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Morphology, Stoichiometry, and Interface Structure of CeO2 Ultrathin Films on Pt(111) P. Luches,*,† F. Pagliuca,†,‡ and S. Valeri†,‡ † ‡
CNRIstituto Nanoscienze, Centro S3, Via G. Campi 213/a, 41125 Modena, Italy Dipartimento di Fisica, Universita di Modena e Reggio Emilia, Via G. Campi 213/a, 41125 Modena, Italy ABSTRACT: Studies of model systems based on cerium oxide are important to improve current understanding of the properties of ceria-based materials, which find wide application based on the ability of cerium oxide to store, release, and transport oxygen. We report a study of CeO2 ultrathin films grown on the Pt(111) surface by reactive deposition of Ce using molecular or atomic oxygen as the oxidizing gas. High-temperature treatments in O2 allowed us to obtain epitaxial structures with a very good quality in terms of morphology, stoichiometry, and structure. The cerium oxide films have a very flat morphology with terraces several tens of nanometers wide. The stoichiometry of the films is mainly CeO2, and the concentration of Ce3þ ions in the film can be reversibly increased by temperature treatments. We propose that the Pt substrate oxidation has a determinant role for the epitaxial stabilization of ceria films.
1. INTRODUCTION Cerium oxide is a material with unique physical and chemical properties, which make it appealing for the application in different fields. It is used in catalysis1 and fuel cells,2 due to its ability to store, transport, and release oxygen depending on the ambient conditions. Other applications include its use as a high dielectric constant material which can be grown epitaxially on Si3 or as a buffer layer for the growth of high-temperature superconducting coatings.4 The microscopic origin of the peculiar properties of cerium oxide is still under debate in a number of experimental and theoretical studies, which are focused on the understanding of the origins of reducibility and on the optimization of the properties of this material. For example, the study of formation and interaction of defects in ceria and the relative electronic structure has reached a good level of understanding,514 which is of great help in studies concerning the identification of the active sites in the separate steps of important reactions, such as, for example, the water gas shift reaction.15,16 When insulating materials such as cerium oxide are concerned, basic studies of bulk crystals are possible,6,8,10,11 although they are sometimes difficult due to charging problems. The use of systems made of ultrathin films/nanostructures on metal substrates has a twofold advantage. The first one is the increase of the charge carriers mobility, which allows for a more straightforward application of the analysis techniques involving a charge transfer to/from the probed sample. Furthermore it allows one to study also the interaction with the metal substrate, which is also a very relevant aspect. Most catalysts based on cerium oxide are in fact made of oxide-supported metal nanoparticles, in which the oxide, besides having a role for itself, also activates the metal. Understanding the metal/oxide interface in its details is therefore a very r 2011 American Chemical Society
important issue also in this respect. Platinum represents a very interesting substrate since the material is employed in combination with CeO2 in catalysts of different kinds.1,16,17 Early studies of the model ceria/Pt system were performed on postoxidized metallic Ce layers or CePt surface alloys on Pt(111) and resulted in films with relatively high Ce3þ concentrations (approximately 20%)1820 and relatively small-sized nanostructures.2123 Very recently Grinter et al. used both reactive deposition and postoxidation of a CePt surface alloy for the growth of cerium oxide ultrathin films on Pt(111) and studied the morphology and the atomic-scale defectivity of the obtained surfaces.14 However, several important aspects of model systems made of ceria layers grown on metal surfaces with hexagonal symmetry are still unclear. For example, the stabilization mechanisms for the obtained epitaxial orientation of the (111) ceria phase with sides parallel to the substrate main surface crystallographic directions in spite of the considerable lattice mismatch (approximately 40% for most of the investigated substrates) and the detailed structure of the ceria/ metal interface still deserve investigation. In this work ceria nanostructures with very flat and wide terraces and low concentrations of defects are produced and studied, with emphasis on the role of interfacial effects on the structure and morphology.
