Catalyzed SnO2 Thin Films: Theoretical and Experimental Insights into

Article pubs.acs.org/JPCC

Catalyzed SnO2 Thin Films: Theoretical and Experimental Insights into Fabrication and Electrocatalytic Properties A. Rabis,† D. Kramer,*,‡ E. Fabbri,*,† M. Worsdale,‡ R. Kötz,† and T. J. Schmidt† †

Electrochemistry Laboratory, Paul Scherrer Institut, 5232 Villigen PSI, Switzerland Engineering Sciences, University of Southampton, SO17 1BJ Southampton, U.K.

‡

S Supporting Information *

ABSTRACT: SnO2 thin films are studied experimentally and from firstprinciples as model supports for Pt nanoparticle catalysts in an acidic environment. SnO2 thin film supports are attractive model systems because composition, microstructure, and surface termination can be tailored by varying the deposition parameters. SnO2 films are synthesized by reactive dc magnetron sputtering, and the effects of the deposition conditions on the physicochemical and electrochemical properties are investigated experimentally and theoretically. Variation of the deposition conditions results in limited longrange order SnO or SnO2 films. Annealing in either case leads to wellcrystallized SnO2 films, but with different growth directions. Films deposited as SnO2 show only growth along the [110] direction, while SnO2 films formed from deposited SnO show no preferred orientations. Hybrid density functional theory (DFT) suggests that growth along the [110] direction is driven by (110) being the lowest energy surface, while the loss of orientation in the SnO derived films originates from an almost degenerate set of surface energies at the SnO|SnO2 equilibrium. The oxygen reduction reaction activity of Pt nanoparticles depends on the SnO2 film orientation. A 2-fold higher catalytic activity is observed for Pt nanoparticles on the SnO2 film without preferential orientation compared to Pt on SnO2 grown along the [110] direction, pointing to the presence of strong surface-dependent metal−support interaction. conditions.9−11 An increase in stability against potential cycling compared to conventional Pt supported on high surface area carbon (Pt/C) catalysts was found by different studies using model electrodes as well as single cell setup.9,12 The group of Sasaki13 performed rotating disk electrode (RDE) measurements on Pt deposited on SnO2 powders and showed only a small decrease (less than 20%) in electrochemical active surface area of the Pt catalyst after 60 000 potential cycles between 0.9 and 1.3 V in 0.1 M HClO4 at 25 °C.13 However, their Pt/SnO2 catalysts showed lower activity toward ORR compared to conventional Pt/C catalysts. Other groups like Saha et al.10 or Watanabe and co-workers 11 used nanostructured SnO 2 supports with high surface areas showing similar or even higher ORR activities of their Pt/SnO2 catalysts compared to Pt/C. They attribute the high eletrocatalytic activities to the unique 3D structures, the electronic properties, and the low impurity level of their SnO2 compared to carbon blacks and to the occurrence of strong interactions between Pt catalyst particles and the SnO2 surface. Strong metal−support interaction (SMSI) was observed for Pt catalysts supported on oxides like SnO2 and is assumed to influence the electrocatalytic activity of the catalyst.14−16 Some models like

1. INTRODUCTION Oxygen reduction reaction (ORR) catalysts in polymer electrolyte fuel cells (PEFCs), typically Pt nanoparticles on carbon blacks, face harsh and strongly oxidizing conditions like potentials up to 1.5 V vs reversible hydrogen electrode (RHE).1 Corrosion of the carbon support is a major problem and limits the lifetime of PEFCs. Therefore, alternative carbon-free supports are of interest. Oxides might be candidate materials to support PEFC cathode catalysts because of their thermodynamic stability against oxidizing environments and high potentials, although it is challenging to engineer materials with sufficient electronic mobility, high surface area, and stability across the relevant potential range (i.e., ranging from about 0 V up to about 1.5 V vs RHE). Various metal oxides (e.g., TiOx, WOx, and SnOx) have been evaluated as possible support materials for PEFC catalysts.2 To investigate the feasibility of using oxides as support for PEFC catalysts, SnO2 has been selected as model system since the Pourbaix diagram suggests thermodynamic stability above 0 V at pH = 1,3 and thin films with electrical conductivities on par with amorphous carbon and approaching 103 S cm−1 can be produced.4 Pt supported on SnO2 is an interesting catalyst system for various chemical reactions, e.g., oxidation of CO5 and lowweight alcohols such as methanol6 and ethanol.7,8 Also, some investigations were performed to examine the activity and stability of Pt/SnO2 catalysts under oxygen reduction © 2014 American Chemical Society

Received: December 8, 2013 Revised: May 2, 2014 Published: May 7, 2014 11292

dx.doi.org/10.1021/jp4120139 | J. Phys. Chem. C 2014, 118, 11292−11302

The Journal of Physical Chemistry C

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the charge transfer17 or the encapsulation18 model have been proposed, but the detailed mechanism of the SMSI phenomenon is not fully understood. SnO2 thin film supports produced by physical deposition techniques are attractive model systems because important physicochemical properties, such as composition, microstructure, and surface termination, can be controlled and studied in detail. Furthermore, a defined support surface allows studying support−metal interactions as a function of surface orientation and termination when catalyst nanoparticles are also deposited in a controlled manner. Therefore, we produced dense SnO2 thin films19 by reactive dc magnetron sputtering tailoring the deposition parameters in order to obtain films featuring different surface stoichiometries (i.e., a fully oxidized and a partially reduced surface) and surface orientations. The electrochemical properties of the bare thin films were investigated, highlighting differences between the fully oxidized and the partially reduced surfaces. Additionally, Pt nanoparticles were deposited by dc magnetron sputtering on both SnO2 supports, and material and catalytic properties of the resulting catalyst systems were investigated in terms of ORR activity and selectivity.

