Tip-Dependent Scanning Tunneling Microscopy Imaging of Ultrathin

ARTICLE pubs.acs.org/JPCC

Tip-Dependent Scanning Tunneling Microscopy Imaging of Ultrathin FeO Films on Pt(111) Lindsay R. Merte,† Lars C. Grabow,‡,§ Guowen Peng,‡ Jan Knudsen,†,|| Helene Zeuthen,† Wilhelmine Kudernatsch,† Soeren Porsgaard,† Erik Lægsgaard,† Manos Mavrikakis,‡ and Flemming Besenbacher*,† †

Interdisciplinary Nanoscience Center (iNANO) and Department of Physics and Astronomy, Aarhus University, DK-8000 Aarhus C, Denmark ‡ Department of Chemical and Biological Engineering, University of Wisconsin;Madison, Madison, Wisconsin 53706, United States ABSTRACT: High-resolution scanning tunneling microscope (STM) images of moire-structured FeO films on Pt(111) were obtained in a number of different tip-dependent imaging modes. For the first time, the STM images are distinguished and interpreted unambiguously with the help of distinct oxygen-vacancy dislocation loops in the FeO moire structure. The experimental STM results are compared with the results of electronic structure calculations within the DFTþU scheme for √ √ a realistic ( 91 91)R5.2° moire FeO unit cell supported on Pt(111) as well as with the results from previous studies. We find that one type of STM imaging mode, showing both Fe and O atoms, agrees well with simulated STM images, indicating that the simple Tersoff-Hamann theory is partially valid for this imaging mode. In addition, we identify other distinct, element-specific imaging modes which reveal a strong dependence on the specific tip apex state and likely result from specific tip-sample chemical interactions. From the present STM results we show that several of the previously published conclusions for the FeO system have to be revisited.

1. INTRODUCTION The study of ultrathin oxide films is important to a number of subfields of basic and applied sciences, on one hand, due to their decisive role in the functional properties of various advanced nanomaterials from electronic devices1 to heterogeneous catalysts2 and, on the other hand, due to the novel physical phenomena observed and anticipated for such ultrathin oxide films.3 It has become quite clear that oxides of only one or a few layers’ thickness cannot be treated simply as thin slices from corresponding bulk materials but rather should be considered unique 2-dimensional (2D) phases with distinct physical and chemical properties.4 For thin oxide films grown on metal surfaces, the presence of an electron reservoir and the restriction to two dimensions often leads to formation of complex oxide structures with unusual metal:oxygen ratios, and in some cases many different structures depending on the growth conditions appear.5-7 The single-layer FeO film grown on Pt(111) has been studied for more than two decades8 and is one of the best-characterized ultrathin oxide films. Although interest in the FeO film was initially based on heteroepitaxy and the study of (bulk-like) iron oxide surface properties,9 attention has more recently been focused on the novel properties of the ultrathin FeO film itself. It has been found, for example, that nanometer-size islands of FeO on Pt(111) catalyze CO oxidation under high-vacuum conditions10 and that a full monolayer of FeO on Pt(111) catalyzes CO oxidation under oxygen-rich conditions at higher pressures11 and reacts r 2011 American Chemical Society

readily with alcohols.12 Also intriguing is the local variability in the film’s properties dependent on the local interface structure, which is periodic on a scale of ∼2.5 nm due to the ∼10% lattice mismatch with the substrate. This characteristic moire superstructure has been shown to induce ordering of single gold atoms,13 MgPc molecules,14 and various metal clusters15 and to exhibit peculiar conductive behavior at low temperatures.16 The scanning tunneling microscope (STM) is an important instrument for elucidating atomic-scale structure and the electronic, chemical, and physical properties of ultrathin oxides,17 including FeO. This is partly due to the important role of atomicscale defects in determining chemical and electronic properties of ultrathin oxides and partly because of the inherent complexity of many oxide structures and the potential for coexistence of several different phases, which can make characterization by traditional averaging diffraction and spectroscopy techniques very challenging. The full power of the STM and its capability to resolve atomicscale features are coupled to the extreme distance dependence of the electron tunnelling between tip apex states and distinct surface states. However, the relationship between the observed STM images and the surface structure is not trivial, and it has Received: October 6, 2010 Revised: December 8, 2010 Published: January 10, 2011 2089

dx.doi.org/10.1021/jp109581a | J. Phys. Chem. C 2011, 115, 2089–2099

The Journal of Physical Chemistry C become clear that STM images may depend on a number of different factors, whose relative importance and explicit manifestation depend on the structure of the STM tip in addition to that of the sample.18-20 It is extremely difficult, from both experimental and computational points of view, to determine the detailed atomic-scale structure of an STM tip apex and its effect on imaging, and therefore, interpretation of STM images is usually based on phenomenological arguments or, if available, comparison with simulations based on simplified STM theories. For a proper theoretical treatment of the three-dimensional STM problem, we require a correct description of the electron potential in the tunnel-gap region between the tip and the sample surface and a detailed description of the electronic states of both the sample and the tip, including the effects of interaction between the two and of the local electric field. However, this is a formidable task, and therefore, a number of approximate models have been developed. In the simple, approximate model of Tersoff and Hamann,21 a direct correlation exists between the STM tunnel current and the energy-integrated local density of states (LDOS) projected to the position of the tip apex above the surface. This approximate approach has been used widely, often with great success, also for the study of ultrathin oxides (see, for example, refs 6, 22, and 23). However, it is important to remember that the Tersoff-Hamann model is a simplification and may fail to qualitatively reproduce the experimental STM images, especially if it appears that the specific tip states influence the recorded experimental STM images (see, for example, refs 24 and 25). Thus, although STM can, in principle, provide extremely detailed information about a surface structure, caution must always be taken in the interpretation of STM images. High-resolution STM images can, in principle, shed light on the unique properties of the monolayer FeO film and how these depend on the interface geometry and thus contribute to the growing understanding of polar ultrathin films. At present, however, a better understanding of the tunneling process and interpretation of the recorded STM images on the FeO ultrathin film is needed to extract detailed information about the geometrical and electronic structure of FeO. Our recent studies have demonstrated that detailed STM imaging of FeO/Pt(111) is far from being fully understood,26,27 and therefore, we have undertaken the present study where we focus especially on how differences in tip state influence the recorded STM images of the moire-structured ultrathin FeO films on Pt(111). A number of different STM images obtained with different tip states are presented and discussed, and we hope this will be helpful for interpretation of recorded STM images in the future. In order to gain further insight into the formation mechanisms of the recorded STM images, we compared them with the electronic structure of the FeO film calculated with the DFTþU method and with simulated STM images using the Tersoff-Hamann model.

