) [d = thickness, x = number of ALD cycles] for a-TiO2 and c-TiO2, respectively. 4.2. Electronic Characterization of TiO2 Films. The sheet resistance of a-TiO2 and c-TiO2 as a function of film thickness was determined using a four-point probe measurement system. Figure 3 shows that sheet resistance increases as a

Figure 3. (a) Sheet resistance data for n-Si(111)−CH3|c-TiO2 (blue triangles) and n-Si(111)−CH3|a-TiO2 (red circles) obtained by the four-point probe measurement.

function of the number of TiO2 ALD cycles for both a-TiO2 and c-TiO2; however, c-TiO2 showed a diminishing trend of rising resistance while the resistance of a-TiO2 continuously rose. These results indicate that c-TiO2 has a higher conductivity/lower resistance compared to a-TiO2. From these results we conclude that the conductivity difference between a-TiO2 and c-TiO2 on Si(111)−CH3 is due to a change in the crystallinity of the films (i.e., atomic defects and/

Figure 4. X-ray photoelectron spectra in the Ti 2p region for (a) n-Si(111)−CH3|c-TiO2-60 and (b) n-Si(111)−CH3|a-TiO2-60. Note the nearly identical extent of Ti3+ (teal) defect density in each sample. 25701


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Figure 5. Bode (a) modulus and (b) phase plots for n-Si(111)−CH3, n-Si(111)−CH3|a-TiO2 ,and n-Si(111)−CH3|c-TiO2-60 cycles of both a-TiO2 and c-TiO2 were selected for EIS analysis. Plots (c) and (d) show the comparisons between experimental results and equivalent circuit fitting for the bare n-Si(111)−CH3 sample.

Table 1. Obtained EIS Parameters Fitted by Equivalent Circuita bare a-TiO2 c-TiO2

Rs (Ω)

R1 (Ω)

CPE1 (F)

n1

R2 (Ω)

CPE2 (F)

n2

W

225.4 249.6 251.0

2.78 × 103 2.52 × 104 2.91 × 104

8.10 × 10−8 5.93 × 10−8 3.83 × 10−8

0.867 0.864 0.914

1.80 × 104 1.39 × 106 6.97 × 105

2.75 × 10−8 2.81 × 10−7 4.49 × 10−7

1.00 0.916 0.826

7.38 × 10−6 8.06 × 10−7 9.81 × 10−7

a

Notation: Rs, series resistance of the electrolyte; R1, resistance of substrate|oxide interface; CPE1, capacitance of substrate|oxide interface; n, exponent value of CPE; R2, charge transfer resistance of redox reaction; CPE2, electrical double layer capacitance; W, Warburg impedance.

Figure 6. XPS results for chemical stability test of (a) n-Si(111)−CH3|c-TiO2-60 and (b) n-Si(111)−CH3|a-TiO2-60 samples. Within each panel, the black and red lines represent the as-prepared sample (pre-EIS) and used sample (post-EIS) for electrochemical measurement, respectively. The corresponding XPS results for bare n-Si(111)−CH3 can be found in ref 17.

CPE exponent related to the angle ((1 − n) × 90°) that is evaluated from the real Z axis in the Nyquist plot. For n = 1 the CPE describes an ideal capacitor, and for n ≈ 0.5 the CPE represents a Warburg impedance with a diffusion character.43 The EIS parameters from the fit with the applied equivalent circuit model (SI Figure S6) are tabulated in Table 1. In fact, the general bare electrode shows a good fit with a Randles circuit − R(CPE,RW) as evidenced in Figure 5c. However, as shown in Figure 5d the bare n-Si(111)−CH3 electrode reveals a closer match with the R(CPE,R)(CPE,RW) circuit, which is the model used for TiO2 covered wafers. Fitting the bare n-Si(111)| CH3 wafer with a R(CPE,R)(CPE,RW) circuit provides a n2

(CPE2) for the electrolyte, and a Warburg impedance (W) for the electrolyte diffusion. Because most materials do not exhibit ideal capacitor behavior, the original capacitance C is not suitable to fit the experimental result. Thus, the constant phase element (CPE) was introduced to describe the nonideal capacitive behavior. The CPE can be defined by the following equation:43 ZCPE =

1 Q (jω)n

(8)

where Q is the CPE constant, ω is the angular frequency (ω = 2πf, f is the frequency), j is an imaginary number, and n is the 25702


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Figure 7. Cyclic voltammograms for (a) n-Si(111)−CH3|c-TiO2 and (b) n-Si(111)−CH3|a-TiO2 obtained with 2 mM ferrocene (Fc0/+) in 0.3 M LiClO4 acetonitrile solution. All data were collected at 100 mV s−1 under AM-1.5G 1-sun illumination.

