The Surface Structure of Cu2O(100): Nature of Defects

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The Surface Structure of CuO(100): Nature of Defects Heloise Tissot, Chunlei Wang, Joakim Halldin Stenlid, Tore Brinck, and Jonas Weissenrieder J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.8b05156 • Publication Date (Web): 18 Jul 2018 Downloaded from http://pubs.acs.org on July 25, 2018

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The Journal of Physical Chemistry

The Surface Structure of Cu2O(100): Nature of Defects

Heloise Tissot1, Chunlei Wang1, Joakim Halldin Stenlid2, Tore Brinck3, Jonas Weissenrieder1* 1

Materials and Nano Physics, School of Engineering Sciences, KTH Royal Institute of Technology, SE-100 44 Stockholm, Sweden

2

Department of Physics, Albanova University Center, Stockholm University, SE-106 91, Stockholm, Sweden

3

Applied Physical Chemistry, School of Chemical Science and Engineering, KTH Royal Institute of Technology, SE-100 44 Stockholm, Sweden

Corresponding author: [email protected]

1 ACS Paragon Plus Environment

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ABSTRACT The Cu2O(100) surface is most favorably terminated by a (3,0;1,1) reconstruction under ultrahigh vacuum conditions. As most oxide surfaces it exhibit defects and it is these sites that are focus of attention in this study. The surface defects are identified, their properties are investigated, and procedures to accurately control their coverage are demonstrated by a combination of scanning tunneling microscopy (STM) and simulations within the framework of density functional theory (DFT). The most prevalent surface defect was identified as an oxygen vacancy. By comparison of experimental results, formation energies, and simulated STM images, the location of the oxygen vacancies was identified as an oxygen vacancy in position B, located in the valley between the two rows of oxygen atoms terminating the unperturbed surface. The coverage of defects is influenced by the surface preparation parameters and the history of the sample. Furthermore, using low energy electron beam bombardment, we show that the oxygen vacancy coverage can be accurately controlled and reach a complete surface coverage (one per unit cell or 1.8 defects per nm2) without modification to the periodicity of the surface, highlighting the importance of using local probes when investigating oxide surfaces.


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INTRODUCTION Due to its excellent thermal and electrical conductivity, corrosion resistance and relatively low cost, copper is used for a wide range of industrial applications, including the fabrication of electrical circuits, conductors, and electronic equipment1 Copper has also been proposed as material in the enclosure of spent nuclear fuel in long-term repositories, and is used as a catalyst in, for instance, the low-temperature water-gas shift-reaction,2,3 and in methanol production.4 The oxides of copper are also active catalysts and are used in e.g. CO oxidation,5–7 propene epoxidation,8–11 and photocatalytic water splitting.12 Understanding the surface properties and chemical behavior of copper oxides is of importance for obtaining efficient catalysts, but also to prevent possible disastrous and costly corrosive degradation of copperbased materials. Despite the industrial importance of copper oxides, and in spite of numerous preceding studies, the nature of Cu2O surfaces remains poorly understood. Previous investigations include single crystals,13–19 oxidized polycrystalline copper,20 oxidized copper single crystals,21,22 and Cu2O nanoparticles.23 Experimental studies of single crystal Cu2O surfaces have been conducted using a range of different methods: low-energy electron diffraction (LEED),15 X-ray photoelectron spectroscopy (XPS),14 and scanning tunneling microscopy (STM).16 For example, the Cu2O(111) surface has been investigated with high resolution XPS and STM.16,17 On this surface, it was shown that oxygen vacancies are created upon annealing in ultra-high vacuum (UHV), resulting in a (√3 × √3) R30° reconstruction of the surface. This reconstruction is interpreted as a loss of one-third of the outermost oxygen ions. Surface reactivity towards water, methanol and SO2 have confirmed that defects play an important role for the reactivity of Cu2O(111).13,17,24 The Cu2O(100) surface has been subject for several theoretical studies in attempts at describing its structures and properties.25–29 However, determination of its surface structure 3 ACS Paragon Plus Environment


