TiO2 Nanosystems by SiO2 ... - ACS Publications

Feb 2, 2010 - Center for AdVanced Modeling and Simulation, Idaho National Laboratory, Idaho Falls, Idaho 83415-3553, and Chemical Sciences DiVision, ...
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J. Phys. Chem. C 2010, 114, 2996–3002

Modification of Au/TiO2 Nanosystems by SiO2 Monolayers: Toward the Control of the Catalyst Activity and Stability Sergey N. Rashkeev,*,† Sheng Dai,‡ and Steven H. Overbury‡ Center for AdVanced Modeling and Simulation, Idaho National Laboratory, Idaho Falls, Idaho 83415-3553, and Chemical Sciences DiVision, Oak Ridge National Laboratory, Oak Ridge, Tennessee 37831 ReceiVed: September 23, 2009; ReVised Manuscript ReceiVed: January 19, 2010

The activity and stability of Au/TiO2 catalysts depend on several different factors such as the anchoring strength of the Au particles at the TiO2 surface, the particle sintering, and the surface mobility of individual gold atoms and/or gold particles. Au/TiO2 catalysts can be made resistant to sintering by atomic layer deposition (ALD) of a layer of SiO2 onto the catalysts. In this study, first-principles density-functional calculations are used to investigate how the stability of Au nanoparticles is modified when a partial monolayer of SiO2 is deposited on a Au/TiO2 catalyst. We find that SiO2 structures deposited on a pure TiO2 substrate exhibit lattice-mismatch instabilities that result in the formation of additional strong anchoring sites for Au atoms/ nanoparticles. An atomic-scale roughness introduced by a partial monolayer of SiO2 can slow the atomic surface diffusion and inhibit Au nanoparticle growth/sintering, in agreement with previous experimental results. 1. Introduction Supported gold nanoparticles (NPs) smaller than 5 nm are extremely active catalysts for low-temperature CO oxidation.1 Different investigations have attributed the enhanced activity of Au catalysts to different NP features (perimeter sites,2,3 lowcoordination atoms,4-9 charge transfer,3,10-16 dynamic structural fluxionality,15,16 nonmetallic properties,17,18 oxidation state19). A recent theoretical study of a large ensemble of TiO2-supported Au nanoparticles20 showed that several nanoscale features collectively result in the observed behavior. Although low coordination is necessary for the Au activity, other features are also essential, such as bonding of either a CO or O2 molecule to a Au nanoparticle that is accompanied by weakening of the Au-Au bonds and facilitates the catalytic reaction. Another key function that needs to be realized for the successful use of a gold catalyst is thermal stability against nanoparticle growth and sintering because the catalytic activity declines with the particle size. Thermal stability depends on several atomic and nanoscale characteristics of the system, such as the anchoring strength of the particle at the oxide substrate, the adsorption energy and migration barrier of an individual gold atom on a nanoparticle or oxide surface, and the barrier for a gold atom jumping from the particle to the substrate. All of these characteristics should depend on the size and shape of the nanoparticle, as well as on the type of oxide substrate and on the presence of impurities and structural defects at the oxide surface. It is known that Au/TiO2 can be highly active in this reaction but suffers from a significant loss of activity at high temperatures (mainly due to the growth and sintering of Au nanoparticles). This feature of gold/titania catalysts can cause significant restrictions in using these systems at elevated temperatures. To improve the thermal stability of gold nanoparticles, several attempts have been made to modify an oxide substrate by * To whom correspondence should be addressed. E-mail: sergey.rashkeev@ inl.gov. † Idaho National Laboratory. ‡ Oak Ridge National Laboratory.

coating it with a thin layer of a second oxide. Such a modification could affect the interaction between the Au particle and the substrate because of the decoration of perimeter sites of the particles by the structures formed by this second oxide. The modified layer could also inhibit the “sliding” of individual metal atoms to the surface of the substrate and reduce the mobility of individual Au atoms at the oxide surface. Examples of such a modification include the modification of SiO2 supports with TiO221-24 or CoOx25-27 coatings, decoration of a TiO2 support with various oxides28-30 or a thin layer of Al2O3 before the applicationi of gold nanoparticles,31 and MgO and BaO modification of Al2O3.32,33 In many cases, an increased thermal stability of the gold particles against agglomeration was observed after the modification. Recently, a Au/TiO2-based catalyst with enhanced thermal stability through postmodification of Au/TiO2 by amorphous SiO2 decoration was reported.29,34 The idea to modify the Au/ TiO2 catalyst by silica was inspired by recent success in the preparation of highly active Au/SiO2 catalysts by a solutionbased technique;35 that is, it seemed quite challenging to combine the most attractive features of both the titania and silica substrates in one catalytic system. It was shown that postmodification (depositing SiO2 after the gold loading) resulted in a more active and stable catalyst than premodification of the TiO2 substrate by a SiO2 thin layer before the gold loading. In this article, we report a theoretical study of the atomic-scale mechanisms and nanoscale features that collectively contribute to the observed behavior of the SiO2-modified Au/TiO2 catalytic system. We find that a SiO2 thin film deposited on a TiO2 substrate exhibits structural instabilities that can result in the appearance of strong anchoring sites for Au atoms/nanoparticles. Pre-existing structural defects at the TiO2 surface also contribute in these processes. Additional SiO2 structures formed on the TiO2 surface promote atomic- or nanoscale texture and roughness of the oxide surface that can significantly reduce the mobility of gold atoms at the surface and inhibit the gold nanoparticle growth and/or sintering.