Figure 8. Plots of the maximum oxidation current density (JOX) as a function of TiO2 thickness for (a) n-Si(111)−CH3|a-TiO2-(x), (b) n-Si(111)− CH3|a-TiO2-(x)|Pt, (c) n-Si(111)−CH3|c-TiO2-(x) and (d) n-Si(111)−CH3|c-TiO2-(x)|Pt samples. The Pt was deposited by ALD and JOX was measured using cyclic voltammetry with 2 mM ferrocene in CH3CN containing 0.3 M LiClO4.

respectively) indicates a better film quality and a negligible presence of structural defects at the substrate|oxide interface.43 The higher R2 (6.97 × 105 Ω) of c-TiO2 compared to a-TiO2 (1.39 × 106 Ω) indicates a faster charge transfer rate between the electrode and the redox molecule in the electrolyte. In addition, XPS was utilized to observe the passivation properties of a-TiO2 and c-TiO2 thin films. Figure 6a,b shows the XPS comparison of oxidized surfaces before and after EIS for c-TiO2 and a-TiO2, respectively. First, following the EIS measurement, the bare n-Si(111)−CH3 sample exhibits a broad feature at 101−104 eV, which is assigned as the Si4+ signal (SiO2, at 103 ± 0.3 eV)46,47 and indicative of SiOx formation (22.7%).25 Such oxidation is due to trace dioxygen and/or water in the electrolyte, or through exposure to ambient air during the experiment. In contrast, on both a-TiO2 and c-TiO2 passivated n-Si(111)−CH3 samples, only a negligible broad peak is observed. The atomic fractions of Si-oxide/Si-total quantified by using XPS signals1for n-Si(111)−CH3|c-TiO260 and n-Si(111)−CH3|a-TiO2-60 were relatively lower, at 12.7 ± 0.03% and 12.1 ± 0.11% (respectively). The n-Si(111)−CH3| c-TiO2-60 sample shows slightly higher SiO2 peak proportion

value (= 1), which represents a high quality capacitive layer. This ideal parallel-plate capacitive property is likely caused by the surface methyl group. The methyl monolayer at ∼100% coverage on the n-Si(111) surface25 allows it to act as a capacitive and resistive layer at the substrate|electrolyte interface. Fits of the R(CPE,R)(CPE,RW) circuit model on Si(111)−CH3 wafers coated with a-TiO2 and c-TiO2 are shown in SI Figure S7a,b. As shown in Table 1, the R1 and R2 of TiO2-passivated samples show higher resistance (∼10-fold) compared to the substrate without TiO2. The ALD-grown TiO2 acts as a resistive layer due to high resistance as shown in Figure 5a, thus the electron transfer rate is decreased compared to bare nSi(111)−CH3 at the substrate|oxide interface and across the substrate|oxide|electrolyte multiface. In general, it is known that the resistance (R1 + R2) is inversely related to the surface corrosion rate,44,45 so an increased R1 + R2 indicates an enhancement in passivation of a n-Si(111)−CH 3 |TiO 2 compared to a bare n-Si(111)−CH3 wafer. The slightly higher R1 (2.91 × 104 Ω) and n1 parameters (0.914) of c-TiO2 compared to those of a-TiO2 (2.52 × 104 Ω and 0.864, 25703


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Figure 9. (a) Schematic of the two separate hole conduction mechanisms used in the simulation (wxAMPS) of the oxidation current versus TiO2 thickness behavior of n-Si(111)−CH3|TiO2 photoelectrodes. Pathway (1) is direct tunneling across the metal oxide passivation layer while pathway (2) is defect-mediated transport through a defect band. (b) Simulated oxidation current versus TiO2 thickness for both tunneling (red dashed line) and defect-mediated transport (blue solid line). The band gap of TiO2 was set to 3.2 and 1.7 eV for tunneling and defect-mediated transport, respectively.