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proved more difficult than for Cu2O(111). First investigated experimentally using LEED, X-ray and ultraviolet photoelectron spectroscopy (XPS/UPS) by Schulz and Cox,15 a copper terminated (3√2 × √2) R45° reconstructed surface structure was identified, as well as a c(2 × 2)-structure with a half terminating layer of oxygen, and an oxygen terminated (1 × 1)-structure depending on the preparation conditions. In a previous paper, we revisited the atomic structure and electronic properties of the Cu2O(100) surface by a combination of multiple experimental techniques and simulations within the framework of density functional theory (DFT).18 Similarly to what was previously observed, three ordered structures were experimentally identiﬁed ((3,0;1,1), c(2 × 2), and (1 × 1)) depending on the sample preparation parameters. However, the copper terminated (3√2 × √2) R45° was not observed, but instead identified as a (3,0;1,1) reconstruction. The (3,0;1,1) reconstructed structure is the energetically most favorable termination under UHV conditions taking both DFT and experimental results into account. In addition to structural studies, the surface reactivity of Cu2O(100) has also been examined. The literature contains a large number of studies concerning adsorption on the Cu2O(100) surface of small molecules such as carbon monoxide,30 water,19,31 or sulfur dioxide.13 As mentioned previously, oxygen vacancies on the surface of Cu2O(111) and have proved to play a decisive role in the surface reactivity. However, the presence and nature of oxygen vacancies on the (100) surface and their influence on the surface reactivity have not been investigated to date. Furthermore, oxygen vacancies have been observed on other oxide surfaces such as TiO2.32–34 The nature of the defects observed on TiO2(110) have been the subject of intense debate. Indeed, as revealed by STM, the quick water dissociation on the surface oxygen vacancies in UHV induces the formation of single or paired hydroxyl groups, making the distinction between oxygen vacancies and OH groups difficult.35–38 In a reactivity point of view, oxygen vacancies greatly influence the surface reactivity towards O2 dissociative adsorption which selectively proceeds at bridging oxygen vacancies, with one O atom healing the vacancy.36,39 Moreover, it has been shown that the photocatalytic properties of TiO2-based metal oxides are closely related to oxygen vacancies.40,41 A complete picture of the Cu2O(100) structure and reactivity cannot be obtained without a precise understanding of the nature of surface defects and their coverage. In this work, we focus on the identification of these defects as well as the control of their coverage by means of


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several methods: STM, LEED, and DFT calculations. The STM results are, with good agreement, compared to DFT simulations. The most favorable site for vacancy formation is identified.

EXPERIMENTAL DETAILS The Cu2O(100) single crystal used in this study is a natural crystal acquired from the Surface Preparation Laboratory, The Netherlands. The cleaning procedure of the crystal consists of cycles of argon ion sputtering (0.5 kV, 20 min) followed by anneal in ultrahigh vacuum (at 640 °C) or in oxygen gas (at 500 °C, PO2 = 3 × 10-6 mbar). For the LEED and STM studies, the crystal cleanliness was confirmed by the sharpness of the diffraction spots in the LEED pattern. The STM-study was carried out using an Omicron VT-STM operated in constant current imaging mode with electrochemically etched tungsten tips. The STM-chamber is attached to a preparation chamber equipped with ion sputter gun, leak-valves for gases and LEED-apparatus. The presented LEED patterns were obtained using this LEED. The base pressure in both preparation chamber and analysis chamber is low 10−10 mbar.

COMPUTATIONAL DETAILS Spin-polarized DFT calculations at the PBE-D3+U level of theory42–45 were conducted with the Vienna Ab-initio Simulation package (VASP)46 using a Hubbard U–j value of 3.6 eV. The valence electrons (Cu: 3d104s1; O: 2s22p4) were treated explicitly by a plane-wave basis-set with a cut-off of 520 eV, while the core states were represented by PBE PAW47,48 (projector augmented wave method) potentials. The Brillouin zone was sampled using the tetrahedron method with Blöchl corrections49 and a Γ-centered 4×4×1 k-point mesh. The Cu2O(100) surface was modeled by the (3,0;1,1) structure of Soldemo et al.18 A six Cu2O layer thick asymmetric slab was used with a vacuum distance of 20 Å. The top two layers were allowed to relax and the forces were converged to 0.03 eV/Å. All results have been compared to results obtained with dipole corrections according to the methods of references

50

and 51 showing negligible differences.



