Nature of the Mixed-Oxide Interface in Ceria–Titania Catalysts

Article pubs.acs.org/JPCC

Nature of the Mixed-Oxide Interface in Ceria−Titania Catalysts: Clusters, Chains, and Nanoparticles Aaron C. Johnston-Peck,† Sanjaya D. Senanayake,‡ José J. Plata,§ Shankhamala Kundu,‡ Wenqian Xu,‡ Laura Barrio,∥ Jesús Graciani,§ Javier Fdez. Sanz,§ Rufino M. Navarro,∥ José L. G. Fierro,∥ Eric A. Stach,*,† and José A. Rodriguez*,‡ †

Center for Functional Nanomaterials and ‡Chemistry Department, Brookhaven National Laboratory, Upton, New York 11973-5000, United States § Department of Physical Chemistry, University of Seville, E-41012 Seville, Spain ∥ CSIC-Institute of Catalysis and Petrochemistry, Cantoblanco, E-28049-Madrid, Spain S Supporting Information *

ABSTRACT: The ceria−titania mixed metal oxide is an important component of catalysts active for the production of hydrogen through the water−gas shift reaction (CO + H2O → H2 + CO2) and the photocatalytic splitting of water (H2O + hv → H2 + 0.5O2). We have found that ceria−titania catalysts prepared through wet chemical methods have a unique hierarchal architecture. Atomic resolution imaging by high-angle annular dark field scanning transmission electron microscopy (HAADF STEM) reveals that ceria supported on titania exhibits a range of morphologies. One can clearly identify ceria structures involving clusters, chains, and nanoparticles, which are distributed inhomogeneously on the titania support. These structures are often below the sensitivity limit of techniques such as X-ray diffraction (XRD), which in this case identifies the average particle size of the ceria and titania nanoparticles (via the Debye−Scherer equation) to be 7.5 and 36 nm, respectively. The fluorite-structured ceria grows epitaxially on the anatase-structured titania, and this epitaxial growth influences the morphology of the nanoparticles. The presence of defects in the ceria such as dislocations and surface stepswas routinely observed in HAADF STEM. Density functional theory (DFT) calculations indicate an energetic preference for the formation of O vacancies and the corresponding Ce3+ sites at the ceria−titania interface. Experimental corroboration by soft X-ray absorption spectroscopy (SXAS) does suggest the presence of Ce3+ sites at the interface.

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INTRODUCTION

When considering the CeOx/TiO2 system, the two oxides on their own have been studied considerably and are prototypical lanthanide (CeO2) and transition metal (TiO2) oxides. Catalytically, both offer excellent redox chemical properties including reducibility, stable oxidation states (Ce4+/Ce3+, Ti4+/ Ti3+), and oxygen storage capacity (OSC).14,15 However, both structurally (CeO2−fluorite versus TiO2−anatase lattice) and electronically (Ce-4f versus Ti-3d valence levels), there are substantial differences between these two materials. In this report, a suite of complementary techniques have been used to probe the local and global properties of CeOx/ TiO2. High-resolution transmission electron microscopy (HRTEM), high-angle annular dark field scanning transmission electron microscopy (HAADF STEM), and X-ray diffraction (XRD) were used to characterize the structural and morphological properties of the ceria−titania system. Soft Xray absorption spectroscopy (SXAS) and calculations based on density functional theory (DFT) were used to study the

Metal oxides are an important class of materials used in industrial processes including energy production (H2), storage (batteries), conversion (fuel cells), remediation of environmental pollutants (automotive emissions), fine chemical synthesis (Fischer−Tropsch), electronic materials (semiconductors), and photovoltaics (solar cells).1−3 Of particular interest are mixed metal oxides generated by the deposition of nanoparticles of a given oxide on top of the surface of a second oxide.4−7 For example, structures of M/CeOx/TiO2 (M = Au, Cu, or Pt) produced H2 by the water−gas shift (WGS) reaction (CO + H2O → CO2 + H2) with considerably better activity than either of the individual oxide counterparts.6,7 The origin of this notable enhancement is likely from a combination of structural and electronic properties unique to the oxide nanoparticles and the interfacial region between the oxide and the metal.8 In principle, these systems could expose reactant molecules to unique structures.9−11 Therefore, efforts are needed to understand the fundamental properties of these nanoscale oxides to better interpret their catalytic behavior and to improve the design of new catalysts.12,13 © XXXX American Chemical Society