Table 1. Summary of SnO2 Thin Film Samples Prepared by DC Magnetron Sputtering and Subsequent Annealing sample

reduced SnO2

O2 content (O2/O2 + Ar) back pressure sputter time sputter power nominal film thickness deposited films crystal structure space group nominal composition preferred orientation annealing temperature annealing atmosphere annealing time annealed films crystal structure space group nominal composition preferred orientation crystallite size

2. EXPERIMENTAL SECTION 2.1. Synthesis. Thin Films. Tin oxide films were prepared by using reactive dc magnetron sputtering according to the procedure described in detail in ref 19. The target was a metallic tin disk (99,99% purity, Umicore AG) with a diameter of 75 mm and 6 mm thickness. The substrate for electrochemical investigations was glassy carbon (GC), and glass substrates were used as reference for material characterization. Prior to the deposition, the GC substrates were pretreated in concentrated HNO3 at 75 °C for about 12 h to ensure optimal adhesion on the amorphous carbon. A 2 nm thick Ti interlayer was deposited by dc magnetron sputtering before the SnOx to further increase adhesion on the substrates. The SnOx deposition was performed without any additional heating of the substrates. Discharge power and gas flows (Ar and O2) were varied to obtain films with predominantly SnO or SnO2 stoichiometry (see Table 1). A film thickness of about 200 nm was achieved by controlling the deposition time. The films were subsequently annealed at 400 °C in an argon atmosphere for 4 h after deposition. With respect to the SnO/SnO2 thermodynamic equilibrium the two different deposition conditions can be described as reducing or oxidizing. The films obtained after annealing via conversion of SnO are called reduced SnO2, and those obtained by direct deposition of SnO2 are called oxidized SnO2 throughout this article (see Table 1). The preparation condition for the reduced SnO2 will therefore be called reducing conditions and those for the oxidized SnO2 oxidizing conditions. Pt Deposition. Pt was deposited by dc magnetron sputtering on SnO2 supports from a metallic Pt target (99.99% purity) after they have been undergone the annealing step. To achieve a Pt loading of 2 μg cm−2, the discharge power, deposition time, and Ar flow were fixed at 50 W, 1 s, and 10 sccm, respectively.20 2.2. Characterization. Physical Characterization. Crystal structure and preferred growth directions were obtained from a Bruker D8 Advance X-ray diffraction (XRD) system (Cu Kα, λ = 1.5418 Å) in Bragg−Brentano geometry (θ/2θ). The Scherrer equation

d = 0.9

% 10−6 mbar min W nm

°C h

nm

oxidized SnO2

30 5 5 100 200

40 5 15 80 200

litharge P4/nmm (129) SnO

rutile P42/mnm (136) SnO2

[101]

[110]

400 argon 4

400 argon 4



none

[110]

∼14

∼14

k1 cos θ B2

was used to determine the average particle diameter d from the X-ray wavelength k1, the angle of the [110] reflection, and the full width at half-maximum B2. Information of the chemical state of Sn, O, and Pt at the surfaces was gained from X-ray photoelectron spectroscopy (XPS) data by analyzing the Sn 3d, O 1s, and Pt 4f signals, respectively. XPS measurements were performed using a VG ESCALAB 220iXL spectrometer (Thermo Fischer Scientific) equipped with an Al Kα monochromatic source (spot size: 500 μm; power: 150 W) and a magnetic lens system. Survey spectra were initially recorded at low resolution, and high-resolution spectra were collected for quantitative analysis of SnO2 and Pt/SnO2 compositions. All spectra were calibrated with the sp2hybridized carbon component of C 1s at 284.5 eV. Background subtraction has been performed according to Shirley,21 and atomic sensitivity factors (ASF) of Scofield22 were applied. Surface electronic states of the SnO2 films were gauged by recording high-resolution XPS valence band spectra at binding energies between 0 and 10 eV. After Shirley-type background subtraction to remove the secondary electron scattering, the dband center relative to the Fermi level was calculated from the density of state based on dcenter =

∫ N (ε)ε d ε ∫ N (ε) d ε

where N is the density of states and ε is the energy of states.23 Ellipsometry was performed using a spectroscopic ellipsometer (MOSS model ES 4G, Sopra, France) at an angle of incidence of 70°. In our measurements the photon energies were scanned from 2.0 to 4.0 eV in 0.1 eV intervals. For the data analysis a stratified layer model was assumed consisting of a Si substrate/ SiO2 interlayer/Pt film. Best fits of the measured data, 11293