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Surface oxygen vacancy formation energies (∆Evac) were evaluated by: ∆Evac = Esurf/vac + ½EO2 – Esurf

(1)

where Esurf/vac, EO2 and Esurf represent the electronic energies of the oxygen vacant surface, O2 in gas phase and the stoichiometric surface respectively. Corrections to Gibbs free energies were obtained by adding vibrational contributions via the harmonic approximation and gas phase rotational and translational corrections via the rigid rotor and ideal gas approximations to the electronic energies as described in Stenlid et al.19 The electronic energy of O2 was described through the equilibrium with H2O and H2 due to the known difficulty of correctly describing ground-state O2 with plane-wave DFT. A partial pressure of 10-15 bar was assumed for O2(g) throughout. The electronic density of states (DOS) calculations on bulk Cu2O with and without O vacancies were conducted using a 18×18×18 k-point mesh and the screened hybrid functional HSE06.52–55 Simulated STM images were generated with the HIVE program,56 using the TersoffHamann57 method as described in Soldemo et al.18 The surface was represented by a p(3 × 3) supercell in the simulated STM in order to allow for a range of low (1/3) to high (full) vacancy coverage.

RESULTS Figure 1 a-c summarize our current understanding of the (3,0;1,1) surface structure for UHV annealed Cu2O(100). It presents a high resolution STM image (Figure 1a), its corresponding DFT simulated image (Figure 1b) and an atomic structure model (Figure 1c).18 All three images show protrusions aligned in rows along the [011] direction attributed primarily to surface oxygen atoms at the top of the rows (Figure 1c, bright red atom labeled A), representing the main contribution for tunneling voltages in excess of 2 V. Due to experimentally undefined tip conditions, the atomic rows can either appear very pronounced (see Figure 2a) or can present a honeycomb-like structure as observed Figure 1a. The surface unit cell is defined in Figure 1 with its corners positioned at the oxygen atoms in the rows of the surface (bright red, labeled A). There are two additional protrusions


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inside the unit cell, the first is located on the bottom left of the unit cell, at a position close to an oxygen atom that is partly concealed by copper atoms at the edge of the row (red atom, labeled C); and the second protrusion in the center of the cell, brighter in the DFT simulated image than on the experimental, associated to another oxygen atom exposed at the bottom of the valley (dark red atom, labeled B) and to surface copper atoms. The experimental and by DFT simulated images are in reasonable agreement. Some discrepancies can be expected since the bias voltage used in the simulated image may not be perfectly aligned with the experiments. Band gaps of transition metal semiconductor oxides are known to be notoriously difficult to estimate by generalized gradient approximation DFT (e.g. PBE), thereby inducing a shift in the STM voltage when comparing DFT and experimental images. Further, since DFT simulated images are obtained by the Tersoff-Hamann approximation, they represent an intrinsic property of the unperturbed surface rather than a property of the joint surface-tip system as in the experimental case. Whereas the latter point is likely to introduce some differences between the experimental and DFT images. Tests at different biases gives us confidence that the errors due to the first point raised can be alleviated by a careful selection of simulation conditions: experimental STM images remain very similar between 2 and 3 V, as do DFT images in a range from approximately 2 to 4 V. We can, furthermore, note that adjusting the DFT simulated voltage towards lower biases (by 1 - 2 V) amplifies the relative contribution of the Cu-dominated spots in the center of the cell at the expense of the ridge oxygen atoms as seen in Figure S1 of the supporting information. In other words, we probe Cu dominated areas at low voltage, and O dominated areas at higher voltage. This would shift the protrusions by half a unit cell in the [0 1 1]-direction going from low to high biases, but cannot be confirmed by in STM images due to the experimental difficulties in fixating the coordinate system.