2. EXPERIMENTAL SECTION The experimental apparatus used for the preparation and characterization of the samples consists of three connected Received: February 3, 2011 Revised: April 21, 2011 Published: May 06, 2011 10718
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(10. Unless otherwise specified, during the annealing steps the sample was kept at the desired temperature for 15 min. XPS measurements were performed at normal emission using Al KR photons as the exciting probe. STM measurements were performed at RT in constant current mode using electrochemically etched tungsten tips electrically grounded and a positive sample bias. The STM images have been processed using the WSXM software.24 The analysis of Ce 3d photoemission spectra is quite complex, the spectra being, in general, given by the contribution of several different initial and final states for each of the two possible Ce oxidation states. To obtain relative concentrations of Ce3þ and Ce4þ in the samples, we fitted Ce 3d spectra following a wellestablished procedure introduced by Skala and co-workers.25,26 As shown in Figure 1, 10 peaks and a Shirley-type background were used to fit the spectra. Six of them are associated to three Ce4þ spinorbit split doublets with three different final states and four are associated to two doublets in the Ce3þ ions.27 The precise assignment of these peaks in terms of inter- and intraatomic many-body effects has been a matter of debate for a long time, and it is still under discussion.2830 Figure 1. Ce 3d spectra of a 7.5 ML ceria sample before (a) and after (b) annealing in O2 at 1040 K and of a 0.45 ML ceria sample before (c) and after (d) the same treatment. The fits of the spectra are shown as solid lines. An example of the different components and background used for the fitting is also shown for spectrum d.
ultrahigh vacuum (UHV) chambers: one allows for the growth of cerium oxide layers, the second one is equipped with facilities for substrate preparation and for Auger electron spectroscopy (AES), X-ray photoelectron spectroscopy (XPS), and low-energy electron diffraction (LEED) analysis, and the third one contains a room temperature (RT) scanning tunneling microscope (STM). The base pressure in the UHV apparatus is 2 1010 Torr. The substrate used for the growth of cerium oxide is a Pt(111) single crystal, which was prepared by repeated cycles of Arþ sputtering (1 keV, 1 μA) and annealing (1040 K) in UHV until the level of contaminants was below the sensitivity of AES and the LEED pattern showed sharp spots. Ceria was grown by reactive evaporation of metallic cerium using either molecular or atomic oxygen as the oxidizing gas. Cerium was evaporated from an electron beam evaporator with a rate, measured by a quartz microbalance, of approximately 0.2 Å/min. O2 was dosed through a nozzle in O2 background pressure of 1 107 Torr. Atomic oxygen was supplied through a thermal cracker operated in an oxygen background pressure of 1 107 Torr and in conditions nominally giving approximately 70% dissociation of O2 molecules. The Pt substrate was kept at RT during ceria growth. The total amount of evaporated ceria in each sample was estimated using the evaporation rate and cross checked by Pt 4f XPS intensity attenuation and STM, the latter only in the case of samples with noncomplete substrate coverage. In the following the amount of deposited ceria is expressed in terms of ceria monolayers (ML), defined as the thickness of an OCeO trilayer in the (111)-oriented fluorite structure, corresponding to a thickness of 3.12 Å. The ceria overlayers were annealed at different temperatures, up to 1040 K, either in UHV or in an O2 pressure of 1 107 Torr. The sample temperatures were measured by a thermocouple attached to the sample holder very close to the sample position and are assumed to be accurate within approximately
3. RESULTS 3.1. Stoichiometry. The stoichiometry of cerium oxide deposited on Pt at the different coverages and its modification with the different treatments was investigated by XPS. In particular, the concentration of Ce3þ and Ce4þ in each sample has been obtained by fitting the Ce 3d spectrum with the procedure described in the Experimental Section and by calculating the ratio between the total area of the peaks related to the single ionic species and the total area of the considered spectrum. This evaluation does not include possible attenuation effects related to the specific location of the different ionic species. The fitting of XPS spectra measured at grazing emission on selected samples indicates that there is a preferential concentration of Ce3þ sites at the surface of the samples; therefore, the evaluated ratio represents an overestimate of the real Ce3þ concentration in the samples. Figure 1 shows the Ce 3d XPS spectra of two representative samples, 7.5 and 0.45 ML thick, grown using molecular oxygen as the oxidizing gas, and the relative fits before and after annealing in O2. In the as-grown samples (spectra a and c) the concentration of Ce3þ is very large for the thinner sample (CCe3þ = 55.7%) and relatively small for the thicker one (CCe3þ = 3.8%). For both samples the obtained stoichiometry has been found to be stable with annealing in O2 up to temperatures of approximately 770 K. An annealing to 1040 K in O2, instead, gives rise to a significant change in the spectrum of the thin sample, whereas the thick one is substantially unaltered (Figure 1, spectra b and d). The fit of spectrum d in Figure 1 indicates that the Ce3þ concentration in the 0.45 ML sample is decreased from 55.7% to 12.6% after high-temperature annealing in O2, but it is still larger than the one of the 7.5 ML sample. The O 1s XPS spectra of the 0.45 and 7.5 ML cerium oxide samples before and after O2 annealing are shown in Figure 2, in comparison with a spectrum acquired on the clean Pt substrate annealed in O2. The presence of the Pt 4p3/2 photoemission peak at 519 eV (not shown in Figure 2) gives origin to the observed sloped background in the O 1s binding energy region of the O2-annealed Pt and the 0.45 ML ceria samples. Before annealing in O2 the O 1s spectrum of the 0.45 ML sample shows a main peak above 530 eV binding energy and a lower intensity peak 10719
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Figure 3. Ce3þ concentration at different ceria nominal thicknesses after annealing in O2 at 1040 K obtained using atomic (red triangles) or molecular (blue dots) oxygen during the growth.