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particularly in thin films, were obtained assuming a certain void fraction in the Pt film, in agreement with the nanoparticulate nature of the deposited Pt catalysts as described in our previous work.20 The optical properties of the Pt film with voids were calculated according to the Bruggeman effective medium theory.24,25 Finally, scanning electrode microscopy (SEM) was used to investigate the SnO2 surface, while transmission electron microscopy (TEM) (FEI Morgagni 268, 120 kV, image acquisition in bright field) images were taken to further investigate the sputtered Pt nanoparticle morphology. Electrochemical Characterization. Electrochemical measurements were conducted in a standard electrochemical cell using an interchangeable ring disk electrode setup with a programmable potentiostat (VMP3 workstation, Biologic France) and a rotation control unit (Pine Instruments). All potentials refer to RHE scale. SnO2 supports were investigated by cyclic voltammetry in Ar-saturated 0.1 M HClO4. ORR polarization curves were obtained in oxygen-saturated electrolyte. The experiments were performed at room temperature, varying the rotation speed (400−2500 rpm). While the potential of the disk was swept from 1.00 to 0.05 V at a scan rate of 5 mV s−1, the ring potential was held at 1.2 V vs RHE in order to detect the formation of H2O2. The fraction of peroxide produced during ORR was estimated from the measured disk and ring currents applying the equation26

calculations were fully relaxed (unit cell and atomic coordinates) within the symmetry group using default convergence criteria. Atomic coordinates within surface slabs were relaxed, but unit cells were fixed at dimensions equivalent to the fully relaxed bulk unit cell.

3. RESULTS 3.1. Thin-Film Structure and Morphology. Two different SnO2 films were sputtered at room temperature on glass using the sputter conditions given in Table 1. We will refer to the two sets of samples as “reduced” and “oxidized” films for reasons that will become apparent shortly. The XRD patterns shown in Figure 1 were obtained from the as-prepared samples. Both patterns have rather broad features,

I

X H2O2 [%] =

2 NR ID +

IR N

× 100%

where XH2O2 is the percentage of peroxide, IR and ID are the ring and disk currents, respectively, and N is the collection efficiency of N = 0.2 ± 0.02 determined experimentally as described before.26 2.3. First-Principles Calculations. Several low index surfaces of SnO2 were investigated by hybrid density functional theory. The CRYSTAL09 software package,27 which is based on the expansion of the crystalline orbitals as a linear combination of atom-centered Gaussian orbitals, was used for all calculations. A triple-valence all-electron basis set with an 8411d(1) contraction (one s, three sp, and one d shell) was used for O; the most diffuse sp and d exponents are αO = 0.168 bohr and αO = 0.45 bohr, respectively. A large core pseudopotential due to Durand and Barthelat28 was used for Sn with a doublevalence basis set having a 21 contraction; the sp exponents are αSn = 0.234 bohr and αSn = 0.12 bohr, respectively. Note that this is insufficient to describe strongly polarized Sn in bulk SnO with acceptable accuracy. A more diffuse basis, however, led to linear dependence in octahedral environments such as in SnO2, and the total energy of bulk SnO2 was converged to within 10 meV using αSn = 0.12 bohr for the most diffuse orbital. We also checked the sensitivity of the surface calculations with respect to the most diffuse exponent of surface Sn, because the undercoordinated environment could potentially lead to stronger polarization,29 and found negligible sensitivity around αSn = 0.12 bohr. The B3LYP hybrid-exchange functional was used. The Coulomb and exchange series were summed directly and truncated using overlap criteria with thresholds of 10−7, 10−7, 10−7, 10−7, and 10−14 (cf. refs 27 and 30). Reciprocal space was covered with a Monkhorst−Pack net with shrinking factor 16. This high shrinking factor became necessary to accurately describe some surfaces with metallic surface states. Self-consistency cycles were converged up to 10−7 Ha. Bulk

Figure 1. XRD patterns of reduced and oxidized SnO2 thin films as deposited and after annealing at 400 °C in argon. The XRD pattern of stoichiometric SnO2 is included as reference.

indicating that films with limited long-range order were deposited. Nonetheless, the two main reflections at 26.5° of the oxidized film and 29.9° of the reduced film can be attributed to the [110] reflection of SnO2 (rutile; P42/mnn) and the [101] reflection of SnO (litharge; P4/nmm), respectively. Cassiterite (i.e., rutile SnO2) was therefore produced by sputtering in an oxygen-rich atmosphere, while the deposition under less O2-containing atmosphere resulted mostly in the formation of litharge with nominal stoichiometry SnO. Both samples were annealed in argon at 400 °C for 4 h. The XRD patterns of the annealed films depicted in Figure 1 show significantly less broadening of the XRD reflections compared to as-deposited films, indicating well-crystallized films. A sharpening of the [110] rutile reflection is the major change of the XRD pattern obtained from the oxidized film before and after annealing. The XRD pattern of the reduced film changes more substantially by annealing the sample. No reflections are seen close to 2Θ = 30° after annealing, which indicates that litharge SnO is unstable under the chosen annealing conditions. Instead, a new triplet of sharp reflections at 26.6°, 33.9°, and 51.7° develops. These match the [110], [101], and [211] reflections of rutile SnO2, suggesting that annealing in Ar atmosphere resulted in a reorganization of the films crystal structure to form the high-temperature stable casserite (i.e., rutile SnO2). The change in crystal structure under an Ar atmosphere leads to the assumption that the as deposited 11294