10.1021/jp9091738  2010 American Chemical Society Published on Web 02/02/2010

Modification of Au/TiO2 Nanosystems by SiO2 Monolayers

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2. Computational Approach To investigate the structural characteristics of a thin SiO2 film grown on a TiO2 substrate, we constructed SiO2 particles with different geometries at the TiO2 surface and optimized their geometries using density-functional theory. The TiO2 substrates were represented by four-layer-thick (110) rutile slabs (with and without surface oxygen vacancies) terminated by nonpolar surfaces. All of the slabs were fully hydrogenated (passivated by hydrogen). Large, periodically repeated supercells (17.75 Å × 12.99 Å × 22.5 Å) were used. The vacuum layer between slabs was >12 Å, which is sufficient to avoid direct interactions between neighboring supercells even in the presence of a SiO2 monolayer and/or small Au nanoparticles on the top of the slab. The relaxed complex oxide SiO2/TiO2 structures were used for accommodation of individual gold atoms or small Au nanoparticles. Gold nanoparticles containing between 5 and 35 atoms on top of the SiO2/TiO2 structures were constructed, as described previously,20 by repeatedly adding single gold atoms to existing structures and allowing the systems to relax again. For simulation of processes on larger nanoparticles, quasi-one-dimensional periodic rod-ike structures containing a well-defined boundary between the nanoparticle and the substrate were also used. We did not perform a systematic search for the lowest-energy structures because (i) the experimental formation of nanoparticles is not necessarily an equilibrium process, and metastable configurations are likely to form; (ii) one can never ensure that the system relaxes to the structure corresponding to the global, not local, energy minimum, as this depends on the choice of the initial structure that, in most cases, is defined by the kinetics of the growing process; and (iii) an ensemble of nanoparticles with diverse local bonding and coordination numbers always provides a wide opportunity for exploring the most favorable conditions for their stability and activity. The calculations were based on the generalized gradient approximation (GGA) for exchange and correlation and plane waves.36 We used the GGA of Perdew, Burke, and Ernzerhof (PBE),37 which gives good results for the chemisorption of molecules at transition-metal surfaces. Ultrasoft scalar relativistic pseudopotentials and the VASP code38 were used. The energy cutoff for the plane-wave basis was set at 400 eV, and all integrations over the Brillouin zone were done using the Monkhorst-Pack scheme with one k point in the relevant irreducible wedge.39 Inclusion of additional k points was found to have minimal effect on the total-energy differences of interest here. The total number of atoms varied between 210 and 320 for different periodic supercells. For each supercell, we relaxed all atoms until the quantum-mechanical force on each atom became smaller than 0.02 eV/Å. We performed both non-spinpolarized and spin-polarized test calculations and did not find a significant difference for the reaction energies and activation barriers of interest. Therefore, to save computer time, most of the calculations did not include spin. Activation barriers were calculated using the nudged-elastic-band method.40 3. Results and Discussion 3.1. Formation of SiO2 Monolayers at TiO2 Surfaces. To understand the mechanism of the formation of SiO2 structures on the TiO2 surface, we positioned individual Si(OH)4 (silicon hydroxide) molecules at a fully hydrogenated (passivated by hydrogen) rutile surface with the (110) orientation. In the bulk rutile crystal, all of the Ti atoms are six-coordinated to O, and all of the O atoms are three-coordinated to Ti.41 When a (110) rutile surface is formed, there appear two different types of undercoordinated atoms: two-fold-coordinated O(2) atoms [that

Figure 1. Fully hydrogenated (110) TiO2 rutile surface. Ti is shown in light gray, O in red, H in white. Undercoordinated surface oxygen atoms O(2) form chains with the [001] orientation. Passivating hydrogen atoms are positioned on top of surface undercoordinated O(2) and surface oxygen vacancies (indicated by red arrows).