Received: December 19, 2012 Revised: June 18, 2013

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tion functional proposed by Perdew et al.19 and the projector augmented wave method (PAW)20 as implemented in the Vienna Ab-initio simulation package (VASP) 5.2 code.20,21 A plane-wave cutoff energy of 400 eV was used. The Ti (3s, 3p, 3d, 4s), Ce (4f, 5s, 5p, 5d, 6s), and O (2s, 2p) electrons were treated as valence states, while the remaining electrons were kept frozen as core states. To obtain faster convergence, thermal smearing of one-electron states (kBT = 0.05 eV) was allowed using the Gaussian smearing method to define the partial occupancies. To adequately represent the electronic structure of Ce (in particular the 4f level of the Ce3+ species), we used the GGA+U formalism. The Hubbard U term (on-site repulsion) was added to the plain GGA functional using the rotationally invariant approach proposed by Dudarev et al.,22 in which the Coulomb U and exchange J parameters are combined into a single parameter Ueff = U − J. For Ce we have used a Ueff of 4.5 eV which was self-consistently estimated by Fabris et al.23 using the linear-response approach of Cococcioni and de Gironcoli24 and which is in the range of values usually proposed in the literature (4.5−5.5 eV) for GGA+U calculations.25−34 To represent the 3d states of Ti, a Ueff parameter of 4.5 eV was used, as it reproduces the experimental values of the gap between the Ce3+ 4f and Ti3+ 3d levels observed in the valence photoemission spectra of the Ce/TiO2(110) system.6 The presence of Ce3+ species was indicated by a characteristic 4f peak in the band gap and later confirmed by the magnetization of the Ce atoms (higher than 0.9 electrons) found in the calculations. These Ueff parameters for Ce and Ti have been used successfully for the CeOx/TiO2(110) systems in previous papers.6,7,35 As shown in section 3.2, the microscopy data indicate the growth of CeO2(001) particles on TiO2(112) surface facets. Consequently, the ceria−titania models were built to mimic this relationship. Because the CeO2(001) and TiO2(112) are polar surfaces, the charges of the bottom and top layers of the slab are opposite, and as a result a dipole moment appears. These surfaces, called type 3 according to Tasker classification,36 are not stable and reconstruct. The usual way to model these surfaces is to move half of the charge from the top layer to the bottom layer quenching the dipole moment of the slab.36 Accordingly, in our models, half of the O atoms of the top O layer (belonging to CeO2) were moved to the Ti layer at the bottom of the slab (TiO2). However, because at the ceria− titania interface a net dipole moment still remains, in the present study we have used a dipole correction based on refs 37 and 38 as implemented in the VASP program. To describe phenomena at the interface between a ceria particle and the titania support, a slab was built with five CeO2 layers on top of five TiO2 layers to model the CeO2(001)/ TiO2(112) system. To have converged properties with the number of k-points in the supercell we used a grid of 3 × 3 × 1 k-points in the reciprocal space for the slab model. To avoid the misfit of both surfaces, we did not use the conventional lattice vectors for the TiO2(−112) a and b but (a+b) and (a−b), directions [110] and [1−10]; it could be denoted as (√2 × √2)R45 TiO2(−112) in Wood’s notation. For the same reason, we used a (2 × 2) unit cell for the CeO2(001) surface; as a result, the misfit in the cell parameters is only −1.7%. The cell is not completely orthorhombic; α and β are 90.00°, but γ is 91.09° because the lattice vector a (5.443 Å) is slightly longer than the b vector (5.340 Å) in the TiO2(112) conventional surface unit cell. Therefore, a and b are orthogonal, while the diagonal vectors (a+b) and (a−b) are not. Thus, the misfit in

electronic properties of the ceria−titania interfaces and the propensity for the formation of O vacancies. Our results show a unique mixture of ceria morphologies on titania and the presence of highly stable Ce3+ species at the oxide−oxide interface. These properties have a significant impact on the performance of the mixed-metal oxide in catalytic and photocatalytic processes.