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reduced film contained an excess of O2 or an amorphous SnO2 phase. An average particle size after annealing of about 14 nm can be estimated for both films from the XRD patterns using the Scherrer equation. This was confirmed visually by SEM investigations (see Figure S1). A full coverage and welldeveloped films free of holes and other damage are observed. As-prepared films are more or less featureless, further corroborating the XRD results which indicated a limited long-range order. Similarly, the well-developed microstructure of the annealed films indicates that grains were crystallized during annealing with an average grain size of >10 nm, which matches well with the grain size of 14 nm estimated from the XRD data. 3.2. Surface Orientation and Termination. Note that the [101] and [211] reflections are largely missing from the XRD patterns of the oxidized film before and after annealing (cf. Figure 1). This suggests preferred orientation of the oxidized film along the [110] direction. The XRD data of the reduced film, on the other hand, do not show strong evidence for a preferred orientation after annealing. The energetics of several low-index surfaces of rutile SnO2 were investigated using hybrid density functional theory to rationalize the preferred orientation of the oxidized film as well as the lack of preferred orientation of the reduced film. Details of the three investigated surfaces, (110), (101), and (211), are given in Table 2. The surfaces were modeled as slabs cut from

the center plane was enforced to ensure equal surfaces (even after relaxation) on both sides of the slab. All slabs were checked for convergence with respect to slab thickness, and we found negligible sensitivity of the surface energy on slab thickness provided slabs are at least 18 atomic layers thick. Two terminations were considered for all surfaces: oxidized and reduced. The fully oxidized surfaces feature Sn4+ throughout, while Sn2+ in the surface layer charge compensates for the lower oxygen coverage of the reduced surfaces. Ball models of the investigated surfaces are depicted in Figure 2. The surface energy is calculated from the total slab energy Eslab by 2Aλ = Eslab − NSn(E bulk + ΓOμO)

with NSn being the number of Sn atoms in the slab unit cell, ΓO being the excess oxygen defined by ΓO = NO/NSn − 2, Ebulk the bulk energy per formula unit, and γ the surface energy.32 Note that the dependence on just the chemical potential of oxygen μO is not an approximation because μSn and μO are not independent variables; they are linked via the bulk energy. We have chosen μO as the independent variable because this better reflects experimental control (e.g., by the use of O2, air, or Ar as atmosphere in the sputtering or annealing chamber). The calculated surface energies for the investigated directions and terminations (four directions and two terminations each) are shown in Figure 3. Note that the surface energies of the fully oxidized surfaces do not depend on the chemical potential because Sn4+ prevails throughout the slab. We find the fully oxidized (110) surface with a surface energy of 1570 mJ m−2 to be energetically favorable under oxidizing conditions, followed by (100) with 1695 mJ m−2 and (101) with 2130 mJ m−2. We also briefly investigated fully oxidized (001) but did not consider this surface further due to the high surface energy of 2620 mJ m−2. The surface energy under reducing conditions depends linearly on the chemical potential because of the nonstoichiometry of the reduced surfaces. We find the (101) surface to be the lowest energy surface under reducing conditions followed by (100) and (110). The relative order, therefore, reverses from (110) < (100) < (101) under oxidizing conditions to (101) < (100) < (110) under reducing conditions. We also investigated partially

Table 2. Details of the Low Index Surfaces Investigated by DFT surface Sn−O coordination

unit cell parameter surface

a [A]

b [A]

γ [deg]

oxidized

reduced

(001) (100) (101) (110)

4.719 3.158 4.719 3.158

4.719 4.719 5.678 6.674

90.0 90.0 90.0 90.0

5 5 5 5/6

3 3 3

the fully relaxed bulk with lattice constants of 4.719 and 3.158 Å. These compare well with the experimental values of 4.737 and 3.186 Å,31 respectively. Mirror symmetry with respect to

Figure 2. Ball and stick models of theoretically investigated surfaces and terminations of rutile SnO2; surface oxygen is shown as large gray, bulk oxygen as large white, Sn4+ as small black, and Sn2+ as small dark-gray circles. 11295



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temperatures >300 °C.34 The spectrum of the O 1s (Figure 4b) can be resolved into two peaks at relatively lower binding energy (530.6 eV) and high binding energy (532.1/532.4 eV) by fitting with two Gaussian distributions. They are assigned to O 1s core peaks of O2− ions in Sn−O band (lattice oxygen) and surface adsorbed hydroxyl groups (OH-band), respectively. A shift to slightly higher BEs can be observed for the OH-band of the reduced SnO2. Integration of the Sn 3d and O 1s spectra leads to different Sn:O ratio of the SnO2 films found in the surface layer. SnO2 thin films directly prepared under oxidizing conditions have a surface Sn:O atomic ratio of 35:65. In contrast, SnO2 indirectly produced by depositing SnO first and annealing later shows a Sn−O surface composition of roughly 45 at. % Sn and 55 at. % O. The surface electronic states for both SnO2 films were examined by XPS. Figure 4c illustrates the partial valence band of the reduced and oxidized SnO2 films. The oxidized SnO2 surface shows a density of states starting at 2.5 eV while the DOS of the reduced SnO2 surface is somewhat higher and starts already close to the Fermi level. The values for the d-band center are 5.88 and 6.07 eV for the reduced and oxidized SnO2, respectively, clearly demonstrating the different surface electronic conditions of the two surfaces. 3.3. Morphology of Catalyzed Films. The films were catalyzed by depositing Pt as described in ref 20. Using magnetron sputtering, 2 μg of Pt per cm2 was deposited. This corresponds to a 0.9 nm thick film, if Pt would form a dense film on SnO2. Spectroscopic ellipsometry was performed on these samples and resulted in a film thickness of 2 nm with a void fraction of about 40%, indicating the presence of Pt clusters or agglomerates which are not fully covering the support. For estimating the Pt particle size TEM images of Pt deposited on carbon film grids were acquired (Figure 5, 2 μg cm−2). Pt is homogeneously distributed; primary particles with diameters of about 2−3 nm are observed together with small agglomerates, consistent with the Pt morphology indicated by the ellipsometry measurements and in agreement with our previous work.20 XPS results are presented in Figure 6. A representative Pt 4f doublet profile of the Pt deposited on SnO2 is shown in Figure 6a with the Pt 4f 7/2 peak maximum at 71.6 eV, which corresponds to Pt0. As shown in Figure 6b for the oxidized SnO2 surface, the binding energies (BE) of the Sn 3d5/2 and Sn 3d7/2 are 486.7 and 495.1 eV, respectively. Only a small shoulder forms at lower binding energy, which can possibly be attributed to metallic Sn0, indicating that mainly Sn4+ ions contribute to the Sn 3d signal with a minor contribution from Sn0. We, therefore, conclude that Pt forms particles on SnO2, although there is