Figure 2. Schematics of the formation mechanism of an individual Si-O-Si bridge structure by two reacting OH groups positioned at neighboring Si(OH)4 molecules. Si atoms are shown in dark gray. (a) Interaction between two neighboring Si(OH)4 molecules; light blue ellipse surrounds atoms that later will form a water molecule. (b) Resulting Si-O-Si bridge structure formed after a free water molecule is combined.

form (001)-oriented rows at the surface] and five-foldcoordinated Ti(5) atoms (Figure 1). Both of these undercoordinated atoms can interact with hydrogen. For an isolated H atom, the calculated O(2)-H binding energy (3.6 eV) is much higher than the Ti(5)-H binding energy (0.8 eV). The H binding to a “regular” three-coordinated surface atom O(3) is even weaker (0.4 eV). Therefore, one would expect that, even at room temperature, passivating hydrogen atoms are not kept for a long time anywhere at the surface, except for the positions atop the O(2) atoms. If the initially formed rutile surface contains some number of oxygen vacancies [created by removing surface O(2) atoms], the binding energy between H and an O(2) vacancy is also quite high (3.3 eV). Therefore, the fully hydrogenated rutile surface should contain H atoms on top of each surface undercoordinated O(2) and surface oxygen vacancy (Figure 1). We chose to position silicon hydroxide molecules [Si(OH)4] on the TiO2 surface for illustrative purposes only: in a real-life synthesis of the SiO2-decorated Au/TiO2 system, other Si precursors are used, such as tetramethyl orthosilicate [TMOS, Si(OCH3)4] or butoxysilanol {BOS, [(CH3)3CO]3SiOH} dissolved in ethanol.34 For TMOS and BOS precursors, dehydrolysis similarly leads to ethers or alcohols that might have somewhat different energetics. However, in this work, we did not investigate the synthesis process in detail. Instead, we tried to “visualize” the initialization and growth of a small SiO2 precipitate at the TiO2 surface. Figure 2 shows the formation of an individual Si2O(OH)6 “bridge” cluster that contains two silicon atoms. The Si-O-Si bridgelike structure is formed when two OH groups positioned

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Figure 4. (a) Growth of the SiO2 precipitate along the O(2) chain ([001] direction) as a result of the piling of Si-O-Si bridges along the O(2) chain at the rutile surface. (b) Infinite SiO2 chain precipitate grown at the TiO2 surface.

Figure 3. Schematics of the formation mechanism of SiO2 structures at the fully hydrogenated (110) TiO2 rutile surface without oxygen vacancies from individual Si-O-Si bridge molecules. (a) Deposition of the bridge structure onto the rutile surface. (b) Attachment of the bridge to the surface with the formation of two free water molecules.

on the two neighboring Si(OH)4 molecules react with each other by dehydrolysis to form a free water molecule. Calculations show that such a process is nearly thermoneutral (neither exothermic of endothermic) because one Si-O (plus one O-H) bond in the initial reactants is broken and one Si-O (plus one O-H) bond is reconnected in the final products of the reaction (one bridge molecule and one water molecule). The calculated energy barrier for this reaction is about 0.9 eV. After the formation of the Si-O-Si bridge molecule, the whole group can attach to the TiO2 surface. For this purpose, two OH groups attached to different Si atoms react with the surface H atoms passivating the rutile surface (Figure 3). Again, two other free water molecules are formed as the two Si atoms from the Si2O(OH)6 group bind to two O(2) atoms positioned in different O(2) rows at the TiO2 surface. This reaction is slightly endothermic: calculations show that the two initially existing Si-O (4.7 eV) and O(2)-H (3.6 eV) bonds are “traded” for the finally formed O-H (4.8 eV) and Si-O(2) (3.2 eV) bonds; that is, an additional energy of 0.3 eV is needed for a formation of each free water molecule. The barrier for such a reaction is 0.95 eV, which is quite comparable to the transition-state energy for the formation of the Si2O(OH)6 bridge molecule. Also, an additional bond (with a binding energy of about 0.5 eV) is formed between the bridge oxygen and one of the five-foldcoordinated Ti(5) atoms at the surface (Figure 3b); that is, the reaction of the attachment of the whole bridge molecule to the rutile substrate is weakly endothermic (with the energy loss of 0.1 eV per each Si-O-Si bridge). Here, we proposed that a Si-O-Si bridgelike structure is formed by dehydrolysis from the Si-containing precursor and that the resulting dimer subsequently attaches to the TiO2 surface. We have not yet investigated whether such a scheme is statistically most likely to occur in actual film growth. We just tried to construct the most symmetric SiO2 precipitates and accommodate them on the rutile surface. We think that even such a preliminary consideration still provides an opportunity