Figure 2. O 1s spectra of a 7.5 ML ceria sample before and after annealing in O2 at 1040 K and of a 0.45 ML ceria sample before and after the same treatment. The O 1s spectrum measured from the Pt(111) substrate after annealing in O2 at 1040 K is also shown.
around 529 eV. After O2 annealing a single broad component is present, probably originating from the contribution of the two peaks present before the annealing with a similar relative weight. Also for the 7.5 ML sample, the high binding energy shoulder of the O 1s peak above 530 eV, evident in the as-grown sample spectrum, is significantly decreased and the 529 eV component is increased in intensity after high-temperature annealing in O2. The Pt sample, annealed in oxygen in the same conditions of the cerium oxide samples, shows a broad O1s peak centered around 531 eV (Figure 2, bottom curve). The effect of the use of atomic oxygen for ceria growth has also been tested, finding also in this case that the low-coverage structures contain a high concentration of Ce3þ sites, which can be decreased to values around 10% after annealing in O2 at 1040 K. As shown in Figure 3, after annealing in O2, the Ce3þ concentration in films grown in atomic oxygen is systematically slightly lower than in films grown using atomic oxygen at all of the investigated coverages. We did not observe any X-ray irradiation-induced reduction of the samples after annealing in O2, at least for irradiation times comparable with the ones used for our measurements (below 1 h at 14 keV, 20 mA). This evidence is consistent with the results obtained in the literature, where ceria samples in the form of compressed powders, hence with a much larger surface area than the ones used for this study, were found to be reduced only after the exposure to similar irradiation doses and energies for times significantly longer than 1 h.35 We investigated the thermal stability of a sample of intermediate thickness, namely, 2.3 ML, after the optimization of its stoichiometry by annealing in O2 at 1040 K. The Ce 3d spectrum measured before UHV annealing is shown in Figure 4, spectrum a. The Ce4þ concentration, evaluated from its fit, is 88%. The film stoichiometry remains approximately stable after 5 min annealing
Figure 4. Top: Ce 3d spectra of a 2.3 ML ceria sample after annealing in O2 at 1040 K (a), after 1 h of annealing in UHV at 1040 K (b), and after reannealing in O2 at 1040 K (c) and relative fits. The data are shown as black dots, and the fit is the red solid line. Bottom: Ce4þ concentration after UHV annealing for 5 min at the indicated temperatures of the 2.3 ML ceria sample (filled red dots). The Ce4þ concentration after 1 h of UHV annealing at 1040 K (empty red dot) and after reannealing in O2 at 1040 K (blue triangle) is also reported.
steps at temperatures up to 770 K, whereas at higher temperatures the Ce4þ concentration slightly decreases reaching 70% at 1040 K (Figure 4, bottom panel). At this temperature a longer annealing time drastically reduces the sample. The stable state at this 10720
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Figure 5. LEED patterns at a primary beam energy of 80 eV of the clean Pt(111) substrate (a) and of ceria samples of different thicknesses after annealing in O2 at 1040 K: (b) 0.7 (inset, zoom around one of the ceria-related spots), (c) 2, (d) 3.4, (e) 11.1, and (f) 2.3 ML after 1 h of annealing in UHV at 1040 K. (g) Evolution of the ceria inplane lattice parameter with ceria thickness; the bulk value is shown as a solid line.