Figure 3. Calculated surface energies of selected low index surfaces of rutile SnO2 as a function of oxygen chemical potential and electrode voltage; fully oxidized (110), (101), (100), and (001), partially reduced (110), and fully reduced (110), (101), and (100) were considered; voltage scale is based on in vacuum calculations and neglects presence of the condensed phase in electrochemical environments.

oxidized [110], but our calculations indicate that this surface termination is generally energetically unfavorable relative to other terminations. These results are in agreement with the experimentally observed growth behavior as discussed above. XPS measurements were performed to evaluate the surface composition of the two SnO2 thin films after the thermal treatment. As shown in Figure 4a, the binding energies (BE) of the Sn 3d5/2 and Sn 3d7/2 are 486.7 and 495.1 eV, respectively. Determining the ratio between Sn2+:Sn4+ at the surface of the SnO2 films is rather questionable by simply comparing Sn 3d BE values due to the overlapping of the 3d binding energy for Sn2+ and Sn4+ (Sn 3d5/2 BE: Sn2+ from 485.5 to 487 eV and Sn4+ from 486 to 488 eV33). On the other hand, from the XPS spectra in Figure 4a we can exclude the presence of tin metal after the thermal treatment within the outward 5−10 nm of the thin films (mean free path of photoelectrons); nevertheless, slow decomposition of SnO is reported for bulk material and

Figure 4. Representative X-ray photoemission spectra of (a) Sn 3d, (b) O 1s, and (c) the valence band for the reduced and oxidized SnO2 thin films. 11296



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100 mV s−1 for 20 times in 0.1 M HClO4 before electrocatalytic activity was investigated. This ensured reproducible results and a C-free surface of the Pt/SnO2 catalyst. The electrochemical characteristics of the catalyzed films closely resemble those of polycrystalline Pt. In general, all CVs of the activated Pt/SnO2 electrodes show the characteristic Pt features visible in Figure 8b such as the hydrogen adsorption/desorption region followed by the double-layer region and oxide formation/reduction. The CV of the Pt deposited on the reduced SnO2 shows a broad peak at E = 0.5 V in the positive going scan and a broadening (starting at a more positive potential) of the hydrogen underpotential deposition (Hupd) region in the negative going scan. These features are most likely related to the reduced SnO2 support, which is showing the same redox couple at slightly lower potentials. That implies that charges related to SnO2 surface processes may contribute to the Pt Hupd charges, and therefore, the determination of the electrochemical active surface area of Pt by integration of Hupd currents is not straightforward. The ORR activity of both catalyzed SnO2 films was evaluated using rotating ring-disk electrodes (RRDE). This is exemplified in Figure 8a, which shows the ring and disk currents obtained from a 2 μg cm−2 Pt on reduced and oxidized SnO2 electrode at a rotation speed of 1600 rpm and at a scan rate of 5 mV s−1. Furthermore, a set of ORR polarization curves at various rotation speeds (400−2500 rpm) were recorded for both samples. Simultaneously, the ring currents have been monitored to detect formation of H2O2.21 The ORR of Pt on both SnO2 supports is under mixed kinetic-diffusion control in the potential range between 0.9 and 0.7 V, followed by a region where diffusion limiting currents (plateau between 0.6 and 0.2 V) can be observed, comparable to those obtained for conventional Pt/C catalyst.30 The oxygen reduction currents decrease below 0.1 V and coincide with a significant increase of the ring current, indicating that H2O2 becomes a significant reaction product at these low potentials. Quantifying the fraction of H2O2 produced results in about 10% H2O2 formation at 0.05 V for Pt deposited on the reduced SnO2, while Pt on the oxidized SnO2 produces double amounts of H2O2 of about 20%. Small fractions of about 4−5% are being produced in the potential range between 0.6 and 0.2 V (cf. Figure 8a), typically also observed on Pt/C catalysts.36,37 The ORR currents can be evaluated according to the equation

Figure 5. TEM image of 2 μg cm−2 Pt deposited by dc magnetron sputtering.

some indication that the particles might be alloyed with metallic Sn. 3.4. Electrochemical Properties. Electrochemical Properties of the Supports. The electrochemical properties of the films in acidic environment were examined by cyclic voltammetry (CV). Figure 7a shows stable cyclic voltammograms of the reduced and oxidized SnO2 film in Ar-saturated 0.1 M HClO4. Since there are no distinct features observed in the CV of the oxidized SnO2 between 0 and 1 V, the fully oxidized SnO2 surface can be considered as redox inactive. On the other hand, the reduced film shows a distinct, stable redox couple at around 0.3−0.4 V, although a decrease of the strength of the associated oxidation/reduction peaks was observed within the first couple of cycles. Thermodynamics suggest that faradaic SnO2 reduction is not favorable at positive potentials relative to RHE,3 and indeed we see faradaic SnO2 reduction/ oxidation currents only if the voltage range is extended to negative potentials relative to RHE (see Figure 7b). The electrochemical response of the reduced and oxidized films in the negative potential range was very similar and is shown in Figure 7b for the oxidized sample. Large currents, typical of a faradaic reaction, were only observed below −0.2 V, showing a redox couple in agreement with reduction of SnO2 (SnO2|SnO = −0.24 V and SnO2|Sn = −0.22 V at pH = 1, calculated using formation energies from Barin35). ORR Activity of Catalyzed Films. All samples were equilibrated by potential cycling between 0.05 and 1.2 V at