to understand some important issues of the growth of the monolayer. The more complete statistical analysis of different structures and growth regimes will be published elsewhere. We do not exclude the possibility that the first steps in SiO2 formation are catalyzed by the surface, that is, individual Si(OH)4 molecules attaching independently to the surface before interacting with each other. When additional Si2O(OH)4 clusters attach along the O(2) row (oriented along the [001] direction at the TiO2 rutile structure), the SiO2 precipitate continues to grow at the TiO2 surface along the [001] direction. On neighboring clusters, OH groups also interact with each other, releasing free water molecules and forming new Si-O-Si bridging bonds along the O(2) row. The OH groups remain attached only to the two edging Si atoms (Figure 4a). Continuing the growth process, one can obtain infinite SiO2 “dimer walls” stretching along the [001] direction (Figure 4b). A collection of parallel walls grown above each O(2) row corresponds to 0.5 monolayer of the SiO2 oxide grown at the defectless (without oxygen vacancies) fully hydrogenated TiO2 rutile structure. Each Si atom in the wall is four-fold-coordinated, the O atoms that connect the nearest Si-O-Si bridges are two-fold-coordinated, and the oxygen atoms that bond to Ti are three-fold-coordinated [they also attach to the Ti(5) atoms]. This structure does not contain hydrogen atoms, as all of the H atoms that were initially present at the hydrogenated rutile surface and the silicon hydroxide molecules formed water molecules and left the surface. Also, the resulting structure does not contain any undercoordinated atoms or other structural defects to which hydrogen atoms could bind strongly (the binding energy of H atom to different sites at a “saturated” Si-O-Si bridge is no higher than 0.5 eV). Therefore, one can expect the growth of any SiO2 structures to stop at this coverage and not continue to form a full monolayer. Here, we assumed that surface attachment of the Si-O-Si bridgelike structures occurs through a symmetric, bidentate structure anchoring to the surface. However, we understand that this is not the only possibility, and monodentate anchoring followed by free rotation of the fragment and subsequent attachment of the other end might be statistically even more favorable. In such a case, the interfacial structures will probably be less ordered. At the present moment, we have not had a chance to investigate all possible ordered and disordered SiO2 structures that can grow on the rutile substrate. Nevertheless,

Modification of Au/TiO2 Nanosystems by SiO2 Monolayers

Figure 5. Full SiO2 monolayer constructed at the rutile surface: (a) metastable symmetric structure, (b) stable asymmetric structure.

our calculations clearly show that, at the initial stage of growth, the SiO2 structures have a tendency to form walls aligned along the [001] substrate direction. To confirm this conclusion, the structural details of the SiO2 monolayer need to be identified through higher-resolution experiments. A hypothetical full SiO2 monolayer might also be constructed at the rutile surface (Figure 5). We call this structure “hypothetical” because we do not discuss here in detail how it was grown. Such a full monolayer could be formed, for example, at a surface that contains a higher number of coordination defects than the ideal rutile surface considered here (see below). A symmetric structure (Figure 5a) is obtained by adding the Si-O-Si bridges between the walls characterizing the half-monolayer structure. However, such a structure is metastable and relaxes to a lesssymmetric configuration (Figure 5b) in which the bond lengths between the three-fold-coordinated bridge oxygen and the two nearest Si atoms become different (1.65 and 1.71 A) compared to the two equal bond lengths of 1.70 Å in the symmetric walls. This means that the SiO2 monolayer has stronger (shorter) and weaker (longer) Si-O bonds; that is, it has a higher tendency toward structural rearrangement and defect formation than a more symmetric half-monolayer collection of the walls. Such a structural reconstruction should affect the anchoring of individual Au atoms and gold clusters at the surface. As stated above, growth of the half-monolayer structure depletes the surface of H; that is, the growth of any SiO2 structures is likely to stop at this coverage. However, we note that, for this model of the TiO2(110) surface, a full monolayer corresponds to 10.5 Si atom/nm2. Previously, it was reported that a single cycle of ALD (using TMOS precursor) yielded a Si/Ti ratio of 0.06 for deposition on Degussa TiO2, and two cycles yielded a slightly higher ratio of 0.11.29 Considering that the TiO2 has a surface area of about 48 m2/g,34 the experimental Si/Ti ratio of 0.060 corresponds to loading of 9.4 Si atom/nm2, approximately a full monolayer. These experiments confirm that a single cycle of ALD will cover the surface in a 90% complete monolayer. The TiO2 powder used in the experiments is expected to contain a much larger number of coordination defects (and is hydrogenated to a higher extent) than the ideal rutile structure used in calculations. The SiO2 structures considered above were constructed at an ideal, defectless rutile structure containing infinite O(2) rows