temperature, reached after approximately 1 h, has a Ce4þ concentration of 34%. The Ce 3d spectrum of this phase and its fit are shown in Figure 4, spectrum b. The Pt 4f peak area (not shown) was unaltered after the UHV annealing process. After reduction by UHV annealing, the sample could be brought back to the original stoichiometry by annealing in O2 at 1040 K (Figure 4, spectrum c). 3.2. Structure. The structure of the samples was investigated by LEED. Before annealing in O2 the ceria samples did not show any LEED pattern (only the Pt spots were visible at low coverage), regardless of the oxidizing gas (O2 or atomic oxygen) used during growth. After annealing in O2 we did not observe significant differences in the LEED patterns of samples of similar
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coverage grown in atomic or molecular oxygen; therefore, in the following we shall not explicitly indicate the oxidizing gas used. Figure 5 reports the LEED patterns, obtained using a primary beam energy of 80 eV, for the clean Pt substrate and for ceria samples of different thicknesses after annealing in O2 (Figure 5ae). The clean substrate shows a pattern with hexagonal symmetry typical of an unreconstructed (111) fcc surface. At low ceria coverages (Figure 5, parts b and c) the Ptrelated spots are still visible, and ceria-related spots appear as a hexagonal lattice with a smaller reciprocal lattice spacing compared to the Pt one and with the in-plane symmetry directions aligned to the Pt ones. The ceria in-plane lattice parameter evaluated from the LEED patterns at different coverages appears 1% contracted with respect to the CeO2 bulk lattice parameter (Figure 5g), and a slight relaxation is observed at larger coverages. Figure 5g is limited to the samples in which a relative measurement of the ceria surface lattice parameter with respect to the Pt one was possible, i.e., the ones for which the Pt spots were still visible in the LEED pattern. At very low coverage two satellites are evident around the ceria-related spots (Figure 5b, and relative inset). As the coverage increases the satellites disappear and merge into single spots, which become progressively broader as the film thickness is increased. Figure 5f shows the LEED pattern of the 2.3 ML sample after UHV annealing at 1040 K for 1 h. Satellites with hexagonal symmetry appear around the ceria-related spots forming a 3 3 superstructure plus additional spots in a quasi 6 6 arrangement. 3.3. Morphology. In order to have a more complete understanding of the morphology of the investigated system, we first describe the morphology of the Pt substrate after exposure to molecular oxygen in the conditions used for ceria growth and postgrowth treatments. The STM-measured topography of the Pt substrate surface (not shown) shows large flat terraces, hundreds of nanometers wide, separated by monatomic steps. After RT exposure of the substrate to oxygen, either molecular or atomic, at the same doses used for ceria growth, i.e., 510 min at 1 107 Torr, no change in the Pt surface morphology was detected. The Pt surface morphology was significantly modified by annealing at 1040 K in 1 107 Torr molecular oxygen pressure. As shown in Figure 6a, a high density of structures appears on the terraces, with some features appearing as holes, some others as slight protrusions. In all cases the structures show straight edges, forming 60 or 120 angles with each other, and have a size which ranges from a few nanometers to a few tens of nanometers. The terrace steps, which were straight before O2 annealing on scales of hundreds of nanometers, appear corrugated after annealing in O2. A close inspection of the edges shows that the erosion exposes straight sections of terrace edges a few tens of nanometers long, which form 60 or 120 angles also in this case. It is interesting to note that a higher density of protrusions is present along straight lines, probably where the Pt terrace edges were located before O2 annealing. A bias-dependent depth/height of the observed holes/protrusions was observed. An example of this effect is given in Figure 6, showing the same sample area measured with a 0.5 V bias (Figure 6a) and using a 2 V bias (Figure 6b). Most of the features which appeared as holes in Figure 6a appear as protrusions in Figure 6b. As shown by the profiles in Figure 6, parts c and d, the height of the protrusions varies from 1 to 34 Å, changing the bias from 0.5 to 2 V. Figure 7 reports the topography of ceria overlayers at low coverages. After deposition at RT ceria has granular surface, with grains of a few nanometers in size and a few angstroms in height, 10721
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Figure 6. STM images of the Pt(111) surface annealed in O2 at 1040 K measured at 0.5 V, 0.1 nA (a) and at 2 V, 0.1 nA (b) and relative height profiles along the blue lines (c and d).