1 1 1 1 1 = + = + 1/2 i iD iK iK BcOω

Figure 6. (a) Pt 4f core XPS spectra performed on Pt/SnO2 thin film catalyst. (b) Comparison of Sn 3d XPS core level spectra of SnO2 thin film (dotted line) and Pt deposited on SnO2 (full line). 11297



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Figure 7. (a) Stable CVs of oxidized and reduced SnO2 thin films at positive potentials relative to RHE, sweep rate at 100 mV/s, electrolyte is Arsaturated 0.1 M HClO4. (b) CV of oxidized SnO2 thin film recorded in 0.1 M HClO4 in different potential range.

Figure 8. (a) Ring and disk current potential curves for 2 μg cm−2 Pt on oxidized and reduced SnO2 in oxygen-saturated 0.1 M HClO4 at 1600 rpm and room temperature, cathodic direction of potential scan. (b) Cyclic voltammograms of 2 μg cm−2 Pt/SnO2 model catalysts in 0.1 M HClO4 with 50 mV/s at room temperature. (c) Tafel plots relative to the ORR curves of Pt on oxidized and reduced SnO2.

According to the above equation, the Levich−Koutecky plots for Pt on oxidized and reduced SnO2 at various potentials were derived from the RRDE polarization curves at different rotation speeds (not shown). The linearity of the Levich−Koutecky plots implies a first-order dependence of ORR kinetics on Pt/ SnO2. From the slope of the Levich−Koutecky plots a B-factor of about 3.0 × 10−2 mA/rpm can be obtained for the different Pt/SnO2 catalysts. For comparison, the theoretical value is about 4.27 × 10−2 mA/rpm, calculated for a four-electron process using literature data (oxygen solubility cO = 1.26 × 10−3

with B = 0.62DO2/3υ−1/6

where iK and iD represent the kinetic and diffusion-limited current densities, respectively, ω is the rotation rate, F the Faradaic constant, n the number of overall transferred electrons, cO the oxygen solubility, DO the oxygen diffusion coefficient, and ν the kinematic viscosity of the electrolyte. 11298



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mol/L, oxygen diffusion coefficient D = 1.93 × 10−5 cm/s, and the kinematic viscosity of the electrolyte v = 1.009 × 10−2 cm/ s).15 Having confirmed that ORR on Pt/SnO2 thin film electrodes follows first-order kinetics, the kinetic current densities can be evaluated from the ORR polarization curves by using the above equation. Tafel plots for Pt deposited on oxidized and reduced SnO2 are shown in Figure 8c. Values at 0.9 V were taken to compare the kinetic activities of the two Pt/SnO2 model catalysts. The obtained values are compared in Table 3. Strikingly evident, the activity of Pt supported on reduced SnO2 is doubled as compared to the activity found for Pt on the oxidized SnO2 support.

deposited directly as SnO2 in an oxidizing environment (cf. Figure 1). Usually, thin-film growth is governed by a competition between surface energy and strain energy.41 We, however, believe it is sufficient to discuss the growth behavior of our films based on surface energy alone. Because of the amorphous or polycrystalline nature of the Ti interlayer, it is unlikely that the substrate will induce strain into the thin film. There is only a very small range of μO where (101) is the lowest energy surface of thermodynamically stable SnO2. This implies demanding requirements regarding the precision of experimental parameters, if one wishes to grow SnO2 films preferentially along [101]. If the atmosphere is just a bit too reducing, SnO will form. If the atmosphere is slightly too oxidizing, growth will be along [110]. The equilibrium between the competing bulk phases at μO = −3 eV coincides with an environment where four different surface states yield very similar surface energies, as highlighted by the dashed line in Figure 3. This might provide a rational for the observed growth behavior under reducing conditions. As SnO was originally deposited under reducing conditions, which only converted into SnO2 via annealing at 400 °C, both phases were present during annealing. Although SnO and SnO2 were not formally in thermodynamic equilibrium during the conversion process, the equilibrium oxygen chemical potential of about −3 eV, nonetheless, reflects the environment under which SnO2 is formed in this case. Under these conditions (110), (100), and (101) all yield roughly the same surface energy. It is, therefore, logical to expect no preferred growth along either of these directions. There is simply no significant thermodynamic advantage in favoring (110) over (100) or (101) or vice versa on the grounds of minimizing the surface energy. This is indeed what the XRD results (cf. Figure 1) suggest: random orientation of the grains. Moreover, our calculations predict some of the competing surfaces of the randomly oriented, indirectly grown SnO2 film to be reduced. This is corroborated by the XPS results, which show a substantial increase of Sn near the surface. Indeed, the ratio of about 1:1.25 comes close to 1:1, the theoretical value for fully reduced surfaces. The experimental ratio of 1:1.25 might be taken as an indication of the presence of fully oxidized (110) in conjunction with reduced (101) and (100), respectively. For the film grown along [110] our calculations imply a fully oxidized, stoichiometric surface layer. This is in agreement with the XPS result, which indicates for that sample a Sn:O ratio close to the theoretical (1:2). 4.2. Surface Redox Behavior. The oxidized SnO2 surface is electrochemically redox-inactive between 0.0 and 1.0 V, while a stable redox couple at 0.2−0.4 V was observed for the reduced SnO2. Faradaic redox behavior of both samples was only observed at negative potentials relative to RHE, in agreement with thermodynamic predictions.3 Interestingly, the equilibrium between Sn2+(aq) and Sn4+(aq) is at 0.151 V,3 making it conceivable that the observed peaks in the CV are due to oxidation/reduction of dissolved Sn. Even though the redox couple was observed only for the reduced SnO2 film, we cannot exclude that differences in the surface termination may result in diverse dissolution stability. Our theoretical calculations suggest that the reduced SnO2 film might feature (101) as well as (110) terminations after annealing, while there is strong evidence that oxidized SnO2 mainly features oxidized (110). It is, therefore, tempting to seek the reason for the activity at 0.2−0.4 V in the presence of (101). Indeed, Figure 3 suggests that the equilibrium potential