2. EXPERIMENTAL AND COMPUTATIONAL METHODS All STM experiments were performed in a UHV chamber with base pressure below 1 10-10 Torr. The Pt(111) crystal was cleaned by cycles of Arþ sputtering and annealing at 1000 K. The FeO film was prepared by postoxidation of room-temperaturedeposited Fe at 870-1000 K in 1 10-6 Torr O2. STM images were obtained with a home-built Aarhus STM,28 with a mechanically cut PtIr tip. The STM images were obtained at 300 K in constant-current mode. Tip preparation was performed in situ by applying voltage pulses (5-10 V) to the sample during

ARTICLE

measurement. When such voltage pulses were unsuccessful in producing a sharp, stable tip, the tip was driven into gentle contact with the sample surface and pulses were again applied. The STM measurements were subsequently obtained in areas well separated from where tip preparation was carried out. The state of the STM tip apex, and hence the imaging mode observed, was found to vary during the course of the STM imaging, though tip changes could often be induced by applying mild voltage pulses. Our general conclusions are based on a collection of several hundred atomically resolved STM images obtained over the course of a large number of different experiments. Spin-polarized density functional theory (DFT) calculations using the DFTþU approach by Dudarev et al.29 were performed using the VASP code.30,31 As in previous DFT studies of the FeO/Pt(111) system,26,32,33 we used the projector-augmented wave (PAW) method34,35 and chose the parameters describing the on-site Coulomb interaction between Fe 3d orbitals as U = 4 eV and J = 1 eV (Ueff = U - J = 3 eV). To model the 1 ML √FeO/ Pt(111) surface we used the experimentally observed ( 91 √ 91)R5.2° unit cell consisting of one layer of Fe atoms and one layer of O atoms (“bilayer”) supported on a√three-layer √ Pt(111) slab. According to a previous study using a ( 84 84)R10.9° unit cell,33 the difference between three and seven supporting Pt layers is negligible, and for that we consider our setup converged with respect to the included number of Pt metal layers. The FeO film and the top Pt layer were fully relaxed. The experimental value for the Pt-Pt spacing in the (111) plane of 2.77 Å (average Fe-Fe spacing = 3.09 Å) was used. The Brillouin zone was sampled with the gamma point only. Exchange and correlation were described by the GGA-PW91 exchange-correlation functional,36 and a kinetic energy cutoff of 400 eV was used. The initial guess for the magnetic structure is based on a row-wise antiferromagnetic (RW-AFM) structure, with a magnetic defect, due to the odd number of Fe atoms, at the Fe atom directly atop a Pt atom at the topmost Pt layer, i.e., at the TOP domain shown in Figure 1. The spin pattern is shown in Figure 1, with Fe spins of opposite orientation marked with a white dot. For investigation of the influence of the magnetic structure on the local electronic structure the defect was also moved to the HCP and FCC domains (see Figure 1).

3. STM MEASUREMENTS The structure of the monolayer FeO film on Pt(111) has been studied in the past and steadily been refined from low-energy electron diffraction (LEED), STM, X-ray photoelectron spectroscopy (XPS), near-edge X-ray absorption fine structure (NEXAFS), and X-ray photoelectron diffraction (XPD) studies.37-42 The resulting structure, which is depicted in Figure 1, consists of a hexagonal Fe-O bilayer with an average lattice parameter of ∼3.1 Å, slightly rotated against the Pt(111) substrate lattice, resulting characteristic moire√ structure,√best approximated as √ in a√ a ( 73 73)R5.8°-FeO/( 91 91)R5.2°-Pt coincidence lattice. Only one of two possibilities, differentiated by the direction of rotation relative to the Pt(111) lattice, is shown in Figure 1. On the basis of primarily Fe-2p and O-1s XPD measurements,40 which are sensitive to the local binding environment of the two species, it was established that the rumpling of the film (i.e., the average spacing between the Fe and O layers) is 0.68 Å, with the O layer buckled away from the surface. Also clearly established in the XPD measurements was the influence of the substrate orientation on the FeO film’s own 2090

dx.doi.org/10.1021/jp109581a |J. Phys. Chem. C 2011, 115, 2089–2099

The Journal of Physical Chemistry C

ARTICLE

Figure 1. Top and side views of the atomic structure of the 1 ML FeO film on Pt(111) as determined by previous STM, LEED, and XPD measurements, including the stacking of atomic layers. The moire-type superstructure unit cell is indicated. Two orientations are possible, distinguished by the placement of the O layer above the Fe layer, but only the orientation on the left is observed experimentally. The six different high-symmetry domains of the moire structure are indicated, along with the definitions used in this work. White dots are used to mark Fe atoms with opposite magnetic moment in the calculations.

orientation, beyond the (near) alignment of close packed rows. The two possible orientations, rendered inequivalent by the placement of the second Pt layer, are depicted in Figure 1 and can be interchanged either by rotation of the overlayer by 60° or, as depicted here, by shifting the O sublattice from one Fe hollow site to another. Due to the incommensurate epitaxial relationship, the in-plane coordinates of Fe and O atoms relative to the highsymmetry surface sites vary continuously across the surface, but the favored and unfavored orientations can nonetheless be defined in terms of the ABC stacking sequence of the underlying Pt substrate. As can be deduced from Figure 1, the relationship between the Fe and the O layers in the favored orientation is the same as between two adjacent close-packed Pt layers, while in the unfavored orientation this relationship is reversed. This difference is reflected in the definitions of the six high-symmetry domains indicated in Figure 1. In the present study, these domains are defined according to the stacking sequence of the top three atomic layers (i.e., TOP and TOP* have Fe atop Pt and O at hollow sites, HCP and HCP* have Fe at hollow sites and O atop Pt, and FCC and FCC* have both Fe and O above hollow sites), since this is expected to have the greatest influence on its structure and physical properties; an asterisk is used to denote the domains in the unfavored orientation. 3.1. STM Image Interpretation. In the 2D projection of the FeO film structure, the Fe and O atoms are symmetrically equivalent, as are the three high-symmetry points of the moire FeO superstructure. Thus, it is not possible, a priori, to distinguish the atomic species or the different domains in STM images of the defect-free FeO film. However, we demonstrated how we can reliably interpret STM images26 by making use of unique vacancy defects which are intentionally introduced by exposure of the FeO film to atomic hydrogen (Hat). As described in ref 27, exposure to Hat leads initially to formation of surface hydroxyls (OH), which react to form H2O, leaving behind oxygen vacancies, when the OH coverage exceeds ∼10-15% ML. These vacancies cause the formation of triangular features observed in STM images.