Figure 7. STM images of (a) a 2 ML ceria sample as grown (2 V, 0.15 nA); (b) a 0.2 ML ceria sample annealed in O2 at 1040 K (1 V, 0.2 nA); (c) same as panel b (0.5 V, 0.2 nA); (d) sketch of the ceria/Pt system at low coverages with PtO2 islands of different thicknesses either below the ceria islands or on the bare substrate; (e) 0.7 ML ceria sample annealed in O2 at 1040 K measured at two different biases (left 0.3 V, 0.1 nA, right 1 V, 0.1 nA); (f) atomically resolved images measured on the ceria islands of the sample shown in panels (b) and (c) (0.8 V, 0.2 nA); the different kinds of defects are evidenced: oxygen vacancy clusters (triangle, zoom in panel g), surface oxygen vacancies (circle, zoom in panel h), and subsurface oxygen vacancies (dashed circle, zoom in panel i). 10722
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Figure 8. STM images of ceria samples of different thicknesses after annealing in O2 at 1040 K: (a and b) 2.85 ML (2 V, 0.1 nA); (c) 3.4 ML (2 V, 0.1 nA); (df) 9.3 ML (4 V, 0.1 nA).
as shown in Figure 7a. The grain size and height do not change significantly with the nominal amount of cerium oxide deposited and with the use of molecular or atomic oxygen as the oxidizing gas during the growth. The obtained morphology is unaltered by annealing up to 770 K, whereas it changes significantly after annealing at 1040 K in molecular oxygen. Figure 7b shows the morphology of a 0.2 ML film after annealing at 1040 K in O2. The flat Pt terraces, which show corrugated edges similar to those observed after annealing in O2 of the bare Pt surface, are covered by islands of two different kinds. The smaller islands, which nucleate both on the terraces and at step edges, have an average lateral size of approximately 10 nm, and they appear approximately 3 Å high at 1 V bias. Their shape is mostly triangular, with the sides forming 60 angles with each other, but 120 kinks are also observed (see Figure 7c). The larger islands in Figure 7, parts b and c, have a lateral size of several tens of nanometers and an apparent height of approximately 6 Å, which corresponds to 2 ceria ML. They are atomically flat, and they preferentially nucleate at step edges. The sides of these structures are mainly straight, and they mostly form 120 angles with each other, whereas angles of 60 are very seldom observed. A close inspection of the larger islands (Figure 7c) shows small areas with a different contrast, as large and as dense as the smaller islands, often located at some of their kinks and sides but also within the larger islands. Both smaller and larger islands change their apparent height with bias: as shown in Figure 7e the smaller islands appear as holes at 0.3 V and as protrusions at 1 V, whereas
the larger islands appear as protrusions, although with a biasdependent apparent height, at all of the bias values used. The measured height value saturates at 67 Å for the larger islands and at 3 Å for the smaller islands for biases above 1 V. Also the small features within the larger islands change their contrast with bias, as shown in Figure 7e (top left and bottom right islands). Atomically resolved STM images (see Figure 7f) could be obtained on the larger islands using a positive sample bias, thus probing empty Ce 4f states, while the quality of small-scale images acquired with negative biases was systematically worse. The atoms are arranged in a hexagonal lattice with an average interatomic distance of 3.53.6 Å, a value contracted with respect to the one of bulk ceria. Different kinds of atomic-scale defects are evident. Defects with a triangular symmetry, point defects, and Y-shaped defects are marked in Figure 7f with a triangle, a dashed circle, and a solid circle, respectively. Zooms of each type of defect are shown in Figure 7gi, respectively. On the smaller islands it was not possible to obtain atomically resolved images. At larger cerium oxide deposited amounts, the islands tend to coalesce into structured formations. Bias values higher than 2 V are necessary to obtain good quality images of thicker samples. After the deposition of 2.8 ML of ceria (Figure 8, parts a and b), the surface is covered by a highly structured layer with several tens of nanometers wide formations, with an apparent height of approximately 6 Å, although also some 9 Å high islands are evident, together with higher structures. The ceria nanostructures 10723