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EXPERIMENTAL AND THEORETICAL METHODS 1. Preparation of Ceria−Titania Catalysts. The ceria− titania powder system was prepared by a wet chemical synthesis.4 The titania powder (Alfa Aesar, anatase, BET = 85 m2/g) was thermally treated at 773 K for 4 h. Following this, the Ce was deposited by wet impregnation of the thermally treated titania powders by using a solution of cerium nitrate (Alfa Aesar, Reacton 99.5%). After the Ce impregnation, the ceria−titania powders were calcined in air at 773 K for 8 h. The nominal Ce content used for the impregnation was a mass fraction of 6%, which corresponds to half of a theoretical monolayer coverage.4 Inductively coupled plasma optical emission spectroscopy (ICP-OES) was run on a Perkin− Elmer Optima 3300 DV instrument to determine the final loading of ceria. Standards of pure CeO2 and TiO2 were also preprepared to obtain fingerprints for all spectroscopy and diffraction measurements. 2. Transmission Electron Microscopy. Scanning transmission electron microscopy images and energy electron loss spectra were collected with a Cs-corrected Hitachi HD-2700C operated at accelerating voltages of 200 and 120 kV with a probe convergence semiangle of 23 mrad. The collection angle for energy electron loss spectroscopy (EELS) was 20 mrad, and the inner collection angle of the annular dark field detector was 53 mrad. A JEOL 2100F operating at 200 kV was used for highresolution transmission electron microscopy (HRTEM). 3. X-ray Diffraction. Powder XRD measurements were performed at beamline X7B (λ = 0.3196 Å) of the National Synchrotron Light Source (NSLS) at Brookhaven National Laboratory. Samples weighing approximately 1 mg were loaded into a Kapton capillary with 0.5 mm ID and measured in transmission geometry. A Perkin-Elmer Amorphous Silicon Detector was used to collect two-dimensional diffraction data, which were processed with the program Fit2D16 to obtain XRD profiles (intensity versus 2θ). Lattice parameters of CeO2 and TiO2 and quantities of each phase were determined by Rietveld refinement using the program GSAS.17,18 4. Soft X-ray Absorption. Near-edge X-ray absorption fine structure measurements (NEXAFS) using soft X-rays from the VUV (vacuum ultraviolet) ring of the NSLS were performed at beamline U4B. This end station is equipped with a custommade partial electron yield (PEY) detector and also a Vortex Si drift detector for partial fluorescence yield (PFY) measurements that were conducted in parallel. The PEY was measured with a single channeltron with a front-end bias of −250 V. All NEXAFS spectra were collected with the photon beam fixed at 55° from the surface normal of the sample. All scans were normalized by the incident beam flux. The cerium M-edge reported here was obtained by scanning of the monochromator from 870 to 930 eV at an energy resolution of 0.2−0.4 eV. Samples were mounted inside the UHV (10−10 Torr) chamber using carbon tape on to the vertical sample manipulator. 5. Density Functional Theory Calculations. The density functional theory (DFT) calculations were made using the generalized gradient approximation (GGA) exchange correlaB

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Figure 1. (a) X-ray diffraction (XRD) pattern from the ceria/titania sample. (b) Ce M-edge NEXAFS of bulk CeO2 and 6 wt % CeOx/TiO2 powder samples obtained via partial electron yield (PEY) and partial fluorescence yield (PFY) measurements.

the γ angle between CeO2(001) (orthogonal) and TiO2(112) (not orthogonal) in this orientation is 1.2%. All of the atoms of the supercell were allowed to fully relax their atomic positions. The slab model is separated from their images by a vacuum of ∼15 Å, considered enough to avoid interaction between them. For building the supercell model we used the optimized lattice parameters for bulk anatase: a = 3.785 Å and c = 9.541 Å.