Figure 2a shows an STM image of a Hat-exposed FeO film close to this OH coverage threshold, where OH groups are observed both inside and outside a triangular dislocation feature. With an overlaid hexagonal lattice, it becomes clear that the OH groups inside the triangle are shifted out of registry with those outside, indicating that the triangular features consist of vacancy dislocation loops in the O sublattice. The resolution in the recorded STM images is furthermore sufficient to determine the direction of the shift, as it is seen that the OH groups inside the triangle occupy the left halves of the unit cells marked. This is compared with STM images of the triangular features without adsorbed hydrogen, obtained in different tip-dependent imaging modes. The STM image in Figure 2b shows protrusions corresponding to Fe atoms, and the overlaid lines show that the Fe atoms inside the triangular feature are in registry with the rest of the Fe sublattice. The STM image depicted in Figure 2c, on the other hand, shows protrusions corresponding to O atoms. Similar overlaid lines indicate that the O atoms inside the triangular feature are shifted out of registry with the surrounding O sublattice, in the same direction as that observed for OH groups in Figure 2a. From this we deduce the structural model as the one depicted in Figure 2d. Here, a dislocation loop is formed when several O vacancies (in this case 7, depicted schematically as pale red disks) accumulate, allowing a triangular group of O atoms (in this case 21, located inside the marked triangle) to shift positions to adjacent hollow sites of the Fe lattice, creating a small domain of FeO in the opposite orientation (cf. Figure 1) to the surrounding film. Such dislocation loops can, in principle, be formed in any conceivable size, depending on the number of O vacancies incorporated. However, all the dislocation loops observed experimentally can be divided into two distinct types, a smaller type incorporating ∼5-8 vacancies and a larger type incorporating ∼12-14 vacancies. The histogram depicted in Figure 2e shows the size distribution observed in a representative experiment and indicates that dislocation loops incorporating 6 or 7 vacancies are by far the most common. The two types of dislocation loops are 2091



ARTICLE

Figure 2. Illustration of the use of oxygen vacancy dislocation loops to interpret STM images of the FeO film. (a) STM image of the FeO film obtained directly after exposure to atomic hydrogen, showing OH groups located inside and outside a triangular dislocation loop, the edges of which are marked with a yellow line. (b and c) STM images with different tip states obtained after exposure to atomic hydrogen and annealing, during which the film is partially reduced, leaving well-ordered dislocation loops. (b) Protrusions are observed to be in registry with the surrounding film, indicating that Fe atoms are imaged; (c) protrusions are out of registry with the surrounding film, indicating that O atoms are imaged. (d) Atomic model corresponding to the oxygen-vacancy dislocation loops shown in b and c. (e) Size distribution of oxygen-vacancy dislocation loops obtained from atomically resolved STM images. (f and g) STM image of a partially reduced FeO film showing one large and three small dislocation loops. The moire unit cell is marked in f, with corners at TOP/TOP* domains. The lattice in g marks the positions of Fe atoms in the structure, and the FCC, HCP, and TOP domains are marked with a square, triangle, and circle, respectively. (h) Ball model showing the surface atomic structure deduced from f and g.

distinguished not only by their size but also by their locations relative to the FeO moire structure. In order to determine which domains these are, we turn back to Figure 1 and examine the relationships between the different domains of the moire structure in the two possible orientations, obtained by shifting the O sublattice while keeping the Fe lattice fixed. The names of the different domains used here are chosen specifically to draw attention to the similarities between domains having identical structures when only the top three atomic layers (i.e., Pt-Fe-O, equivalent to stating that fcc and hcp hollow sites of the Pt(111) surface are identical) are taken into account. Here, it should be particularly noted that the TOP domain is unique in that displacement of the O lattice leaves the basic structure unchanged, as the O atoms shift from fcc to hcp hollow sites of the Pt(111) surface, leaving the Fe atoms in atop positions with respect to Pt sites. In contrast, the FCC and HCP domains are effectively inverted, as the O atoms shift from sites atop Pt to hollow sites and vice versa. Since the main features of the film’s physical properties are expected to be determined by its interaction with the top Pt layer, with differences due to the deeper layers being much weaker, we can use the “invariance” of the TOP/TOP* domains to identify them on the surface. The application of this scheme to interpret STM images of the FeO film is demonstrated in Figure 2f-h, which shows an STM image of a region of the film with one large and three small dislocation loops. We note first that the domain at the center of the large dislocation loop exhibits very distinct protrusions, very similar to one type of domain found outside the vacancy dislocation loop. The large moire cell superimposed upon Figure 2f indicates that this domain is “invariant” upon inversion of orientation and thus must be the TOP/TOP* domain. Noting, now, that the orientation of the dislocation loops, which is directly given by the orientation of the film itself, is sufficient to determine the locations

of the different domains relative to one another, we can deduce that the small dislocation loops are located at HCP domains, with edges forming the boundary with adjacent TOP domains. Taking this analysis further, we are also able to determine the atomic positions in the STM image by noting once again that the positions of the Fe atoms are fixed and therefore can be used as reference points in the STM image. In Figure 2g, a hexagonal lattice has been superimposed using the bright protrusions observed at the TOP domains. The bright protrusions found at the TOP* domain inside the large dislocation loop fall on this lattice and hence must be associated with Fe atoms. Consequently, the O atoms inside the dislocation loops must be located in the right halves of the small FeO unit cell (indicated in green in Figure 2d), and outside dislocation loops they must be located in the left halves of the small unit cell. Comparing now, for example, an FCC domain of the film with an FCC* domain within a vacancy dislocation loop, we measure a higher apparent height above the O atoms (on the left half of the cell for FCC and on the right for FCC*) and lower apparent height above the hollow site. A complete atomic model, which thus can be deduced from the STM image analysis, is depicted in Figure 2h. This assignment of the different atoms and high-symmetry domains based on the invariance of the TOP domain is fully consistent with all our measurements and with a range of calculations. STM measurements of the FeO film with dislocation loops at field emission conditions26 indicate that the TOP and TOP* domains show the lowest local surface potential, in full agreement with all calculations to date. We furthermore determined the orientation of the Pt(111) single crystals used in our studies by Laue X-ray diffraction and thus established that the orientation of the FeO film deduced from that of the vacancy dislocation loops, relative to that of the substrate, is indeed consistent with the photoelectron diffraction measurements of 2092



Figure 3. (a and b) STM images (þ65 mV, 3 nA) in two similar but distinct tip-dependent imaging modes obtained on the FeO film. Arrows begin at FCC domains and end at TOP domains. (c) Line profiles extracted from images a and b taken across the small (3 Å) unit cell diagonal, showing the variation across the high-symmetry domains. Positions of Fe and O atoms are marked with black and red circles and lines, respectively.

Kim et al.40 We can thus conclude that, although the differentiation between the FCC and HCP domains (and thus also the relative positions of Fe and O atoms) may appear to be a very delicate issue, only one possibility is consistent with measurements, and the assignments here must be considered accurate with very high confidence. 3.2. “Mixed” Imaging Modes. The STM image depicted in Figure 2f and 2g is representative of the most common imaging mode observed in our STM studies of FeO. We refer to this imaging mode as “mixed” because the protrusions observed in STM images correspond to either Fe or O, depending on the high-symmetry domain. Two further examples of this type of imaging mode are presented in Figure 3a and 3b, together with line profiles measured along the small FeO unit cell diagonal. The most distinctive features of these STM images are the protrusions over Fe atoms at the TOP domains. At the FCC domains, the highest points locally correspond to O positions, whereas at the HCP domains, nearly equal tip heights are observed over Fe and O. Nowhere on the surface do local maxima correspond to hollow sites. The most striking difference between the two images depicted in Figure 3a and 3b, aside from the corrugation, is the apparent height of the TOP domain, which appears darker (i.e., with lower