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The Journal of Physical Chemistry C have straight borders forming mainly angles of 120 with each other. The surfaces are atomically flat with linear and extended defects (Figure 8b). By increasing the coverage to 3.4 ML (Figure 8c) almost closed films are obtained, although still highly structured and with holes which expose the underlying layers or even the Pt substrate. Four different contrast layers can be distinguished in Figure 8c. The darkest one probably corresponds to the Pt substrate. The first ceria layer appears 6 Å high, whereas the second and the third appear 9 and 12 Å high approximately. The edges of the holes form angles of 120, and a high density of linear defects, several tens of nanometers long, is present also at this coverage. In order to have good quality images of thicker films a 4 V bias was necessary. At a coverage of 9.3 ML (Figure 8df) most of the surface is completely covered by an almost continuous film with protrusions, holes, and linear defects. The stepped structure of the Pt substrate is still visible in large-scale images (Figure 8d). The holes are a few nanometers wide and more than 10 Å deep. The surface of the film is covered by islands, several tens of nanometers wide, which protrude by 3 Å from the underlying film. The islands edges are relatively straight, although slightly less straight than at lower coverages. The islands appear triangular at a first glance, but a closer inspection (Figure 8, parts e and f) evidences several 120 kinks separated by short (a few nanometers) edges. Several linear defects are present on the surface (Figure 8f). The large-scale morphology of a 2.3 ML film before and after 1 h UHV annealing (image not shown) is substantially unaltered: the ceria structures have very similar lateral sizes, heights, and shapes.
4. DISCUSSION Several aspects of the results exposed in the previous section can be understood by a cross comparison of the information obtained by the different techniques. The lower relative concentration of Ce3þ in as-grown samples compared to samples annealed in O2 (Figure 1) can be understood in terms of morphology. Figure 7a shows that after RT growth the formations obtained have a highly structured surface, with a large number of defects and low-coordination sites in which Ce is more likely to be in the 3þ oxidation state. After annealing in O2, instead, the structures coalesce into larger islands with flat surfaces and a much lower density of defects. At the lowest coverages investigated the Ce3þ concentration after annealing in O2, still larger than 10%, is due to a relatively large weight of defects and low-coordination sites, such as step edges, expected to be more reducible than fully coordinated sites. The Ce3þ concentration in the films becomes really low (2%) already at a few monolayers coverage, indicating that the ultrathin films are highly stoichiometric. It has to be noted that the values obtained for the relative concentration of the single ionic species are affected by non-negligible intrinsic errors (of the order of a few percent), due to the complex structure of Ce 3d spectra which contain a number of low-intensity many-body final states as satellites of the most intense structures,30 which are very difficult to be correctly deconvolved, due to their superposition with the higher intensity peaks of different origin. The evolution of the film stoichiometry with annealing in O2 and thickness is reflected also in the shape of the O 1s peaks (Figure 2). These spectra contain at least two peaks, which change their relative weight with annealing in O2 and with thickness. Taking into account also some literature results analyzing O 1s
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spectra of ceria ultrathin films,18,3134 we assign the 529 eV peak to O2 species bound to Ce4þ and the ∼1 eV higher binding energy peak to oxygen species either bound to Ce3þ in the 2 oxidation state or with a slightly lower valence in lowcoordination or defect sites. Also, oxygen bound to Pt atoms (see below) has a binding energy of approximately 530 eV as shown in Figure 2. Absorbed species, such as, for example, OH or CO, have a significantly larger O 1s binding energy.33,34 An inspection of the measured spectra allows us to exclude a relevant presence of foreign adsorbed species on the samples. The decrease in the ratio of the peak related to oxygen ions bound to Ce3þ atoms and the one related to oxygen ions bound to Ce4þ with increasing thickness and/or after annealing in O2 is in agreement with the observed changes in morphology. In the literature atomically resolved images of ceria surfaces have been acquired using either positive or negative biases.6,8,1214,21,36 The topography of the atomic-scale defects observed in this study can be compared to the one of previous works using bulk single crystals8 or ceria films grown on Ru13 in which the empty states are probed. On this basis we assign the point defect marked as a circle in Figure 7f (zoom in Figure 7h) to single surface oxygen vacancies, the Y-shaped defect marked as a dashed circle (zoom in Figure 7i) to subsurface oxygen vacancies, and the defect marked as a triangle (zoom in Figure 7g) to oxygen vacancy