than the n-Si(111)−CH3|a-TiO2-60 sample, which is consistent with the R1 + R2 trend that represents the corrosion rate in the EIS measurement (shown in Table 1). Overall, these results demonstrate that an ultrathin film of c-TiO2 or a-TiO2 improves the stability of the n-Si(111)−CH3 surface during ambient storage and device operation (with Fc/Fc+ in MeCN). 4.3. Photoelectrochemical Characterization. Under illumination, the n-Si(111) wafer absorbs photons, which excites electrons to the conduction band and leaves holes behind. The holes generated at the valence band of n-Si(111) were scavenged by the oxidation of ferrocene. As shown in Figure 7 panels a and b, the anodic photocurrent JOX of nSi(111)−CH3|x-TiO2 wafer systematically decreased as the number of TiO2 ALD cycles increased. A c-TiO2 layer as thick as 27 Å (80 cycles c-TiO2) still allowed a significant amount of current, ∼0.4 mA cm−2. However, beyond this thickness the TiO2 layer significantly blocked the hole transfer, resulting in negligible current. In contrast, the a-TiO2 layer allowed only ∼0.26 mA cm−2 current passing through at a diminished thickness of 16.5 Å (50 cycles a-TiO2) as shown in Figure 7b. We investigated the ability of Pt nanoparticles (NPs) to improve hole conduction through n-Si(111)−CH 3 |TiO 2 electrodes, which could be described as metal|insulator| semiconductor (MIS) devices. To introduce Pt-NPs on the TiO2-modified n-Si semiconductor substrate, 20 cycles of ALDPt were carried out resulting in sparsely dispersed Pt nanoparticles.25 The maximum oxidation current (JOX) was determined using CV and 2 mM ferrocene in a 0.3 M LiClO4 MeCN solution. Figure 8 panels a−d show the JOX current for n-Si(111)−CH3|TiO2 and n-Si(111)−CH3|TiO2|Pt at varying thicknesses of both c-TiO2 and a-TiO2. The original cyclic voltammograms of n-Si(111)−CH3|TiO2|Pt are provided as SI Figure S8 and resemble those of n-Si(111)−CH3|TiO2 shown in Figure 7a,b. As shown in Figure 8a,b, the photocurrent of n-Si(111)− CH3|a-TiO2 diminished at 20.6 Å (60 cycles of a-TiO2) and this hole conductive maximum thickness (dmax) was increased by the addition of Pt nanoparticles up to 45.2 Å (130 cycles of a-TiO2), analogous to the charge transfer extension in our previous report using Al2O3 blocking layers.25 In comparison and shown in Figure 8c,d, the unplatinized crystalline sample of n-Si(111)−CH3|c-TiO2 exhibits a dmax of 36.3 Å (100 cycles of c-TiO2), which was increased to 50.5 Å (140 cycles of c-TiO2) with the addition of Pt nanoparticles. This device also exhibited

a more gradual decrease in current over a broader range of thicknesses (20−55 Å) compared to the sharp decline on the platinized a-TiO2 sample (30−40 Å). [Note that the curves through the data in Figure 8 are sigmoidal fits to the experimental results, and not the simulation results. Additionally, all electrodes discussed in this work show negligible electrochemical response under dark conditions.] 4.4. Simulation of Charge Transfer Phenomena. Theoretically, the large valence band offset between TiO2 and Si should prevent holes from passing through the TiO2. However, tunneling has been suggested as a possible hole conduction mechanism through large-gap metal oxides and demonstrated experimentally through TiO2 films grown by ALD on Pt UMEs.47,48 A simulation of tunneling-based hole transfer through the TiO2 (pathway 1 in Figure 9a) was attempted in wxAMPS using an intraband tunneling mechanism.49 As shown by the red dashed curve in Figure 9b, the tunneling simulation results in an immediately and exponentially decreasing current versus TiO2 thickness: no current through the device is predicted for thicknesses greater than ∼1 nm of TiO2 and only negligible current is predicted even at 0.5 nm (5 Å). These results are inconsistent with both the a-TiO2 and c-TiO2 experimental data sets. In a recent work, a-TiO2 films grown by ALD on silicon substrates4,8 exhibited hole transport through films much greater than 2 nm, which is essentially the thickness limit for effective hole tunneling.33 Campet et al. proposed a qualitative “leaky dielectric” mechanism to explain such behavior, in which holes transport through a defect band in the band gap of the wide gap material (pathway 2 in Figure 9a).50 A leaky dielectric hole transfer mechanism through the TiO2 layer was modeled in wxAMPS by lowering the band gap of the TiO2 to approximate a low energy defect band. The simulated current versus TiO2 thickness behavior with the Eg set to 1.7 eV is shown in Figure 9b. The simulated behavior using a leaky dielectric model is consistent with both the a-TiO2 and c-TiO2 data sets in which a thickness-independent charge transfer regime is followed by thickness-dependent transfer. In this model, the carriers thermionically emit over the potential barrier at the Si-CH3|TiO2 interface, followed by drifting across the defect band and to the redox couple. This is in agreement with a study by Lewis et al. in which evidence of a semiconductor/defect-band heterojunction model of a n-Si/ 25704