ARTICLE

average apparent height) than the FCC domain in Figure 3b but brighter (higher average apparent height) than the FCC domain in Figure 3a. Although STM images similar to Figure 3a, where the FCC domain is imaged with lowest apparent height (darkest), are found to be the most common in our studies, variability in the domains’ relative apparent heights at low bias is frequently observed. To differentiate between these two similar STM imaging modes, we designate the one shown in Figure 3a, characterized at low bias by large corrugation and bright TOP domains, as “mixed mode 1” (hereafter referred to as MM1), and the one in Figure 3b, characterized by low corrugation and dark TOP domains, as “mixed mode 2” (hereafter, MM2). In addition to being the most frequently observed, the mixed imaging modes exhibit relatively high stability against tip changes over a wide range of tunneling conditions and are the only modes for which we successfully obtained reproducible bias-dependent STM measurements. Two such bias-dependent STM data sets are depicted in Figure 4, where the left column was obtained in MM1, similar to Figure 3a, and the center column was obtained in MM2, similar to Figure 3b. In the right column we also show simulated STM images based on the DFTþU calculations using the Tersoff-Hamann approach for comparison. Both the experimentally recorded and the simulated STM images show similar trends. The apparent height of the TOP domains (relative to the cell average) shows the most distinct bias dependence, increasing with increased positive bias to become the brightest domain at U = þ1500 mV and decreasing with increased negative bias to become the darkest domain at U e -1500 mV. The apparent heights of the FCC and HCP domains at large positive and negative bias values are more similar and appear to be more sensitive to the tip state, as the MM2 series shows FCC < HCP at þ1500 mV and HCP < FCC at -1500 mV, while the MM1 series shows HCP < FCC at þ1500 mV and FCC < HCP at -1736 mV. The simulated image at U = þ500 mV agrees qualitatively well with the images taken in MM1, but for U = -500 mV and at large positive and negative bias a better agreement with MM2 is seen. The simulations at low bias ((65 mV) do not show strong variations in apparent height. However, the TOP domains appear brightest at both voltages, whereas the HCP domain appears brightest in both experimental images. The much higher corrugation observed in the MM1 series (∼25 pm at 65 mV, see Figure 3c) allowed atomic features to be distinguished across nearly the entire range of bias values, the exception being at þ1500 mV, where the corrugation decreased significantly and has therefore been compensated with an increase in current set point. Over this range, the local maxima observed at TOP domains correspond to Fe positions, while those at FCC domains correspond to O positions. Although for low bias voltages the apparent heights over O and Fe atoms in HCP domains are similar, at larger positive and negative bias values the contribution due to Fe decreases and only protrusions over O atoms are observed. 3.3. Chemically Specific Imaging Modes. In addition to the “mixed modes” described above, we observe several other imaging modes which appear to be chemically specific, showing protrusions over either Fe or O atoms. These imaging modes have previously proven useful in the characterization of oxygen-deficient FeOx ultrathin films.27 Three of these element-specific imaging modes are depicted in Figure 5. The fact that STM images obtained in these modes reveal only one type of atom appearing as protrusions makes the interpretation more challenging than was the 2093



ARTICLE

Figure 5. STM images obtained in three different chemically specific tip-dependent imaging modes. (a-c) Images obtained of the same region of a mildly reduced FeO film between spontaneous tip changes. White lines mark an atomic row outside the large dislocation loop. The registry shift observed in c indicates that O atoms are imaged. (d-f) Higher resolution STM images obtained in imaging modes corresponding to those in a-c. Arrows begin at FCC domains and end at TOP domains. (g) Line profiles extracted from STM images d-f along the small (3 Å) unit cell diagonal. Positions of Fe and O atoms are marked with black and red circles and lines, respectively. Figure 4. Bias dependence of STM imaging in the two mixed modes and DFTþU-simulated images using the Tersoff-Hamann theory in the natural unit cell, with electron density isosurfaces taken at 1 10-5 e/Å3. Symbols mark the three high-symmetry domains: (0) FCC, (4) HCP, and (O) TOP. Blue parallel lines on the left column mark O rows. All images obtained at 3 nA current, except that of Mixed Mode #1 at 1500 mV, where the current was set to 8.4 nA to maintain atomic resolution.

case for the mixed modes, as both an oxygen-specific and an ironspecific tip state will produce an STM image of a hexagonal lattice with some modulation of height and/or corrugation due to the FeO moire structure. As above, the existence of the vacancy dislocation loops aids us in the interpretation of STM images, as the presence or absence, as well as the direction, of a lattice shift can be used to distinguish the different atoms. Figure 5a-c shows three STM images of a single dislocation loop following successive tip changes. In Figure 5a and 5b we clearly image the Fe atoms as protrusions, although in Figure 5a the domain at the center of the dislocation loops is dark while in Figure 5b this domain is bright.

In Figure 5c, on the other hand, we image the O atoms as protrusions. Higher resolution STM images of the FeO film in these three element-specific imaging modes are depicted in Figure 5d-f, together with corresponding line profiles (Figure 5g). In contrast to the mixed modes, these chemically specific imaging modes show lower stability and typically spontaneous tip changes occur after only a few minutes of measurement. High-resolution images are obtained only over a fairly narrow range of bias and current settings where the specific tip is semistable. Consequently, the presented images were not acquired under exactly identical conditions, which makes quantitative comparisons somewhat challenging. All of the images presented in Figure 5d-f, however, show relatively high corrugation amplitudes, in each case equal to or greater than that for the more common mixed modes MM1 and MM2 discussed above (Figure 3a). Particularly noteworthy are two observations. 2094


The Journal of Physical Chemistry C (1) The Fe mode 1 and the O mode STM images in Figure 5d and 5f, respectively, although depicting different species, exhibit similar trends in the apparent heights of the highsymmetry domains, with the FCC domains appearing highest and TOP domains lowest. (2) The two Fe modes 1 and 2 depicted in Figure 5d and 5e, respectively, reveal the same species (Fe) although with different apparent heights for the different high-symmetry domains. It is therefore important to have some reference features to indicate both the type of lattice imaged and the locations of the different domains, since otherwise these three imaging modes are nearly indistinguishable.