clusters. The predominance of 120 kinks in ceria islands at low coverages (Figures 7 and 8) was observed also on islands at bulk ceria single-crystals surfaces and ascribed to the formation of step edges exposing different crystal facets.37 Ceria layers grown on other metal substrates, such as Cu,38 Rh,36 and Ru,13 also show a similar morphology in terms of step edges. Assuming that the straight edges are oriented along the high-symmetry Æ110æ surface directions, the occurrence of 120 kinks instead of 60 ones is compatible with ceria islands with edges exposing adjacent (111) and (100) facets, the only ones which may form a 120 kink and simultaneously form angles larger than 90 with the surface. An analysis of the shape of the structures shows that the sides of the islands have a similar extension. This evidence implies that, although on the bulk the (111) face is expectedly the most stable one, low-dimensional and substrate-induced effects make the stability of the (100) facet comparable to the (111) one at low coverages. Furthermore, the prevalence of 120 kinks indicates that kinks between the two different facets are more stable than kinks at 60 between steps exposing the same facets. At larger coverages a lower abundance of one of the two edges, possibly the less stable (100) face, is observed (see Figure 8df). This can be explained by assuming a compensating role of the metal substrate, which stabilizes less stable orientations at low coverages, as found for other systems in which polar step edges are observed.39 The linear defects appearing at larger coverages (Figure 8) can be assigned to domain boundaries between islands with edges exposing different facets which get into contact. The appearance of an O 1s peak in the XPS spectra (Figure 2), after annealing the Pt(111) surface in O2 in the same conditions used for ceria film annealing, indicates an oxygen uptake by the substrate. Although Pt exposure to oxygen has been widely studied in different experimental conditions (see, for example, ref 40 and references therein), only a recent study investigates the morphology of Pt exposed to doses and temperatures high enough to oxidize Pt.41 In conditions similar to the ones used in our study for ceria annealing in O2, the authors observe the formation of triangular nanoclusters nucleated at Pt step edges, 10724
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The Journal of Physical Chemistry C which are ascribed to an R-PtO2 phase with a 30 in-plane rotation with respect to Pt, to minimize the epitaxial strain.41 These structures are similar in size and shape to the ones we observe (Figure 6). We observe a high density of structures also on the Pt terraces and a higher degree of erosion in some areas of the Pt terrace step edges, which is probably due to an effectively higher exposure/temperature in our experiment, although the nominal values are very similar in the two studies. We believe that the mechanisms for Pt oxidation and the Pt oxide phase we obtained are the same as the ones described in ref 41. Although a definite assignment of the structures requires a calculation of the density of empty states for the R-PtO2/Pt(111) system, we tentatively assign the features appearing as holes at low bias (Figure 6a) to PtO2 islands thicker than the ones appearing as slight protrusions, the former being expected to have a lower density of empty states at low energies than the latter. Both kinds of structures appear 3 Å high in the images acquired with a high bias, this value being probably close to the real protruding height for both kinds of structures, corresponding to one OPtO layer. After ceria growth and O2 annealing the PtO2 islands are visible also in the regions where the ceria islands are present (see Figure 7, parts c and e). We suggest that the R-PtO2 nanostructures may play a role in the stabilization of the observed ceria epitaxial phase. Given the very large lattice mismatch between Pt and ceria (38%), the driving force for the observed epitaxial relation is not obvious. The protruding OPtO layers may then act as nucleating centers for the formation of ceria nanoislands, which assume an epitaxial orientation mediated by the presence of the PtO2 nanoislands. Ceria can minimize its surface strain by a 30 in-plane rotation with respect to PtO2, which corresponds to an alignment with the Pt(111) surface symmetry directions. This epitaxial orientation requires a 3% contraction of the lattice constant of ceria. The value measured by LEED is contracted compared to the bulk value, although not as much as in the case of perfect epitaxy with the PtO2. A certain degree of relaxation is expected, however, since the ceria islands are much larger in size than the PtO2 islands. The presence of the PtO2 islands may also be responsible for the fact that the