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Figure 10. (a) Simulated thickness dependence for the short-circuit current density (JSC for hole conduction) for n-Si(111)−CH3|TiO2 photoelectrodes with varying defect band energies and a constant hole mobility of μ = 0.0001 V cm−2 s−1. (b) Simulated behavior for the JSC for hole conduction for n-Si(111)−CH3|TiO2 photoelectrodes with varying hole mobilities (μp) and a fixed defect band energy (Eg = 1.7 eV).

0.0001 cm2 V−1 s−1 (Figure 10a). Despite modulating the energy of the defect band by 0.1 eV, the simulated dmax shifts by only 1 nm. More notably, the biggest influence of the defect band energy is on the shape of the current versus TiO2 behavior. A larger barrier height at the Si-CH3|TiO2 interface (Eg = 1.7 eV) provides a smoother, more gradual decline in current versus thickness than a smaller barrier height (Eg = 1.65 or 1.6 eV). However, it is not clear if such a “shape” effect can be experimentally resolved based on the precision of the present data. The effect of hole mobility in the TiO2 layer was also modeled by adjusting the input value for the hole mobility (μp) with the Eg set constant at 1.7 eV (this value most closely simulates the a-TiO2 behavior). Indeed, the hole mobility parameter results in a pronounced influence on dmax as shown in Figure 10b. Specifically, a hole mobility of 0.0001, 0.001, and 0.01 cm2 V−1 s−1 results in a dmax of ∼1, 2, and 3 nm, respectively. Because the drift current is proportional to the product of mobility and electric field (μ × E), the extent of current through the TiO2 at steady-state increases with an increase in hole mobility. Therefore, it takes a greater thickness of TiO2 to induce the change in behavior from carriers generated in the bulk Si limiting current (case 1) to fieldassisted thermionic emission across the Si−CH3|TiO2 potentialenergy barrier limiting current (case 2). The effect of recombination in the TiO2 layer was modeled by adjusting the input value for the concentration of donor defects at midgap with the Eg set constant at 1.7 eV and hole mobility set constant at 0.01 cm2 V−1 s−1 as shown in Figure 10c. The defects were simulated as both traps (σp = 1 × 10−10, σn = 1 × 10−20) and recombination centers (σp = 1 × 10−10, σn = 1 × 10−10), where σp and σn are the capture cross sectional areas for holes and electrons, respectively. However, this difference had a negligible effect on the current output. The value of the defect concentration results in the same

“leaky” TiO2 interface was observed using electrical, photoelectrochemical, and photoelectron spectroscopic techniques.9 The shape of the simulated oxidation current versus TiO2 thickness behavior can be explained by separating the current into three regimes that represent the three parts of the sigmoidal curve. Case (1) represents the saturated oxidation current at the thickness-independent maximum value in very thin layers of TiO2. In this regime, the flux of carriers from the bulk Si to the Si−CH3 surface limits the current, but not the potential barrier at the Si−CH3|TiO2 interface. Under the conditions of case (1), the potential barrier at the Si-CH3|TiO2 interface is insufficient at limiting the current across the electrode given the conductivity of the relatively thin TiO2 and the electric field promoting hole transfer to the electrolyte. Case (2) represents an exponentially decreasing oxidation current versus TiO2 thickness. In this regime the extent of fieldassisted thermionic emission over the potential barrier at the Si−CH3|TiO2 interface limits the current (as described by eqs 6 and 7). Under the conditions of case (2), the increasing thickness of the TiO2 progressively diminishes the field to the point that holes cannot emit across the Si-CH3|TiO2 interface. Case (3) represents the situation of negligible oxidation current. In this regime, a relatively wide barrier, coupled with a low electric field disables the holes from emitting across the Si-CH3|TiO2 interface. The maximum hole conduction thickness (dmax) is that which provides the minimum magnitude of electric field at the Si−CH3|TiO2 interface necessary to initiate thermionic emission. To suggest an explanation for the greater dmax observed experimentally for c-TiO2 (40 Å) compared to a-TiO2 (20 Å), the effects of different electronic parameters on the behavior of current versus TiO2 thickness were simulated. First, the effect of defect band energy or barrier height of the Si−CH3|TiO2 interface was modeled by simulating a 0.1 eV range in the input value for the Eg of TiO2 with the hole mobility set constant at 25705