4. DFT CALCULATIONS Relatively recently, DFT calculations have been performed on the FeO film supported on Pt(111), yielding further insight into its properties, particularly with regard to the effect of the inhomogeneous interface structure.32,33 Such calculations have been necessarily limited by the need to balance accuracy against computational cost, and hence, approximations have been made in each case to limit the number of atoms per unit cell. Despite the recently demonstrated possibility of simulating the FeO film with unit cells approaching realistic size, models of the film’s structure are still restricted due to the as-yet undetermined spin pattern of the Fe lattice and the imposed translational symmetry. Calculations so far suggest a strong preference for antiferromagnetic ordering,32 consistent with ferromagnetic resonance (FMR) measurements, which detected no signs of ferromagnetism.16 The geometrical moire unit cells observed experimentally (or rather their closest approximations) are compatible neither with antiferromagnetic collinear structures (i.e., alternating rows of up and down spins) nor with noncollinear Neel-type structures (adjacent spins are oriented 120° from each other, so that the net spin over any triplet is 0), although strict commensurability between the magnetic and the geometric superstructures is not necessarily required for the real film. The further possibility of even more complex magnetic structures makes the task of simulation even more challenging. In our calculations, the spin pattern used is shown in Figure 1, with Fe spins of opposite orientation marked with a white dot. This spin pattern is based on a row-wise antiferromagnetic (RW-AFM) structure, with a defect, due to the odd number of Fe atoms, at the TOP domain. √ √ 4.1. FeO Atomic Structure. Using the very large ( 91 91)R5.2° unit cell for both possible orientations of the FeO film and the row-wise antiferromagnetic ordering described above as an initial guess for the magnetic structure, we performed a geometric optimization of the surface. In the optimized structure we find a variation of the Fe-Fe lattice parameter across the surface with the largest distance in the TOP domain (3.18 Å) and the smallest distance in the FCC domain (3.00 Å). The Fe-Fe distance in the HCP domain is equal to the cell average of 3.09 Å. The local contraction of the FCC domain is an indication that it is energetically most stable with the strongest FeO-substrate interaction. On the contrary, the TOP domain is the least stable and a local lattice expansion is observed. Previous DFT calculations in smaller, pseudomorphic unit cells, which allow separate simulations of the individual domains by changing the position of the FeO film, support this interpretation, showing that FCC domains are more stable than HCP domains by 0.1-0.2 eV/FeO unit and HCP domains are more stable than TOP domains by 0.03-0.2 eV/FeO unit, depending on the model.26,32 The geometric differences between the domains are

ARTICLE

Table 1. Geometric and Electronic Properties √ √of 1 ML FeO(111)/Pt(111) Calculated in the ( 91 91)R5.2° Unit Cella domain

l/Å

zFe-O/Å

Δz/eV

δV/eV

avg.

3.09

2.67

0.68

0

TOP

3.18

2.78

0.58

-0.40

HCP

3.09

2.62

0.71

0.10

FCC

3.00

2.64

0.77

0.37

avg.* TOP*

3.09 3.18

2.67 2.77

0.68 0.59

0 -0.41

HCP*

3.10

2.62

0.72

0.13

FCC*

3.01

2.65

0.76

0.35

a

l is the FeO lattice parameter; zFe-O = (zFe þ zO)/2 is the average interface height; Δz = zO - zFe is the surface rumpling; δV is the variation of the electrostatic potential with respect to the cell average (5.65 eV for TOP, HCP, FCC; 5.64 eV for TOP*, HCP*, FCC*).

not limited to in-plane variations; zFe-O, the average height of the FeO layer above the top Pt layer, increases from 2.62 Å in the HCP to 2.64 Å in the FCC domain and eventually reaches 2.78 Å in the TOP domain. Lastly, the rumpling of the surface Δz, defined as the distance between the Fe and the O layers, is largest in the FCC domain (0.77 Å) and smallest in the TOP domain (0.58 Å). This trend in the FeO film rumpling matches that of the electrostatic potential above the surface, in agreement with experiment26,43 and with the calculations of Zhang et al.33 In addition, the calculated average rumpling of 0.68 Å is in excellent agreement with the value (0.68 ( 0.05 Å) measured by Kim et al.40 The calculated geometric properties of the Pt-supported FeO √ film√agree well with the calculations by Zhang et al. in a ( 84 84)R10.9° unit cell33 and are summarized in Table 1. 4.2. FeO Electronic Structure. Before discussing the electronic structure of the individual FeO domains in detail it is necessary to investigate the influence of the magnetic defect on the local electronic structure. Hence, we calculated the partial density of states (PDOS) with the magnetic defect located in the HCP and FCC domain while keeping the optimized geometry of the FeO/ Pt(111) system fixed. Although spin-polarized PDOS data were calculated, we do not distinguish between spin states and rather present the sum of both states. This is motivated by the spin state ambiguity introduced by the magnetic defect in our structure. The PDOS in each region was then extracted as appropriate either from one or three Fe (or O) atom(s) at the center of each domain. We find that the Fe-3dz2 and O-2pz states dominate at the Fermi level, in agreement with earlier publications.32,33 In Figure 6 the Fe-3dz2 and O-2pz PDOS in the TOP, HCP, and FCC domains are shown for three possible locations of the magnetic defect. For the TOP domain the blue line, representing the LDOS with the magnetic defect in the TOP domain, deviates from the green and red line, which represent the LDOS when the magnetic defect is located in the HCP and FCC domain, respectively. Similarly, the green line differs from the blue and red line in the HCP domain, while in the FCC domain the blue and green line coincide and the red line deviates. This clearly shows that the magnetic defect directly influences the local electronic structure. Since the actual spin pattern of the FeO film is not known, for our further analysis we use the PDOS, which is obtained by averaging the PDOS in each domain while neglecting the PDOS with the magnetic defect located in the same domain. For example, the PDOS in the TOP domain is obtained as the average of the PDOS with the 2095



Figure 6. Partial density of states (PDOS) of the Fe-3dz2 and O-2pz states in the TOP, HCP, and FCC domain for three different locations of the magnetic defect.

magnetic defect located in the HCP and FCC domain. The resulting PDOS can be considered “defect-free”. The “defect-free” PDOS of the Fe-3dz2 and O-2pz states in the TOP, HCP, and FCC domains is depicted in Figure 7. Below the Fermi level the Fe-3dz2 PDOS is similar in all three regions, and strong features are only found above the Fermi level. At ca. 0.3 eV above the Fermi level, the TOP domain shows a strong Fe-3dz2 feature which coincides with a small √ maximum √ in the O-2pz PDOS. Zhang et al.,33 working with a ( 84 84)R10.9° unit cell, previously reported a similar feature located at 0.4 eV above EF. The HCP region has peaks at 0.5 and 1.1 eV in the Fe-3dz2 PDOS and at -0.4 and 1.1 eV in the O-2pz PDOS. The large O-2pz PDOS for negative energies agrees well with the observed brightness of the HCP region in our STM images at negative bias in both mixed modes and the simulated images (see Figure 4). Lastly, the FCC domain shows dominant features in the Fe-3dz2 and O-2pz PDOS at >1.2 eV above the Fermi level, which may explain the brightening of the FCC region obtained in MM1 at fairly large positive tunnel bias. As a first-order theoretical approximation to the experimental STM line profiles presented in Figure 3 we studied the partial charge density at the surface in the ranges [-1.5 eV, -0.5 eV], [-0.5 eV, 0.5 eV], and [0.5 eV, 1.5 eV]. Vertical cuts of the charge density along the small FeO unit cell diagonal are shown in Figure 8a. For exclusively negative and exclusively positive energy ranges ([-1.5 eV, -0.5 eV], [0.5 eV, 1.5 eV]) the highest charge density is always found over O atoms in all three domains. However, when a symmetric range around the Fermi level is chosen ([-0.5 eV, 0.5 eV]), then the highest charge density in the TOP region is located over Fe atoms. This agrees well with the experimental STM line profile for the MM1 in Figure 3a, showing local maxima over O atoms in the FCC domain and over Fe atoms in the TOP domain. Further similarities with MM1 are revealed in Figure 8b, showing a horizontal cut of the partial charge density in the range from -0.5 to 0.5 eV taken at 4.89 Å above the Pt layer. As in MM1, the FCC region appears dark whereas the HCP and TOP domains appear bright. Thus, depending on the energy range used for calculating the partial charge density we are able to simulate STM images within the Tersoff-Hamann approximation that resemble the STM images recorded in MM1 or MM2 (Figure 4). This suggests that STM

ARTICLE

Figure 7. Fe-3dz2 and O-2pz partial density of states after correcting for the magnetic defect in the TOP (blue), HCP (green), and FCC (red) domains (see text).

images recorded in these mixed modes are indeed representative of the LDOS at the surface and are not strongly influenced by the state of the STM tip apex. In turn, this is consistent with the stability of these imaging modes and the possibility to obtain a series of bias-dependent images as shown in Figure 4.