spectroscopy (ICP-OES) measures a mass fraction of 5.3%. Measurements by XRD are only sensitive to crystalline materials, while ICP-OES will detect ceria regardless of its crystalline state. As the HAADF STEM measurements reveal, there is ceria present without long-range order, and as a result the XRD measurements underestimate the amount of ceria present in the sample. 2. Transmission Electron Microscopy. Transmission electron microscopy (TEM) is used to study the structure of these materials on an atomic level, thereby providing a more local probe than the X-ray techniques utilized. The atomic number sensitive contrast of high-angle annular dark field scanning transmission electron microscopy is particularly useful because the large difference in atomic weight between cerium (Z = 58) and titanium (Z = 22) will make the ceria appear brighter than the titania support and easy to distinguish. However, the signal from a weakly scattering light element like O (Z = 8) will be much less than a heavy element, and as a

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RESULTS 1. X-ray Diffraction. X-ray diffraction measurements of the ceria−titania sample (Figure 1a) detect the presence of anatase TiO2 and cubic fluorite CeO2 with corresponding lattice parameters of {a = b = 3.779 Å, c = 9.503 Å} and {a = 5.404 Å}, respectively. The average crystallite sizes calculated based on XRD peak broadening, accounting for strain, are 7.5 (0.3) nm and 36 (1) nm for ceria and titania, respectively. The Rietveld analysis indicates the mass fraction of ceria is 4.1%. By comparison inductively coupled plasma optical emission C

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Figure 2. HAADF STEM images showing the different morphologies of ceria that can be found in the sample: clusters (a), constrained structures (chains) (b), and nanoparticles (c) are all present (scale bars = 2 nm). Insets in the lower left corner are atomic models illustrating the approximate structure of the different ceria morphologies shown in the HAADF STEM images; for clarity oxygen atoms are not shown. Arrows in (a) and (b) point to the relevant structures.

Figure 3. HAADF STEM images of ceria chains (scales bar = 2 nm). Below each image is a schematic depicting an approximation of the ceria structures observed in the image directly above. The oxygen atoms in the ceria structures were not included in the schematic. The coordinate axes describe the orientation of the titania.

at a lower acceleration voltage of 120 kV which lessens the possibility of knock-on damage. These results suggest these clusters structures are naturally occurring and not an electron beam artifact. In addition to clusters, ceria is also present as morphologies (chains) that are ordered but constrained (less than ∼1 nm) in one or two dimensions (Figure 2b). The chain structures as shown in Figures 2a, 3a, and 3b favor growth along the [101] directions of the titania, and the Ce atoms attach on the Ti atomic columns rather than the troughs in between. Our measurements indicate that the chain structures can be as thin as one or two atomic layers. Even though quantitative STEM techniques were not implemented, it can be estimated, using single ceria clusters as benchmarks, that the out-of-plane dimension can also be very thin (on the order of a few atomic layers). EELS confirmed that these structures and the corresponding changes in contrast were a consequence of the presence of Ce (Supporting Information Figure 2) and not due to other phenomena (e.g., strain) which can affect contrast in HAADF images. Finally, ceria nanoparticles (Figure 2c) were also observed with dimensions upward of ∼8 nm. Figure 4 shows three

result, the atomic resolution HAADF STEM images only reveal the position of the Ce and Ti atoms and not the positions of the O atoms. Therefore, from these images, the coordination between O and Ce atoms in nonperiodic structures (e.g., surfaces, noncrystalline structures) cannot be determined. From S/TEM imaging, the titania support particles are found to be single crystalline and faceted, varying in size from ∼5 to ∼45 nm. The HAADF STEM imaging revealed that a diverse range of ceria morphologies are present within the sample including clusters, chains, and nanoparticles. Representative HAADF STEM images depicting the different morphologies are shown in Figure 2. Figure 2a shows many Ce atoms (HAADF STEM will not visualize the O atoms present), appearing as bright dots, on the surface and at the junction of two titania particles. Electron energy loss spectroscopy (EELS) confirmed that clusters like those shown in Figure 2a were CeOx (see Supporting Information Figure 1). While these clusters can be confidently assigned as ceria, no further insight into their structure can be provided using HAADF STEM or EELS because of technique and sample limitations. It should be noted that care was taken to limit electron beam exposure, and these clusters were also observed when the sample was imaged D