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Figure 6. Formation of double defect [oxygen vacancy plus threecoordinated Si(3) atom] at the rutile surface with a single O(2) vacancy when (a) an individual Si2O(OH)4 cluster is positioned on top of the oxygen vacancy and (b) a pile of Si-O-Si bridged structures is deposited at the same surface. The oxygen vacancy (pink circle surrounded by a dashed line) and the three-coordinated Si(3) atom are indicated by blue arrows.

aligned with the [001] orientation. These structures take into account some general features that any SiO2 structure positioned at a TiO2 substrate should exhibit (each Si atom is fourcoordinated; each O atom is two- or three-coordinated). In reality, the TiO2 surface contains some number of O(2) vacancies and grain boundaries and might not be fully saturated with H. In addition, the initial Si2O(OH)4 clusters can nucleate in adjacent rows, leading to antiphase boundaries. In such a case, instead of infinite SiO2 walls, one would expect to observe a SiO2 “maze” consisting of SiO2 walls that change their orientation at surface grain boundaries and other defects and that might contain SiOH terminations. Therefore, one can expect that the formation of SiO2 walls at TiO2 creates an additional roughness at the surface and additional mazes and labyrinths that any atom or cluster migrating along the surface should feel. Such a roughness could inevitably inhibit any surface migration process. The formation of oxygen vacancies in the O(2) rows requires much less energy than removing a three-fold-coordinated oxygen atom from the surface. O vacancies can affect the structures of the “decorating” SiO2 layer grown at the TiO2 surface. In particular, when an individual Si2O(OH)4 cluster is positioned on top of an oxygen vacancy in an O(2) row, the Si atom does not bind to the vacancy as it binds to an O(2) atom (Figure 6a). When individual Si2O(OH)4 groups cross-link along the O(2) row to form a SiO2 wall as discussed above, the Si atom positioned on the top of the oxygen vacancy still does not bind to the vacancy. Therefore, each oxygen vacancy positioned in one of the O(2) rows can potentially form a defect with a complex structure that consists of an oxygen vacancy and a three-fold-coordinated dangling Si atom positioned on top of the vacancy (Figure 6b). Such a double defect might be an attractive anchoring site for a metal atom or small nanoparticle. 3.2. Gold Anchoring at Complex SiO2/TiO2 Oxide Surfaces. We considered several possible anchoring sites for a gold atom positioned at 0.5 and 1 ML SiO2-decorated rutile surfaces. The binding energy for Au atom at these sites was found to differ in a wide range. First, we found that the Au atom does not bind to any two-fold- or three-fold-coordinated oxygen atom that belongs to any SiO2 structure positioned above the TiO2

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Figure 7. Schematics of a Au atom anchoring at a complex SiO2/ TiO2 oxide surface. (a) Gold atom attaches to a four-fold-coordinated Si atom positioned at the wall as a fifth atom only (without breaking any of the four existing Si-O bonds). (b) Gold atom attaches to the asymmetric SiO2 monolayer positioned at a defectless TiO2 surface; one of the four Si-O bonds breaks, and the Au atom attaches to the thus-created Si dangling bond (the three other O atoms to which Si is attached are indicated by green arrows). (c) Gold atom binds to a double defect (oxygen vacancy plus Si dangling bond). Gold atoms are shown in yellow.

substrate. This could be expected because all of the bonds of such an oxygen atom are saturated. Second, the interaction between Au and Si atoms depends on two main factors, namely, the coordination of the Si atom (i.e., the presence or absence of dangling bonds), and the ability of the local structure (around the Si atom) to reconstruct, which might also change the Si coordination (dynamic structural fluxionality). As mentioned above, the 0.5 ML SiO2 wall-like structure oriented along the [001] direction is strong and has no tendency to undergo any structural relaxation. Therefore, a Au atom can attach to a fourfold-coordinated Si atom positioned at the wall as a fifth atom only (without breaking any of the four existing Si-O bonds, Figure 7a). This Si-Au bond is not strong (the binding energy is about 0.5 eV only); i.e., the saturated Si atom is not efficient in anchoring gold atoms. An asymmetric SiO2 monolayer constructed on a defectless TiO2 surface (Figure 7b) has a higher anchoring ability because of its tendency for structural reconstruction. Before Au attachment, all of the Si atoms are four-fold-coordinated. However, the asymmetry of the monolayer surfaces leads to two different Si-O bond lengths. When a gold atom is deposited, the weaker (longer) of the Si-O bonds breaks, which creates a Si-Au bond (Figure 7b). The binding energy of such a Si-Au bond is 1.9 eV; that is, such a bond is stronger than the calculated anchoring energy of an individual Au atom at an oxygen vacancy at a TiO2 surface (∼1.2 eV). Such an oxygen vacancy site is considered as the most likely anchoring site for gold atom at a titania surface (see, e.g., ref 20). A pre-existing oxygen vacancy at the TiO2 surface and the formation of a double defect (oxygen vacancy plus Si dangling bond, Figure 6) during SiO2 deposition at the rutile surface can significantly change the gold anchoring energy. Figure 7c shows