5. DISCUSSION Although STM images of FeO films on Pt(111) have been frequently published in the literature, it is generally the case that the interpretation of these images has been uncertain. On the basis of the current results, we can examine previously published experimental and theoretical results regarding STM imaging of the ultrathin FeO layer grown on Pt(111) and in some cases offer reinterpretation. The fact that the recorded STM images may depend strongly on the tip state complicates comparison of recorded STM images from different groups considerably, but reviewing the published literature reveals a good degree of consistency. The most often-observed feature of published STM images of FeO/Pt(111), which was noted even in the first article utilizing this technique,38 is the distinctively higher atomic-scale corrugation observed at one domain of the FeO moire structure compared to the others. Such STM images have been reported in numerous previous publications by, e.g., Galloway et al.,44,45 Rienks et al.,43 Giordano et al.,32 Lemire et al.,46 Nilius et al.,13 Fu et al.,10 and Yao et al., 47 in addition to publications from our own laboratory.26,27 Comparison with our own present data suggests that all of these published images should be classified as “mixed mode”, and the distinctive protrusions observed in such STM images should therefore be associated with Fe atoms at TOP domains. In this case, the remaining two domains can be identified based on their apparent heights, as HCP domains are always found to appear brighter than FCC domains in lowbias STM images in the mixed modes. STM images in other tip-dependent imaging modes, revealing a single atomic species only, are less common in the literature but can be found in publications by Ritter et al.41 and Ma et al.48 As described above, it is difficult to interpret such element-specific imaging modes without an internal reference, and thus, it is unclear whether Fe or O atoms are depicted in such STM images. 2096



ARTICLE

Figure 8. (a) Vertical cuts of the calculated partial charge density in the energy ranges [-1.5 eV, -0.5 eV], [-0.5 eV, 0.5 eV], and [0.5 eV, 1.5 eV] with respect to the Fermi level. The cuts are taken along the small unit cell diagonal indicated with the blue arrow in b. The location of the high-symmetry domains is indicated by the symbols: (0) FCC, (4) HCP, and (O) TOP. The average heights z of the Pt, Fe, and O layers are 0.11, 2.44, and 3.12 Å, respectively. (b) Horizontal cut of the calculated partial charge density in the range [-0.5 eV, 0.5 eV] at z = 5 Å (4.89 Å above the Pt layer).

We are aware of three previously published studies specifically investigating STM imaging on FeO films on Pt(111). The electron-scattering quantum chemistry (ESQC) STM simulations of Galloway et al.45 were the earliest of these and remain, to date, the only simulations reported which considered the structure of the STM tip explicitly. The model tips considered included one terminated by a Pt monomer, one with a Pt trimer, and a Pt trimer with an adsorbed O atom. The FeO film was modeled using the observed FeO average lattice parameter for both the overlayer and the substrate, and the tunneling current was calculated for the different domains by changing the registry between the film and substrate lattices. These calculations indicated that for all tips and all domains O-2pz states make the dominant contribution to the tunneling current, leading to protrusions over O atoms for the Pt tips and over Fe atoms for the O-terminated tip, due to tunneling between neighboring O-2pz states of the film and the supporting Pt trimer. These calculations further indicated that the atomic-scale corrugation of the TOP domains is highest, while those of the FCC and HCP domains are lower by about one-half, the latter being lowest overall. It is worth noting that our STM images recorded in the mixed modes show some agreement with the STM simulations by Galloway et al.45 In particular, the magnitudes of the corrugations of the three domains, in MM2, are very close to those calculated with a Pt monomer tip, and the prediction that the highest point in an STM topograph is found at an HCP domain is reproduced, although the measured corrugation of the FeO moire structure is much larger than that of the calculation (∼85 vs ∼25 pm calculated). The variation in average apparent height in MM1 agrees well both qualitatively and quantitatively with calculations for an O-terminated tip. The observation that the highest and lowest points in the STM topograph are found at TOP and FCC domains, respectively, is also reproduced in the calculations. The measured atomic-scale corrugation is, however, much higher (about a factor of 2 higher than the calculated value at the TOP domain) and shows little variation across the moire unit cell. However, most disconcerting in light of the present STM results for the MM1 and MM2 modes is the fact that these ESQC simulations predict protrusions over one atomic species only and thus apparently miss the definitive feature of the mixed modes (the fact that Fe is imaged at TOP domains and O at FCC

domains). In this regard it should be stressed that the three element-specific imaging modes observed experimentally above do not exhibit the calculated variability of the atomic-scale corrugation, and the magnitudes of the corrugations are also higher than predicted in the calculations. Furthermore, the chemically specific imaging modes presented above, Fe mode 1 and O mode, both exhibit global maxima at FCC domains which none of the calculations by Galloway et al.45 predict. The Fe 2 mode shows qualitative agreement with the calculations for a Pt monomer regarding the order of apparent heights of the three domains, but of course protrusions are observed over Fe atoms in these images, while the calculations by Galloway et al. predicted protrusions over O atoms. Thus, it is evident that the early ESQC calculations by Galloway et al. do not reproduce any of the STM results completely, neither for the mixed tip modes nor for the chemically specific modes. It would therefore be interesting to compare these fairly old theoretical calculations with calculations using more elaborate theoretical approaches which have recently become feasible, taking the specific tip state into account while also allowing full geometric and electronic structure optimization. The two more recent studies by Giordano et al.32 and Zhang et al.33 utilized DFT and focused primarily on the average apparent heights of the high-symmetry domains and their variation with tunneling bias. However, in these cases the theoretical simulations of the STM images have been based on the simple Tersoff-Hamann approximation, where the effect of tip structure is neglected and measured constant-current images are assumed to directly reflect the surface electronic structure (i.e., the LDOS above the surface, integrated from EF to the eU). The simulations by Giordano et al. avoided the use of large unit cells without putting drastic constraints on the substrate by considering an FeO √ lattice√rotated 19° relative to the Pt(111) lattice, producing a ( 7 7)R19° FeO unit cell that is nearly commensurate with a p(33) Pt(111) unit cell and still exhibits local features similar to the three high-symmetry domains. The distinctive feature observed in STM simulations based on this somewhat simplified model was the increase in relative apparent height of the TOP domain relative to the FCC and HCP domains between þ500 and þ1500 mV bias, a feature which was attributed to a state localized at TOP domains of primarily Fe-3dz2 character approximately 700 meV above the Fermi 2097



that the Fe- and O-specific images, respectively, are formed with tips terminated by electronegative and electropositive atoms picked up from the surface while the “mixed mode” images are formed with ordinary metallic tip apexes, such as Pt. More elaborate theoretical treatment of STM imaging on the FeO film is required to shed further light on this topic.