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Figure 4. HAADF STEM and HRTEM images of epitaxial ceria nanoparticles on titania (scale bars = 2 nm). The coordinate axes for describing the orientation of the ceria and titania are in black and white, respectively.

different ceria nanoparticles with cubic fluorite structure have an epitaxial relationship with the titania support. The following relationship describes the epitaxy between the particles and support: (020)CeO2∥(−1−12)TiO2 and [001]CeO2∥[111]TiO2. These particles all have the same relationship with the titania, and any apparent differences are from changes in orientation with respect to the electron beam. While this epitaxial orientation between the ceria and titania is predominant, a small portion of nanoparticles are believed to take on a different relationship with the titania support. Growth of the ceria nanoparticles is favored along specific crystallographic directions of the titania, and as a result the ceria nanoparticles are oblong. When simultaneously observing titania and a ceria along the ⟨111⟩ and ⟨001⟩ zone axes, respectively, as shown in Figure 5a, this anisotropic growth becomes very clear. Even

Figure 6. EELS linescan data collected across the interface of a ceria nanoparticle on titania (scale bar = 2 nm). The ADF signal is that recorded by the annular dark field detector, while the Ce and Ti signals are those recorded by the EELS spectrometer.

measurements indicate that Ti has diffused into the ceria nanoparticle. There may be a small amount of Ce diffusion into the titania, but this result depends on the nature of the particular ceria−titania interface and was not always observed. The result of Ti diffusion into ceria is consistent with the fact that ceria and titania can form solid solutions of the Ce1−xTixO2 (x < 0.2) type.5,34 3. Soft X-ray Absorption Spectroscopy. Near edge X-ray absorption fine structure measurements (NEXAFS) of the Ce L3-edge indicate that a large fraction of the Ce in the system is in the +3 oxidation state.4,5 The Ce L3-edge line-shape from the ceria−titania sample is not equal to that of a CeO2 standard,7 showing that the incorporation of the ceria onto titania surface is changing the charge transfer between O and Ce. Figure 1b shows Ce M-edge spectra collected in partial fluorescence yield (PFY) and partial electron yield (PEY) modes, for bulk ceria and ceria deposited on the titania support. Partial electron yield gives information from the surface with a depth of less than 2 nm because only low-energy electrons, which are near the minimum of the “universal curve” for electron escape depth, are detected. By comparison, partial fluorescence measurements of

Figure 5. HAADF STEM images of ceria nanoparticles with defects including (a) a dislocation and (b) surface defects such as steps (scale bars = 2 nm). The coordinate axes for describing the orientation of the ceria and titania are in black and white, respectively.

though the ceria nanoparticle is terminated by {001} facets on all sides, it has preferentially extended along the [1−10] direction of the titania in comparison to the [11−2] direction. In this same image the presence of a dislocation is identified. Extended crystallographic defects are present in some of the nanoparticles, as well as high energy sites with reduced coordination such as surface steps (Figure 5b). The interface between the ceria nanoparticles and titania support is not chemically sharp. Measurements made by STEM EELS across a ceria−titania interface identify spatially overlapping chemical signatures of the Ti L2,3- and Ce M4,5-edges (Figure 6). It should be noted that the ceria and titania at this interface appear well ordered with respect to each other. Additionally, both the ceria and titania are well orientated with respect to the beam (i.e., on a zone axis), which should minimize any geometrically induced overlapping signals. These E