Rashkeev et al. the binding of a gold atom to this defect for a 0.5 ML SiO2 structure: a Au atom binds strongly (∼4.5 eV) to both the oxygen vacancy and Si dangling bond. Further growth of a gold nanocluster continues around the anchored Au atom, and the binding energy of each of clusters constructed in such a way is higher than the binding energy of the cluster containing the same number of atoms to a single oxygen vacancy at the titania surface. This means that a gold NP that is grown around such a defect is unlikely to be detached from its anchoring site and travel around the surface. However, one should take into account the possibility that both the defects (long Si-O bonds and double defects) might be passivated by hydrogen in the solution before these defects get a chance to react with individual gold atom or pre-existing Au NP. In this case, it is very unlikely that one would be able to observe any Si-Au bonds shown in Figure 7b,c. Also, one could expect that incomplete SiO2 layers (e.g., 0.3 ML) might also contain different defects and structural features that anchor Au NPs. Another factor that could control the anchoring of a Au atom at the defect and the Au NP growth around the anchored atom is charge exchange between the Au atom and the substrate. We calculated the charge of an anchored Au atom in the three considered cases shown in Figure 7 by integrating the charge density in real space using Bader topological analysis and did not find any significant charge accumulation (higher than 0.25 electron change in absolute value) at the Au atoms. Therefore, we do not think that electric charge is important in controlling the activity of a Au site. The situation may probably change in the presence of an external electric field directed perpendicular to the rutile surface. However, such an analysis goes beyond the scope of this publication. 3.3. Sintering of Au NPs at SiO2/TiO2 Substrates. It has been experimentally shown that postmodification of Au/TiO2 by SiO2 adlayers can significantly inhibit or even mitigate the sintering of gold nanoparticles.29,34 Activity for CO conversion was much less degraded by calcination treatments for the postmodified catalyst than was the unmodified Au/TiO2 catalyst. Both XRD patterns and TEM data indicate that the gold particles do not significantly grow during the calcination process (even at 700 °C) as they do in unmodified Au/TiO2 systems. These data confirm the stabilization effect of the modification of Au/ TiO2 systems by SiO2 adlayers. To understand this stabilization phenomenon, we should mention two main mechanisms of gold nanoparticle sintering: diffusion of individual atoms and migration of clusters as a whole and their coalescence.42 The first mechanism is initiated by transfer of an individual Au atom from a nanoparticle to an oxide substrate and its further migration along the oxide surface. This migration can finish when the migrating atom attaches to another Au NP particle (or the NP of its origin) or the atom gets trapped at some strong binding anchoring site (e.g., at a site shown in Figure 7c). In the latter case, a new nanoparticle may nucleate around this atom. The second mechanism of sintering may be initiated by detaching of the whole nanoparticle away from its initial anchoring site that may be not sufficiently strong to keep the particle. Such a “free” particle migrates at the surface as a whole and can be trapped by other anchoring site or coalesce with another migrating or anchored nanoparticle. As discussed above, if the TiO2 surface contains some number of oxygen vacancies and grain boundaries, one could expect to observe SiO2 mazes consisting of SiO2 walls that change their orientation at surface grain boundaries, antiphase boundaries and other defects. The formation of such SiO2 structures at the

Modification of Au/TiO2 Nanosystems by SiO2 Monolayers

Figure 8. Eleven-atom Au nanoparticle grown at the (110) rutile surface on top of an oxygen vacancy. The NP definitely shows more than one bond with atoms positioned at the substrate.