Table 3. ORR Kinetic Parameters of Pt on SnO2 (SD = Standard Deviation) support

Pt loading [μg/cm2]

ikin [μA/cm2geo]

ispec [μA/cm2Pt]

SD

reduced SnO2 oxidized SnO2

2 2

26.2 10.2

25.6 10

4 7

4. DISCUSSION 4.1. Film Growth Behavior. Directly depositing SnO2 under oxidizing conditions led to SnO2 thin films with a preferential orientation along the [110] direction. Indirect deposition via growing SnO and subsequent conversion into SnO2 lifted the strong preference for growth along [110]. This can be rationalized in terms of surface energy. Since Sn4+ in SnO2 is the highest oxidation state of Sn, there is no strict upper limit for the oxygen chemical potential (i.e., there is no competing Sn oxide with an even higher oxidation state). However, we have chosen the gaseous O2 dimer at 0 K as the reference state as an indication of what might conveniently be achieved experimentally. Note that an oxygen atmosphere at elevated temperatures will be slightly less oxidizing than the chosen reference state, mostly due to entropic effects. A lower boundary, however, is strictly defined by the phase equilibrium between SnO and SnO2. A lower boundary of about μ(O) = −3 eV can be calculated using experimental formation energies.35 SnO is the thermodynamically favored bulk phase below this value. According to our calculations, there are two thermodynamically favorable lowest energy surfaces: oxidized (110) under oxidizing conditions and reduced (101) under reducing conditions. This is in qualitative agreement with previous DFT studies by Bergermayer and Tanaka, although they predict the equilibrium between reduced and oxidized surfaces to be at a more oxidizing chemical oxygen potential of about −1.6 eV,38 while we predict a value of about −3 eV. The comprehensive study of Batzill et al. likewise used the general gradient approximation (GGA) and shows a similar trend toward higher chemical potentials for the equilibria between reduced and oxidized surfaces of SnO2.39 The discrepancy might be a consequence of their reliance on the GGA rather than a hybrid functional. Especially, the energetics of the oxygen dimer reference state are known to be unreliable in the GGA.40 In fact, our results agree also quantitatively with the study of Bergermayer and Tanaka38 and Batzill et al.39 if the correction of −1.36 eV proposed by Wang, Maxisch, and Ceder for the oxygen dimer in the GGA is taken into consideration.40 The surface energetics depicted in Figure 3 naturally explains the observed preferential growth along [110] if films are 11299



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for oxidizing/reducing (101) is about 250 mV more positive than the oxidation/reduction of SnO2. Although the experimentally observed potential for the positive oxidation peak is somewhat more positive than what we can account for based on our calculations, our results are at least qualitatively in agreement with attributing this redox peak to the oxidation/ reduction of (101). We acknowledge that more work is needed to corroborate this interpretation. For instance, we did not consider OH-terminated surfaces. However, if this redox peak can be attributed to the presence of (101), its gradual decline with cycle number might indicate a gradual recrystallization of the (101) surface under electrochemical conditions in favor of the lowest energy surface at positive potentials, namely oxidized (110). 4.3. Sn−Pt Alloy Formation. Examinations on the surface chemical composition of the Pt/SnO2 catalysts revealed the presence of Sn0 species after deposition of Pt. This is most probably a result of the Pt deposition by dc magnetron sputtering where the backscattering of high energetic Ar ions from the Pt target or the deposition of the heavy Pt clusters reduces the oxygen in the tin oxide lattice. As already reported for different Pt/oxide systems,42 one could argue that alloy formation occurs; however, the XPS results are not fully conclusive. The XPS Pt signals do not give an indication for alloy formation, neither for Pt deposited on reduced nor oxidized SnO2. The same conclusion, based on XPS analysis, was also reported by Katsiev et al.43 for a similar catalyst/ support system, namely Pd deposited on SnO2 by vapor deposition at room temperature. The CVs of the Pt/SnO2 thin film electrodes (cf. Figure 8b) provide more evidence for alloy formation. The Pt−O formation features (expected at E = 0.8−0.9 V) are strongly suppressed for both electrodes, which is observed for PtSn alloys.44,45 Also, the rather featureless Hupd region is indicative of a rough interface with high defect density. This might be a consequence of surface Sn leaching from the alloy in the electrochemical environment.46 These are some indications for PtSn alloy formation, but we wish to stress that our investigations are not fully conclusive on the presence of PtSn alloys. 4.4. ORR Activity of Pt/SnO2 Catalysts. The reduction of oxygen is known to proceed either in a serial pathway through the formation of H2O2 intermediates which can be further reduced to water (transfer of 2 + 2 e−) or on a pathway without formation of peroxo species (transfer of 4 consecutive e−). The general assumptions based on DFT calculations, however, point to the energetically more favorable pathway for the ORR to water through the peroxo state.47,48 Owing to the deleterious effect of H2O2 on the stability of polymer electrolyte membranes and ionomers, the extent of H2O2 formation on the cathode catalyst in the potential region above 0.7 V is a critical criterion for the choice of suitable catalysts. The peroxide formation above 0.7 V is less than 1.0% for the investigated Pt/SnO2 catalysts, and almost a complete reduction of oxygen to water is given. The maximum peroxide formation occurs at potentials below 0.1 V, reaching levels of 10% and 20% for Pt on reduced and oxidized SnO2, respectively. The increase of H2O2 in the Hupd region is generally attributed to hydrogen (co)adsorption, which blocks the dissociation of peroxo species49 and not specific for our system. Similar amounts of peroxide formation below 0.1 V were found for polycrystalline Pt (5−15%) and several Pt/C catalysts (10−20%).30