minimum cerium oxide thickness observed is 2 ML. The ceria islands start nucleating epitaxially on top of the Pt oxide nanostructures, 3 Å high; therefore, the ceria structures protrude by a minimum of 6 Å from the Pt surface, being partly 2 CeO2 ML thick and partly 1 CeO2 ML on top of a PtO2 ML. A sketch summarizing our model for the ceria/Pt interface is shown in Figure 7d. The presence of PtO2 nanostructures also below the ceria islands and the different conductivity of a ceria bilayer compared to a single CeO2 ML on top of a PtO2 island probably give origin to the contrast which appears on the ceria island surface at low values of bias (Figure 7e). We suggest that a similar mechanism, involving the formation of an oxidized phase of the metal substrate, may come into play for the stabilization of the observed epitaxial orientation for ceria films also on the other metallic substrates.13,36,38 The satellites in the LEED patterns at low coverage (Figure 4b and relative inset) are consistent with the presence of additional domains rotated by (5 with each other. The mechanism for the stabilization of these domains can be qualitatively understood assuming that for ceria islands of small size adsorption sites different from the ones along the main symmetry azimuths may be stabilized. These become progressively less stable as the island size increases. Annealing in UHV has been observed to modify the stoichiometry of cerium oxide ultrathin films and bulk samples, and in the
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case of films it may lead to the formation of alloys with the substrate.21,36,42 Of course the resulting stoichiometry is expected to strongly depend not only on the annealing temperature but also on the stability of alloyed phases involving cerium, oxygen, and substrate atoms. Our results are in qualitative agreement with previous results indicating a substantial reduction of ceria films annealed in UHV.18 The formation of cerium alloys with substrates as Pt21 and Rh36 after high-temperature annealing of CeO2 films has been reported, whereas on Cu a CeO2Ce2O3 transition is observed above 900 K.42 The temperature stability range of our samples is in agreement with the one given by Berner and Schierbaum,21 although we have no direct evidence for alloy formation after the last step of UHV annealing. The area of Pt 4f XPS peak is in fact unaltered after annealing, and the LEED pattern cannot be associated with any of the ones reported for CePt surface alloys.43,44 The LEED pattern does not correspond to the one expected for Ce2O3, which has a surface lattice parameter close to the one of CeO2.42 It is probably due to an intermediate transition phase, which deserves further investigation.
5. CONCLUSIONS We obtained epitaxial cerium oxide ultrathin films on Pt(111) by reactive growth of Ce in molecular or atomic oxygen followed by high-temperature treatments in O2. The structures have a lateral size of several tens of nanometers, a minimum height of two OCeO layers, atomically flat surfaces, and well-defined edges. The obtained nanostructures have the fluorite structure with the (111) surface orientation and the CeO2Æ110æ//PtÆ110æ epitaxial orientation. The stoichiometry of the nanostructures is mainly CeO2, although a small fraction of Ce3þ ions are also detected. The Ce3þ concentration can be minimized by the use of atomic oxygen instead of molecular oxygen as oxidizing gas during the growth. The surface lattice parameter of the obtained ceria overlayers is systematically smaller than the bulk one at all of the investigated thicknesses. Substrate oxidation is proposed to have a relevant role in stabilizing the observed epitaxial orientation. The ceria films are found to be thermally stable up to 950 K. At higher temperatures the Ce3þ concentration increases, but the stoichiometry can be brought back to the original one by annealing in oxygen. ’ ACKNOWLEDGMENT The authors thank Valdimir Matolín for fruitful discussion on the methodology for Ce 3d spectra fitting. We also acknowledge the use of the XPSmania software analysis developed by Francesco Bruno (Aloisa beamline staff, CNR-IOM, Trieste, 2007). ’ REFERENCES (1) Trovarelli, A. Catalysis by Ceria and Related Materials; Imperial College Press: London, 2002. (2) Park, S.; Vohs, J. M.; Gorte, R. J. Nature 2000, 404, 265. (3) Chen, C.-H.; Chang, I. Y.-K.; Leea, J. Y.-M.; Chiu, F.-C. Appl. Phys. Lett. 2008, 92, 043507. (4) Jia, Q. X.; Foltyn, S. R.; Arendt, P. N.; Smith, J. F. Appl. Phys. Lett. 2002, 80, 1601. (5) Conesa, J. C. Surf. Sci. 1995, 339, 337. (6) N€orenberg, H.; Briggs, G. A. D. Phys. Rev. Lett. 1997, 79, 4222. (7) Skorodumova, N. V.; Simak, S. I.; Lundqvist, B. I.; Abrikosov, I. A.; Johansson, B. Phys. Rev. Lett. 2002, 89, 166601. 10725
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