6. CONCLUSIONS We investigated the influence of tip state, bias voltage, and interface structure on the STM imaging of single-layer moirestructured FeO films on Pt(111) and compared the high-resolution STM images with simulated STM images of the film’s electronic structure, calculated by DFTþU, using the TersoffHamann model. We presented a simple scheme by means of which we are able to interpret the STM images of the FeO structure unambiguously with the help of distinct oxygen-vacancy dislocation loops, thereby eliminating ambiguity regarding the assignment of features in the experimental STM images. Two similar tip states, referred to as “mixed modes”, result in STM images which show strong qualitative similarity to the simulated STM images, indicating that these STM images are well accounted for by the simple theory of Tersoff and Hamann. The majority of previously published STM images in the literature indeed exhibit the characteristic features of these mixed mode tip states. Three other tip states have been presented which result in STM images showing either Fe or O exclusively and thus bear no resemblance to the simulated STM images. In order to account for these STM images recorded with the element-specific tip states, more elaborate STM simulations which take explicit account for the tip are requested. On the basis of the present thorough study of the influence of the tip states, we reviewed previously published results for the FeO/Pt(111) system and in certain cases found that the conclusions have to be revisited. We have shown that subtler features of the FeO film, such as the magnetic structure, cannot reliably be deduced from spin-integrated STM measurements, and to resolve this, direct measurements of the film’s magnetic structure, such as spin-polarized STM, should be made. ’ AUTHOR INFORMATION Corresponding Author

*Phone: þ45 8942 3604. Fax: þ45 8612 0740. E-mail: fbe@ inano.au.dk. Present Addresses §

Current address: Department of Chemical Engineering, Stanford University, Stanford, California 94305, United States. Current address: Division of Synchrotron Radiation Research, Department of Physics, Lund University, Box 118, S-221 00 Lund, Sweden.

)

energy (EF). The experimental data presented by Giordano et al. appeared to show a rather complex evolution of the observed surface corrugation from low to high positive bias, with the order of apparent heights changing from TOP < HCP < FCC at þ65 mV to TOP < FCC < HCP at þ500 mV to FCC ≈ HCP < TOP at þ1500 mV, according to the authors’ assignment of the domains. From a detailed examination of these STM images, we observe very strong similarity with our own STM data recorded in mixed mode MM2. However, to make STM images recorded by Giordano et al. consistent with those we presented here in the mixed tip mode, a reassignment of the domains by Giordano et al. is required such that FCC domains and HCP domains are reversed in the 65 and 1500 mV images and the order of apparent heights at 500 mV follows the order FCC < TOP < HCP. This revision of the interpretation of STM results by Giordano et al. in fact eliminates the FCC/HCP reversal between þ65 and þ500 mV, in much better agreement with their own calculations. The STM simulations based on DFTþU calculations and the Tersoff-Hamann model reported by Zhang et al.,33 using a more realistic unit cell for the thin FeO layer on Pt(111), showed that the LDOS is a sensitive function of the magnetic structure of the FeO film. Zhang et al.33 were able to reproduce the increase in brightness of the TOP domain reported by Giordano et al. by choosing a particular domain-type antiferromagnetic spin pattern, where the magnetic moments of 7 Fe atoms centered at TOP domains were oriented opposite to the magnetic moments of the rest of the Fe atoms. Comparing their STM simulations with the experimental data of Giordano et al.,32 Zhang et al. proposed a reassignment of the different domains reported by Giordano et al., consistent with our suggestions above. The main parameter here determining STM image evolution at positive bias, as reflected in all the DFT calculations carried out so far, is the position of the strong peak in the Fe-3dz2 PDOS at TOP domains. In our electronic structure calculations, this peak appears at ∼0.3 eV, while in the calculations of Zhang et al.33 the peak appears at ∼0.4 eV, and in the calculations by Giordano et al. it appears at ∼0.7 eV. The position of the Fe-3dz2 PDOS at the TOP domains thus appears to be strongly model dependent, and considering that different tip states suggest different relative contributions of these Fe-3dz2 states to STM imaging at low bias, which would correspond to different positions for the Fe-3dz2 feature, we find it difficult to make conclusions regarding the finer details of the FeO film’s electronic structure based on our recorded STM images. Considering that the “mixed mode” images are reasonably well accounted for by the Tersoff-Hamann theory, the other elementspecific imaging modes clearly indicate that the specific tip state plays a crucial role in the imaging mechanism. Considering the relatively large and uniform corrugation values observed in these modes and the stark differences to the mixed modes, it appears likely that direct chemical interactions between the tip and the surface atoms are responsible for the image contrast. This is also consistent with the observation that stable imaging is only obtained within a relatively narrow range of bias and current settings. Previous STM studies on oxidized transition-metal surfaces24,49 and metal alloy surfaces25 as well as non-contact atomic force microscopy studies on insulating surfaces50,51 have noted similar chemical specificity in atomically resolved images and attributed this to differences in the chemical affinity of the atoms of the surface for the atom forming the tip apex. Following this reasoning, we suggest

ARTICLE

’ ACKNOWLEDGMENT Work at iNANO was supported by the Danish Research Agency, the Strategic Research Council, the Danish Council for Independent Research, as well as the Carlsberg and Villum Kahn Rasmussen Foundations and the European Research Council. Work at the University of Wisconsin;Madison was supported by DOE-BES, Chemical Sciences Division. Supercomputing time was utilized at NERSC, PNNL, ORNL, and ANL, all supported by the US-DOE. 2098