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core level absorption have an increased sampling depth sensitivity of about 10 nm.5 Hence these two modes of measurement used in parallel can identify differences between the surface and bulk states. In pure ceria only Ce4+ is present, both in PFY and PEY modes.5 In the PEY spectrum of ceria− titania, Ce3+ is present along with Ce4+, indicating that the cerium is partially reduced. In the case of the PFY spectrum, the Ce3+ peak is more intense when compared to the Ce4+ peak. The PFY measurement probes more deeply into the ceria, thereby detecting the oxidation state of the material not only at the surface but also in the core and at the ceria−titania interface. This suggests the cerium exists in a reduced state not at the surface but potentially near the ceria−titania interface. 4. Density Functional Theory Calculations. DFT calculations were performed to analyze the interaction between both oxides. The energy of oxygen vacancy formation, Ev, was calculated as a function of layer position, at the interface (labeled as I in Figure 7) as well as in the ceria (labeled as I+1, I

Figure 8. Computed Ev for the CeOx(001)/TiO2(112) system with a different number of CeO2 layers.

and 2.45 eV, are remarkably close to previous calculations computed for the bare (100) CeO2 surface using a similar setup, 2.27 eV.28 Interestingly, Ev in the TiO2 (I−1) falls between 2.81 and 3.08 eV, significantly lower than the value of 5.14 eV that we have estimated for bulk anatase, while the vacancy energy formation for the ceria−titania interface is computed to be between 2.83 and 2.88 eV, also noticeably lower than the value reported by Nolan et al. for bulk CeO2 of 3.39 eV.39 This behavior can be understood if we analyze the localization of the two cerium ions bearing the 4f electrons that appear when a vacancy is generated. The relaxation of the coordination sphere around the Ce3+ ions creates a small polaron whose stability could be responsible for the stability of the vacancy. In fact, it has been reported that charge distribution around the vacancy can modify the Ev values for the same site by more than 0.5 eV.40 The largest reduction in Ev is found when Ce3+ are arranged as next neighbors to the vacancy on one side and at sites that facilitate their characteristic breathing polaronic expansion, on the other side. Because of the smaller ionic radius of Ti, the polaronic relaxation of Ce3+ in the interface is facilitated. This idea is actually supported by the fact that an electron charge analysis shows that the 4f electrons are located on Ce atoms belonging to the interface even when the vacancy is created in the I+1 and I+2 oxygen planes. Additionally, when the O vacancy is created in the tiatania−oxygen planes, the electrons are found occupying 4f orbitals of cerium atoms located at the interface. As a result the large stabilization of the vacancies around the interface with respect to bulk titania would lead to a diffusion of the vacancies toward the interface and their subsequent accumulation.

Figure 7. CeOx(001)/TiO2(112) system as (a) 5 CeO2 layers on top of 5 TiO2 layers model, (b) 5 CeO2 layers on top of 5 TiO2 layers with solid solution interface model, and (c) CeO2 nanocluster model. I+x labels indicate the different oxygen layer in which the vacancy has been created. Atom color: O = red, Ti = gray, Ce = white.

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DISCUSSION 1. Structure of the Ceria−Titania Catalyst. It has been reported that the electronic and physical properties of ceria nanoparticles are size dependent.41,42 Several studies conclude that the lattice parameter is inversely proportional to size.43,44 While there is some debate as to the cause of this phenomenon,45 it has been speculated that it may be the result of changes in oxidation state of surface species from Ce4+ to Ce3+.44 Here, the morphology of the ceria has a unique hierarchical configuration; the sample has clusters (Figure 2a), chains (Figure 2b), and nanoparticles (Figure 2c). This is

+2, ...) and titania layers (labeled as I−1 and I−2). The values of Ev range between 2.45 and 4.24 eV, corresponding to the I+5 and I−2 planes, respectively. However, these values depend on the thickness of the ceria slab used in the model. In Figure 8, Ev for specific positions is reported as a function of model size. While there is a dependence on the ceria thickness, the most favorable position to create a vacancy is the ceria surface, regardless of model size. The values of Ev at the surface, 2.35 F