Pt nanoparticles are about twice as active on the reduced SnO2 than on the oxidized SnO2. The enhanced ORR activity is potentially a result of variations in the physical−chemical characteristics of Pt nanoparticles caused by or related to the two different SnO2 support surfaces. A variety of crystallographic and electronic effects can be envisioned, and further studies are needed to identify the precise interaction mechanisms. The deposition of Pt on a SnO2 support with or without preferential orientation might result in Pt nanoparticles with a predominant facet. Studies performed on Pt single crystals of the three low index faces showed the variation of ORR activity, with the (100) surface being less active than (110) or (111) by a factor of 4−5 in perchloric acid50 (note that the 110 facet of Pt is not the same as for SnO2). Alternatively, interactions between the terminating species of the SnO2 support and the Pt catalyst might impact ORR activity. The oxidized SnO2 has mainly oxygen terminations, while the reduced SnO2 exposes some Sn2+ species at the surface (see Figure 3). It could be envisioned that supporting Pt on a fully oxidized surface might encourage the formation of Pt oxides, which is likely to reduce ORR activity, while having Sn2+ species available in the surface layer might lead to the predominant adsorption of oxygenated species from water on the support surrounding the Pt nanoparticles, generating more active sites for the ORR on Pt (see e.g. the slightly positively shifted Pt-oxide reduction peak observed on the reduced surface, Figure 8b) or to the formation of PtSn alloys at the support-Pt nanoparticle interface with positive bearings on ORR activity.51 On the basis of the different electronic band structure and d-band centers we determined from the data in Figure 4c, it is reasonable to assume that both aforementioned effects could be a consequence of facet/termination specific electronic interactions between catalyst and support surface. DFT modeling performed by Zyubin et al.52 indicates that depositing Pt on reduced SnO2 surfaces leads to an increase in energy of interaction of the Pt/SnO2 interface. Therefore, stronger electronic interactions between the electronic states of Pt and the surface of the reduced SnO2 might lead to the increased ORR activity compared to Pt on oxidized SnO2.

5. CONCLUSIONS Well-crystallized SnO2 thin films with different orientations were successfully synthesized and catalyzed with Pt nanoparticles using dc magnetron sputtering. Adjusting the sputtering conditions can control film morphology. Preferential orientation along [110] is obtained under oxidizing conditions, while randomly oriented SnO2 films can be grown indirectly by first depositing SnO and subsequent annealing. Hybrid density functional theory calculations suggest that the two-phase equilibrium between SnO and SnO2 during annealing provides conditions where multiple oxidized and reduced facets of SnO2 have almost degenerate surface energies, while fully oxidized (110) is the lowest energy surface under more oxidizing conditions. This rationalizes the experimentally observed growth behavior. XPS confirms the presence of reduced surfaces in the randomly oriented film. The experimentally obtained ratio of 1:1.25 even suggests that the majority of the surface is reduced. A Sn:O ratio of 1:2 can be inferred from XPS for the film grown along [110], confirming fully oxidized (110) termination. Pt nanoparticles deposited by magnetron sputtering on the SnO2 films show different oxygen reduction activity (ORR) in 11300



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acidic media depending on the SnO2 surface properties. The presence of a multioriented, partially reduced SnO2 surfaces leads to a 2-fold increase of the ORR activity of Pt nanoparticles compared to the ORR activity of the same catalyst supported on fully oxidized SnO2 with preferential [110] orientation. This result suggests a surface-dependent strong metal−support interaction (SMSI).

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ASSOCIATED CONTENT

* Supporting Information S

A SEM image of a postannealed SnO2 thin film (S1) and the equilibrium crystal shape of reduced and oxidized SnO2, constructed using Wulff’s method (S2). This material is available free of charge via the Internet at http://pubs.acs.org.

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AUTHOR INFORMATION

Corresponding Authors

*E-mail [email protected]; Tel +44 2380 59 8410 (D.K.). *E-mail [email protected]; Tel +41 56 310 2795 (E.F.). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors thank Umicore GmbH & Co KG and the Competence Center for Electricity and Mobility Switzerland (CCEM) for financial support within the project DuraCat.

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REFERENCES

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