’ REFERENCES (1) Ramesh, R.; Schlom, D. G. MRS Bull. 2008, 33, 1006. (2) Freund, H.-J.; Pacchioni, G. Chem. Soc. Rev. 2008, 37, 2224. (3) Netzer, F. P.; Allegretti, F.; Surnev, S. J. Vac. Sci. Technol., B 2010, 28, 1. (4) Freysoldt, C.; Rinke, P.; Scheffler, M. Phys. Rev. Lett. 2007, 99, 086101. (5) Li, F.; Parteder, G.; Allegretti, F.; Franchini, C.; Podloucky, R.; Surnev, S.; Netzer, F. P. J. Phys.: Condens. Matter 2009, 21, 134008. (6) Schoiswohl, J.; Surnev, S.; Netzer, F. P.; Kresse, G. J. Phys.: Condens. Matter 2006, 18, R1. (7) Sedona, F.; Rizzi, G. A.; Agnoli, S.; Llabres i Xamena, F. X.; Papageorgiou, A.; Ostermann, D.; Sambi, M.; Finetti, P.; Schierbaum, K.; Granozzi, G. J. Phys. Chem. B 2005, 109, 24411. (8) Vurens, G. H.; Salmeron, M.; Somorjai, G. A. Surf. Sci. 1988, 201, 129. (9) Weiss, W.; Ranke, W. Prog. Surf. Sci. 2002, 70, 1. (10) Fu, Q.; Li, W. X.; Yao, Y. X.; Liu, H. Y.; Su, H. Y.; Ma, D.; Gu, X. K.; Chen, L. M.; Wang, Z.; Zhang, H.; Wang, B.; Bao, X. Science 2010, 328, 1141. (11) Sun, Y. N.; Qin, Z. H.; Lewandowski, M.; Carrasco, E.; Sterrer, M.; Shaikhutdinov, S.; Freund, H. J. J. Catal. 2009, 266, 359. (12) Kim, Y. K.; Zhang, Z.; Parkinson, G. S.; Li, S.-C.; Kay, B. D.; Dohnalek, Z. J. Phys. Chem. C 2009, 113, 20020. (13) Nilius, N.; Rienks, E. D. L.; Rust, H. P.; Freund, H.-J. Phys. Rev. Lett. 2005, 95, 066101. (14) Lin, X.; Nilius, N. J. Phys. Chem. C 2008, 112, 15325. (15) Berdunov, N.; Mariotto, G.; Balakrishnan, K.; Murphy, S.; Shvets, I. V. Surf. Sci. 2006, 600, L287. (16) Rienks, E. D. L.; Nilius, N.; Giordano, L.; Goniakowski, J.; Pacchioni, G.; Felicissimo, M. P.; Risse, T.; Rust, H.-P.; Freund, H.-J. Phys. Rev. B 2007, 75, 205443. (17) Nilius, N. Surf. Sci. Rep. 2009, 64, 595. (18) Hofer, W. A. Prog. Surf. Sci. 2003, 71, 147. (19) Hofer, W. A.; Foster, A. S.; Shluger, A. L. Rev. Mod. Phys. 2003, 75, 1287. (20) Blanco, J. M.; Flores, F.; Perez, R. Prog. Surf. Sci. 2006, 81, 403. (21) Tersoff, J.; Hamann, D. R. Phys. Rev. B 1985, 31, 805. (22) Barcaro, G.; Agnoli, S.; Sedona, F.; Rizzi, G. A.; Fortunelli, A.; Granozzi, G. J. Phys. Chem. C 2009, 113, 5721. (23) Schmid, M.; Kresse, G.; Buchsbaum, a.; Napetschnig, E.; Gritschneder, S.; Reichling, M.; Varga, P. Phys. Rev. Lett. 2007, 99, 196104. (24) Ruan, L.; Besenbacher, F.; Stensgaard, I.; Laegsgaard, E. Phys. Rev. Lett. 1993, 70, 4079. (25) Varga, P.; Schmid, M. Appl. Surf. Sci. 1999, 141, 287. (26) Merte, L. R.; Knudsen, J.; Grabow, L. C.; Vang, R. T.; Lægsgaard, E.; Mavrikakis, M.; Besenbacher, F. Surf. Sci. 2009, 603, L15. (27) Knudsen, J.; Merte, L. R.; Grabow, L. C.; Eichhorn, F. M.; Porsgaard, S.; Zeuthen, H.; Vang, R. T.; Lægsgaard, E.; Mavrikakis, M.; Besenbacher, F. Surf. Sci. 2010, 604, 11. (28) Lægsgaard, E.; Besenbacher, F.; Mortensen, K.; Stensgaard, I. J. Microsc. 1988, 152, 663. (29) Dudarev, S. L.; Botton, G. A.; Savrasov, S. Y.; Humphreys, C. J. Phys. Rev. B 1998, 57, 1505. (30) Kresse, G.; Furthm€uller, J. Comput. Mater. Sci. 1996, 6, 15. (31) Kresse, G.; Furthm€uller, J. Phys. Rev. B 1996, 54, 11169. (32) Giordano, L.; Pacchioni, G.; Goniakowski, J.; Nilius, N.; Rienks, E. D. L.; Freund, H.-J. Phys. Rev. B 2007, 76, 075416. (33) Zhang, W.; Li, Z.; Luo, Y.; Yang, J. J. Phys. Chem. C 2009, 111, 8302. (34) Bl€ochl, P. E. Phys. Rev. B 1994, 50, 17953. (35) Kresse, G.; Joubert, D. Phys. Rev. B 1999, 59, 1758. (36) Perdew, J. P.; Wang, Y. Phys. Rev. B 1992, 45, 13244. (37) Vurens, G. H.; Maurice, V.; Salmeron, M.; Somorjai, G. A. Surf. Sci. 1992, 268, 170. (38) Galloway, H. C.; Benitez, J. J.; Salmeron, M. Surf. Sci. 1993, 298, 127.

ARTICLE

(39) Schedel-Niedrig, T.; Weiss, W.; Schl€ogl, R. Phys. Rev. B 1995, 52, 17449. (40) Kim, Y. J.; Westphal, C.; Ynzunza, R. X.; Wang, Z.; Galloway, H. C.; Salmeron, M.; Van Hove, M. A.; Fadley, C. S. Surf. Sci. 1998, 416, 68. (41) Ritter, M.; Ranke, W.; Weiss, W. Phys. Rev. B 1998, 57, 7240. (42) Ranke, W.; Ritter, M.; Weiss, W. Phys. Rev. B 1999, 60, 1527. (43) Rienks, E. D. L.; Nilius, N.; Rust, H.-P.; Freund, H.-J. Phys. Rev. B 2005, 71, 241404(R). (44) Galloway, H. C.; Benitez, J. J.; Salmeron, M. J. Vac. Sci. Technol., A 1994, 12, 2302. (45) Galloway, H. C.; Sautet, P.; Salmeron, M. Phys. Rev. B 1996, 54, R11145. (46) Lemire, C.; Meyer, R.; Shaikhutdinov, S.; Freund, H.-J. Angew. Chem., Int. Ed. 2004, 43, 118. (47) Yao, Y.; Fu, Q.; Wang, Z.; Tan, D.; Bao, X. J. Phys. Chem. C 2010, 114, 17069. (48) Ma, T.; Fu, Q.; Su, H.-Y.; Liu, H.-Y.; Cui, Y.; Wang, Z.; Mu, R.T.; Li, W.-X.; Bao, X.-H. ChemPhysChem 2009, 10, 1013. (49) Calleja, F.; Arnau, A.; Hinarejos, J. J.; Vazquez de Parga, A. L.; Hofer, W. A.; Echenique, P. M.; Miranda, R. Phys. Rev. Lett. 2004, 92, 206101. (50) Foster, A. S.; Barth, C.; Shluger, A. L.; Reichling, M. Phys. Rev. Lett. 2001, 86, 2373. (51) Lauritsen, J. V.; Foster, A. S.; Olesen, G. H.; Christensen, M. C.; K€uhnle, A.; Helveg, S.; Rostrup-Nielsen, J. R.; Clausen, B. S.; Reichling, M.; Besenbacher, F. Nanotechnology 2006, 17, 3436.

2099


Tip-Dependent Scanning Tunneling Microscopy Imaging of Ultrathin

Recommend Documents