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titania can form solid solutions.5,34 Moreover, in STM studies of ceria on a TiO2(110) surface,7 there was evidence for intermixing of metal cations at the oxide−oxide interface. This cation intermixing could be associated with the facile formation of Ce3+ sites detected in the NEXAFS data and was not taken into account in those previous calculations. To reproduce this intermixing in our model we have interchanged two Ce cations with two Ti cations to generate a solid solution interface two layers thick (see Figure 7b). Additionally, to take into account the effects of border and step sites, a new model was developed in which a ceria nanoparticle rather than a slab is adsorbed on a titania surface (Figure 7c). The vacancy formation energies, Ev, for these two new models are included in Table 1. The Ev

significant because the properties of the ceria within the system will have a spectrum of properties reflecting the range of morphologies present. Moreover, these aforementioned reports studied isolated, unsupported ceria; however, in this study the ceria is supported on titania. The titania modifies the growth and properties of the ceria, creating structures that may not be observed otherwise. The lattice mismatch between the cubic fluorite ceria and anatase titania generates stress in the growing ceria and influences the growth process. It was observed that constrained chain-like structures are present over a range of in-plane widths from one or two atomic planes (Figure 3a) to over a nanometer (Figure 3c). The further addition of monomers could facilitate additional growth in plane, spreading the ceria over the surface of the titania. However, large expansive lateral growths of ceria were not observed, rather only island-like nanoparticles. This is because additional in-plane growth comes at a cost of increasing strain energy due to the lattice mismatch between the ceria and titania. As a result, the monomers will attach in the out-of-plane direction to minimize the increase in total added strain energy, thereby limiting in-plane growth and triggering island growth to occur causing the chains to grow into nanoparticles. This growth scheme is depicted in Figure 3d−f where each cartoon illustrates the structure of the ceria present in the HAADF STEM image directly above. The strain generated by the lattice mismatched epitaxial relationship between ceria and titania causes island-like growth, but it also influences the shape of the ceria. The ceria preferentially grows along specific crystallographic directions of the titania and forms oblong structures (Figure 5a). The ceria nanoparticle has preferentially grown along the [110] direction compared to the [112] direction. An epitaxial growth will initially grow in an isotropic manner; however, as additional material is added an instability will occur, and the isotropic particle grows into an oblong structure because an oblong shape better facilitates elastic relaxation of the stress.46 In cases where an anisotropic lattice mismatch exists between the substrate and deposit (as we have here), only one energy minima exists and is along the less strained direction.47 The calculated misfit, using the measured lattice parameter values from the XRD measurements, between {002}CeO2 and {110}TiO2, is −1.1% in comparison to the mismatch between {002}CeO2 and {112}TiO2 which is −13.8%. Therefore, the [110] direction corresponds to the energetic minima, and the particle would be expected to preferentially grow in this direction as we observe. Strain energy can also be reduced by the introduction of dislocations, and while the ceria nanoparticles observed are generally coherent with the titania substrate, some are incoherent and have dislocations, as shown in Figure 5a. In addition to dislocations the nanoparticles also contain highenergy surface defects such as steps and kinks as shown in Figure 5b. Both dislocations and surface defects may act as active sites for catalysis.48,49 2. Nature of the Ceria−Titania Interface. Comparison of the bulk sensitive PFY and surface sensitive PEY measurements (Figure 1b) suggests that the amount of Ce3+ in the inner region, i.e., the interface region, is higher. This finding suggests that Ce3+ centers, and therefore oxygen vacancies, appear to be particularly stabilized in the interface with respect to the surface, which is not consistent with the DFT results reported in Figure 8. However, the STEM EELS measurements (Figure 6) indicate intermixing of metal cations at the oxide−oxide interface. The result is consistent with the fact that ceria and

Table 1. Ev Values for Nanocluster Model and Solid Solution Interfacea cluster

solid solution

site/layers

Ev/eV

tip I+1(in) I+1(surf) I(in) I(border) TiO2 5 layers 2 layers

2.38 2.93 2.29 2.82 1.29 2.48 1.98 1.37

a

In and surf labels indicate if the vacancy is inside of the ceria nanocluster or in one of its faces.