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Figure 11. (a) Simulated behavior for the short-circuit current density (JSC for hole conduction) for n-Si(111)|CH3|TiO2 photoelectrodes with the contact potentials set to 5.0, 5.2, and 5.4 eV with fixed Eg = 1.7 eV and μp = 0.0001 V cm−2 s−1. (b) Valence band energy for n-Si(111)|CH3|TiO2 photoelectrodes and 5 nm of TiO2 with the contact potentials set to 5.0, 5.2, and 5.4 eV.

energy throughout n-Si(111)−CH3|TiO2 will equilibrate to the work function of the Pt-NPs instead of the formal redox energy of the electrolyte. To simulate this effect, the contact potential was systematically modulated. The simulated short-circuit current density (JSC) versus TiO2 thickness behavior as a function of contact potential is shown in Figure 11a. As the contact potential is increased, the JSC versus TiO2 thickness curve systematically shifts to greater thicknesses, while maintaining its shape. The reason for the shift is due to a change in the electric field across the TiO2. The electric field is proportional to the change in electrostatic potential across the TiO2. As shown in Figure 11b, as the contact potential increases from 5.0 to 5.4 eV, the electric field also increases (as evidenced by the increase in slope of the TiO2 valence band). Therefore, we conclude that the addition of Pt-NPs to the surface of the TiO2 increases the electric field across the TiO2 layer, which facilitates the field-assisted thermionic emission of holes over the potential-energy barrier at the Si−CH3|TiO2 interface and the drift current across the TiO2. As discussed in the Introduction, previous researchers have shown that the deposition of Pt-NPs on metal|insulator|metal (MIM) constructs recovers the electrochemical kinetics via enhanced tunneling probability across the insulating layer. This has clearly been shown to be caused by the large overlap in the DOS in the NPs and redox species in solution.16 However, our scenario utilizes a semiconductor|insulator|metal (MIS) construct, which has a much lower tunneling probability than a MIM construct. Tunneling is proportional to the DOS of each region involved in the tunneling process and DOSsemiconductor ≪ DOSmetal; therefore, tunneling from a semiconductor to a metal is much less probable than tunneling from a metal to a metal.33 Because of the leaky dielectric attribute of the ALD-grown TiO2 films, an alternative and more probable charge transport pathway than tunneling is dominant. Therefore, Pt-NPs not only enhance the tunneling probability across insulating layers, but also can enhance additional charge transfer mechanisms, such as the field-assisted thermionic emission observed in the present n-Si(111)−CH3|“leaky”|TiO2 constructs.

pronounced influence on dmax as that observed with hole mobility. Specifically, a defect concentration of 1 × 1021, 2 × 1020, and 5 × 1018 cm−3 results in a dmax of ∼1.0, 2, and 2.5 nm, respectively. Because traps and recombination centers capture holes traveling across the defect band in the TiO2, an increase in defect concentration will decrease the total hole current through the TiO2 at steady-state. Therefore, it takes a smaller thickness of TiO2 to induce a change in behavior from the current limited by bulk Si (case 1) to limited by thermionic emission at the Si|TiO2 interface (case 2). Even though changing the defect from a trap center to a recombination center had a negligible effect on the calculated current through the device, the defect type had a significant influence on the simulated value for the open-circuit potential, Voc, as shown in Figure 10d. For both types of defects, the Voc remains constant until a critical thickness is reached at which the Voc becomes linear with TiO2 thickness; however, a lower critical thickness is observed for recombination centers and the slope of Voc versus TiO2 thickness is much greater than for traps. This simulated value can be related to the onset potential for the oxidation reaction, Von,A, by the following equation51 Voc = E(A/A−) − Von,A