TiO2 substrate creates an additional roughness at the surface that tends to inhibit surface migration. To elucidate the details of gold nanoparticle growth and sintering at such an atomically “rough” surface, one should perform extensive statistical modeling and simulations that take into account distributions in the size and shape of gold nanoparticles, as well as distributions in the sizes and heights of the surface roughness. The present extensive first-principles calculations show that individual gold atom can transfer from a Au NP to the TiO2 substrate through a multitude of possible pathways. The exact transfer path depends on the position and orientation of the Au NP at the substrate, and the transfer barrier varies in the energy range between 0.8 and 1.5 eV. The binding energy of the transferred gold atom to the NP depends on the coordination of the perimeter site at the NP from which the atom slides to the substrate. Typically, such sliding is a multistep process: the atom decreases step-by-step the number of bonds that connect it to other Au atoms at the NP and binds more and more to undercoordinated atoms and/or oxygen vacancies at the TiO2 substrate (Figure 8). The binding energy of an individual Au atom to an O(2) atom at the substrate is 0.8 eV, that to a Ti(5) atom is -0.4 eV, and that to an O(2) vacancy is -1.2 eV. The best-case scenario for sliding to the substrate (which corresponds to the lowest barrier of 0.8 eV) occurs when the atom can slide directly from the NP to one of the pre-existing oxygen vacancy sites at the substrate. The highest sliding energy barrier (1.5 eV) corresponds to the case when there is no pre-existing O(2) or oxygen vacancy site nearby and the sliding atom should attach to one or more undercoordinated Ti atoms. The migration barrier of a Au atom across a pure TiO2 substrate (without SiO2 mazes) depends on available (dehydrogenated) migration paths. The lowest migration barriers correspond to migration between sites of similar nature: The motion between two neighboring Ti(5) sites has a barrier of 0.2 eV, and the motion between two dehydrogenated O(2) sites (if they are available) has a barrier of 0.25 eV; in both cases, the Au atom moves along the [001] direction. The transition state for a jump from a Ti(5) chain to an O(2) chain costs a bit more energy (∼0.3 eV), whereas the barrier for a jump from a dehydrogenated O(2) to the nearest dehydrogenated O(2) vacancy is 0.15 eV only. Comparing these numbers, one can say that migration preferably occurs along the [001] direction and that the transition energy for jumps between chains is just 0.1 eV lower than the Au-Ti(5) binding energy (0.3 eV vs 0.4 eV); that is, Au atoms still can move between Ti(5) to O(2) chains without detaching from the surface. Also, dehydrogenated O(2) vacancies (if any) always tend to trap Au particles during their motion along the dehydrogenated O(2) chain. The presence or absence of certain migration paths depends mainly on the degree of hydrogenation of the surface and on the concentration of oxygen vacancies and it can vary with temperature or chemical processes at the surface. Anyway, with such low diffusion barriers, a Au atom that was transferred to the substrate could easily diffuse away from the original Au NP.

J. Phys. Chem. C, Vol. 114, No. 7, 2010 3001 These energetics are changed if the gold NP is surrounded by rough SiO2 maze- or labyrinthlike structures. In such a case, the migrating Au atom statistically prefers to stay within some restricted area (e.g., within a maze “cell”) because any jump across the “fences” formed by the SiO2 structures costs extra energy. The barrier for climbing from a defect-free TiO2 surface to a SiO2 wall (shown in Figure 4b) depends on the presence of Si undercoordinated defects at the wall. If all of the Si atoms at the wall are four-coordinated, the moving Au atom should jump to the wall from one of the Ti(5) atoms [the defectless wall does not have any O(2) or O(2) vacancy species as nearest neighbors] and attach as a fifth atom to one of Si atoms at the wall (Figure 7a). Calculations show that the migration path for such a jump through intermediate oxygen atoms is about 0.4 eV, which is comparable to the Ti(5)-Au binding energy. Therefore, the desorption of a Au atom is not statistically less likely than its move to the Si wall. However, the desorbed Au atom is unlikely to return and participate in the formation of Au structures at the surface. Therefore, leaving the cell of the maze is equivalent to the complete loss of the gold atom from the catalytic system. Therefore, we expect that, after some period of moving within the restricted area (maze cell), the Au atom (if it does not go away from the system) will return back to the gold nanoparticle from which it originated (or to another NP within the same cell). Therefore, the sintering will occur within one cell only, and any “global” agglomeration of NPs at the substrate is excluded. Another possible outcome of moving across a wall is trapping of the migrating atom at some defect site associated with the SiO2 barrier (at an undercoordinated Si defect as shown in Figure 7b or at a double defect as shown in Figure 7c). In that case, this atom will be trapped and could initiate the formation of another nanoparticle. Of course, such a trapping can occur only if these defects have not been passivated by hydrogen in the solution before they get a chance to react with gold atoms. Anyway, in both scenarios, the gold atom will not contribute into the Au NPs sintering and/or growth processes. Also, we do not expect that the migration of the NPs as a whole will contribute to the NP growth and sintering processes. As one can see from Figure 8, a Au NP is typically attached to the rutile surface with more than one gold atom; that is, several Au atoms form bonds with the substrate atoms. Therefore, to perform a migration step, one should break and reconnect more than one bond. In addition, Au NPs are attached to at least one oxygen vacancy even before the SiO2 deposition. To start the migration process of the NP, one should move it away from this vacancy, which usually costs more energy than moving the particle away from other undercoordinated surface defects. Also, it is unlikely that a NP could jump across that SiO2 with lower energy barrier than an individual Au atom. Therefore, if the surface diffusion of Au NPs as a whole can contribute to the sintering/growth processes, the sintering should occur within one cell; that is, it also will not contribute into the global agglomeration. Therefore, simple arguments indicate that gold nanoparticles at SiO2/TiO2 substrates have a tendency to maintain their size or even to shrink as a result of nucleation of additional NPs. Both the factors favor their high catalytic activity and are consistent with XRD and TEM experiments that show that gold particles do not significantly grow during calcination process even at very high temperatures. 4. Conclusions In summary, we have presented a comprehensive theoretical/ computational study of the structural stability of Au/TiO2