values obtained are similar to those reported above for the twoslab model (Figure 7a) in the tip, inside the nanoparticle, and at the interface; however, it decreases drastically at the border of the nanoparticle and at the solid solution interface. This result gives further support to the higher Ce3+/Ce4+ ratio found with PFY mode in Figure 1b. Notice that this solid solution interface further favors the polaronic expansion of Ce3+ without modifying significantly the coordination properties of metal centers. The theoretical and experimental results indicate that there is a strong stabilization of Ce3+ located at the interface region. To illustrate this point, in Figure 9 the calculated oxygen vacancy

Figure 9. Thermodynamic barrier for polaron-vacancy migration in CeOx(001)/TiO2(112) with five and two CeO2 layers systems with well-separated phases or producing a solid solution in the interface. G

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formation energies (Ev) are plotted as a function of the different oxygen planes for several models. For a perfect interface (without intermixing), the calculations indicate the interface is a local minimum, while the surface layer is the global minimum of the system. Thus, the interface can still trap O vacancies because there is a significant barrier to diffusion from the ceria− titania interface to the surface. By comparison, when the solid solution model is considered, which we believe to be a more accurate physical representation of the system, the surface is no longer the global minima, but the interface layer is now the most preferential site for oxygen vacancies. Oxygen vacancy migration from the surface to the interface would be favorable, leading to a higher amount of Ce3+ at the mixed interface than at the surface. This is in very good agreement with the experimental results. The calculated values strongly depend on the thickness of the CeO2 film, but regardless of the model used the intermixed interface remains the global energy minimum. Therefore, vacancies may more easily migrate depending on the size of the ceria, but the interface remains the most preferred site. The migration of Ti4+ into the ceria lattice at the oxide− oxide interface of ceria−titania facilitates the formation of Ce3+ sites, which can be very useful in catalytic processes. For example, the existence of Ce3+ in M/CeOx/TiO2 (M = Cu, Au, Pt) facilitates the dissociation of water, making these systems excellent catalysts for the production of hydrogen through the water−gas shift reaction (H2O + CO → H2 + CO2).4,7 Furthermore, in ceria−titania, the Ce3+ ions introduce additional states in the band gap, corresponding to the partially occupied 4f levels, and thereby reduce the band gap to about 2.2 eV.5,7 This value is significantly lowered with respect to the clean surfaces of both ceria and titania. UV−vis spectra show that ceria−titania can absorb photons in the visible region, and when combined with Pt it can be used in the photocatalytic splitting of water.7

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CONCLUSIONS

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ASSOCIATED CONTENT

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AUTHOR INFORMATION

Corresponding Author

*E. A. Stach: phone, 631-344-2618; e-mail, [email protected]. J. A. Rodriguez: phone, 631-344-2246; e-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The research carried out at the Center for Functional Nanomaterials and the Chemistry Department of Brookhaven National Laboratory was supported by the U.S. Department of Energy, Office of Basic Energy Sciences, under Contract No. DE-AC02-98CH10886. The theoretical studies were funded by the Ministry of Economy and Competitiveness (Spain, grants MAT2012-31526 and CSD2008-0023) and EU FEDER. Computational resources were provided by the Barcelona Supercomputing Center/Centro Nacional de Supercomputación (Spain).

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REFERENCES

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In summary, we have studied the structural, morphological, and electronic properties of a ceria−titania system formed by the dispersion of ceria nanostructures on anatase powders. HAADF STEM imaging reveals that ceria exists in a hierarchical structure with clusters, chains, and nanoparticles on the titania support. Furthermore, it was found that the growth of the ceria is influenced by the strain originating from the lattice mismatch between the ceria and titania. This is significant because a component (ceria) of the catalyst with size- and straindependent properties is found in many different states as well as containing many high-energy defect sites. The results of NEXAFS and DFT calculations indicate that the formation of a CeO2−TiO2 interface substantially favors the formation of Ce3+ centers. Therefore, the deposition of ceria on titania leads to materials which have novel structural and electronic properties and are quite attractive for applications in catalytic processes.

S Supporting Information *

TEM images of the titania particles and EELS data from clusters and wires. This material is available free of charge via the Internet at http://pubs.acs.org. H

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