(9)

where E(A/A−) would be the formal redox potential of ferrocene; therefore, the trend in Voc is directly related to the trend in Von,A. By comparing the trend in simulated Voc with the trend in experimental onset potential (CVs in Figure 7a,b), trap states are found to match the behavior observed with a-TiO2 while recombination centers match the behavior observed with c-TiO2. On the basis of the excellent agreement between the simulations and experimental observations, we conclude that the increase in dmax between a-TiO2 and c-TiO2 is due to a combination of enhanced hole mobility across the TiO2 defectmediated pathway and a change in the defect concentration in the TiO2 layer. The enhanced mobility is supported by the four-point sheet resistance measurements (vide supra) in which a higher conductivity was observed in c-TiO2 versus a-TiO2. One possible explanation is that the Ti3+ defects that generate the defect band lie along the grain boundaries in c-TiO2. Carrier transport across grain boundaries in this case would be much faster than if, for example, the same number of Ti3+ defects were positioned within an amorphous structure. It is likely that the deposition of platinum nanoparticles (PtNPs) will affect the band structure of n-Si(111)−CH3|TiO2 by locally modifying the contact potential. Because the density of states (DOS) in the platinum is greater than the DOS associated with the redox couple, it is likely that the Fermi

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CONCLUSION The findings of the present work are summarized as follows: 1. Ultrathin amorphous and crystalline TiO2 passivating layers (5−60 Å) were controllably prepared on an nSi(111)|CH3 substrate by atomic layer deposition. Both modes of TiO2 deposition exhibit a nearly identical density of electronic defect Ti3+ sites; the Ti3+ defect sites are ascribed to the presence of Ti−N bonds in the ALD precursor (TDMAT). Therefore, the difference in 25706



2. 3.

4.

5.

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conductivity of both amorphous and crystalline TiO2 is not caused by the density of defective Ti3+ sites, but is instead due to a difference in the magnitude of carrier mobility in the material. Both c-TiO2 and a-TiO2 prevented corrosion of Si during anodic operation in MeCN to a similar extent. Charge transport across c-TiO2 and a-TiO2 to a dissolved redox agent (Fc) was inhibited at ∼40 and ∼20 Å, respectively, the former exceeding the limit of chargetransfer that can be attributed to tunneling alone (generally dmax ≈ 20 Å for tunneling) The deposition of Pt-NPs resulted in a significant extension of charge transport distance for c-TiO2 and a-TiO2 (to 50 and 45 Å, respectively) wherein both dmax values are greater than can be attributed to tunneling behavior. Simulation of the charge-transfer behavior using wxAMPS suggests that the photogenerated holes in the bulk Si drift to the Si(111)−CH3|TiO2 interface, at which point the holes thermionically emit over a potential barrier and into an energetic defect band, which is likely established by Ti3+ defects in the TiO2. The deposition of Pt-NPs onto the “leaky”-TiO2 layer significantly increases the electric field in the TiO2, which enhances the extent of charge-transfer through the defect band (Jdrift α μ × E) and the extent of thermionic emission across the Si(111)−CH3|TiO2 interface. The result is chargetransfer across much greater distances than can be achieved with tunneling alone.

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ASSOCIATED CONTENT

* Supporting Information S

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.jpcc.6b08096. Raman spectra, AFM images, SEM images, circuit model for EIS, thickness dependent XPS and thickness dependent CVs (PDF)

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Article

AUTHOR INFORMATION

Corresponding Author

*Tel.: 512-471-4456. E-mail: [email protected]. Author Contributions #

H.J.K. and K.L.K. contributed equally.

Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This research was supported by the US Office of Naval Research (N00014-13-1-0530), the Robert A. Welch Foundation (F-1822), and the UT Austin College of Natural Sciences. The authors also acknowledge the support of the National Science Foundation under Grant No. CBET-1022615, and the International Institute for Carbon Neutral Energy Research (WPI-I2CNER), sponsored by the Japanese Ministry of Education, Culture, Sports, Science and Technology. Funding for the Kratos Axis Ultra XPS was provided by a grant from the National Science Foundation (MRI-0618242), and we acknowledge Mr Ryan Pekarek for assistance in obtaining and analyzing the XPS and Raman spectra, as well as AFM images. 25707


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Charge-Transfer through Ultrathin Film TiO2 on n-Si(111

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