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nanosystems modified by deposition of an additional atomicthick SiO2 oxide layer. We found that an atomic-scale roughness introduced in such a way can slow the atomic surface diffusion and inhibit Au nanoparticle growth and/or sintering. The resulting SiO2 decoration stabilizes an active Au/TiO2 catalyst mainly because of the high thermal stability of the SiO2 component. Acknowledgment. One of us (S.N.R.) acknowledges the INL Laboratory Directed Research and Development program and the U.S. Department of Energy, Office of Nuclear Energy, under DOE Idaho Operations Office Contract DE-AC07-051D14517 for financial support. Research was sponsored by the Division of Chemical Sciences, Geosciences, and Biosciences, Office of Basic Energy Sciences, U.S. Department of Energy, under Contract DE-AC05-00OR22725 with Oak Ridge National Laboratory, managed and operated by UT-Battelle, LLC. This research used resources of the National Energy Research Scientific Computing Center (NERSC), which is supported by the Office of Science of the U.S. Department of Energy under Contract DE-AC02-05CH11231. It was also supported in part by a grant of computer time from the High Performance Computer Center at Idaho National Laboratory. References and Notes (1) Haruta, M.; Tsubota, S.; Kobayashi, T.; Kageyama, H.; Genet, M. J. B.; Delmon, B. J. Catal. 1993, 144, 175–192. (2) Haruta, M. J. New Mater. Electrochem. Syst. 2004, 7, 163–172. (3) Molina, L. M.; Hammer, B. Phys. ReV. Lett. 2003, 90, 206102. (4) Remediakis, I. N.; Lopez, N.; Norskov, J. K. Angew. Chem. 2005, 44, 1824–1826. (5) Remediakis, I. N.; Lopez, N.; Norskov, J. K. Appl. Catal. A 2005, 291, 13–20. (6) Lopez, N.; Janssens, T. V. W.; Clausen, B. S.; Xu, Y.; Mavrikakis, M.; Bligaard, T.; Nørskov, J. K. J. Catal. 2004, 223, 232–235. (7) Lopez, N.; Nørskov, J. K.; Janssens, T. V. W.; Carlsson, A.; PuigMolina, A.; Clausen, B. S.; Grunwaldt, J.-D. J. Catal. 2004, 225, 86–94. (8) Lopez, N.; Norskov, J. K. J. Am. Chem. Soc. 2002, 124, 11262– 11263. (9) Overbury, S. H.; Schwartz, V.; Mullins, D. R.; Yan, W. F.; Dai, S. J. Catal. 2006, 241, 56–65. (10) Liu, Z. P.; Gong, X. Q.; Kohanoff, J.; Sanchez, C.; Hu, P. Phys. ReV. Lett. 2003, 91, 266102. (11) Liu, Z. P.; Hu, P. Top. Catal. 2004, 28, 71–78. (12) Liu, Z. P.; Hu, P.; Alavi, A. J. Am. Chem. Soc. 2002, 124, 14770– 14779. (13) Molina, L. M.; Rasmussen, M. D.; Hammer, B. J. Chem. Phys. 2004, 120, 7673–7680. (14) Molina, L. M.; Hammer, B. Appl. Catal. A 2005, 291, 21–31.

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