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Figure 1: a) STM image (2.1 V, 3.5 nA) of an as-prepared Cu2O(100) (3,0;1,1) surface, b) DFTsimulated STM image (Bias 3 V distance 2.5 Å), c) Atomic structure model deduced from DFT calculations. The three unique oxygen atoms are labeled as A (bright red), B (deep red) and C (semi-dark red). Oxygen atoms in layers deeper in the bulk are red, whereas surface Cu are shaded to distinguish them from the bulk Cu atoms. The surface unit cell is represented in blue on each image. Larger STM images show, in addition to the clean unperturbed surface structure, a low concentration of protrusions associated to surface defects. More precisely, as will be discussed further on, the protrusions can be attributed to oxygen vacancies (Figure 2). As presented in Figure 2c, the observed defects are located in-between two oxygen rows of the surface. The dimension of a protrusion is of ~ 7 Å along the oxygen rows and ~7 Å perpendicular to these rows (average value obtained for dimension measurement on 10 isolated protrusions). Scanning Tunneling Spectroscopy (STS) analysis performed on the defects sites (the curve presented is an average of 7 measurements) show states in the band gap at approximately 0.5 - 1 V above the Fermi level (red curve in Figure 2b) while STS performed on the non-defective rows does not present any states at similar bias (blue curve). This is in line with the expected electronic effect of introducing oxygen defects in bulk Cu2O seen in valence density of states obtained at the HSE06 hybrid functional level of theory (see Discussion section below). It is also in agreement with the above mentioned finding that the lower bias states of the conduction band originate from surface Cu (SI).


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Figure 2: a) STM image (2.4 V, 0.12 nA) and LEED pattern of the Cu2O(100) clean surface. b) STS analysis on a surface defect (red) and on a clean row (blue) c) STM Image (2.1V, 0.35 nA) and d) height profile along the blue line in c. Oxygen vacancies observed on TiO2(110) have also been reported to induce band gap states, as observed by XPS/UPS. The theoretical picture regarding the origin of the band-gap state is not completely settled: most calculations predict that oxygen vacancies should lead to the formation of the band-gap state

34,58

, while a recent spin-polarized hybrid DFT calculation also

suggests that Ti3+ interstitials can contribute to the band-gap state.59,60 On the other hand, UPS in combination with scanning tunneling microscopy corroborates that the band-gap state in TiO2(110) originates mainly from bridging oxygen vacancies by showing that the population of the band gap state is related to the density of oxygen vacancies.33 Figure 3 presents a sequence of images recorded on the same area of the surface. The defect consisting of two holes at the bottom right of each image in Figure 3 (indicated by the black rectangle) is used as a reference position for sample drift The series of images show that the protrusions are mobile at room temperature (blue circles and arrows). The defects either exhibit thermally excited diffusion at room temperature or alternatively the vacancy hopping can be tip induced, due to the potential applied between surface and STM tip.17 On these images, the 9 ACS Paragon Plus Environment


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surface reconstruction is cleanly resolved but some bright protrusions appear blurry or cut in half suggesting tip effects. However, in both cases, this observation suggests a relatively flexible surface lattice.

Figure 3: STM images from the same surface area of the Cu2O(100) clean surface with defects (2.4 V, 0.12 nA). The defect in the at the bottom right of each image (indicated by a black square) serves as a reference position for sample drift. Arrows and circles indicate regions with mobile protrusions/defects. The mobility of vacancies has previously been investigated on titanium and cerium oxide surfaces. Namai et al.61–64 showed that surface oxygen vacancies on CeO2(111) are mobile at room temperature. Similar high mobility of oxygen vacancies was also reported on both the (√3 × √3)R30°-Cu2O(111)16 and TiO2(110)65 surfaces. Interestingly, vacancies tend to cluster into linear chains on TiO2(110), a similar behavior can be observed for the defects on the Cu2O(100) surface in Figure 4, suggesting a non-random defect distribution.


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A careful investigation of the vacancies on Cu2O(100) and their origin shows that their coverage depends largely on preparation conditions of the surface. Repeated cycles of sputtering and annealing increase the coverage of defects as well as extended periods at high annealing temperature in UHV. Figure 4 shows three images obtained after one cleaning cycle with an annealing of 10 min at 650°C (Figure 4a), after a second cycle with an annealing of 40 min at 700°C (Figure 4b) and after several cycles of annealing at 700°C for 20 min (Figure 4c). The defect coverage increases with annealing time from 0.025 defects per nm2 (~ 0.012 defects per (3,0;1,1) surface unit cell) in Figure 4a, to 0.05 defects per nm2 (~0.025 per surface unit cell) in Figure 4b and finally to 0.24 defects per nm2 (~0.12 per surface unit cell) in Figure 4c. The results indicate that annealing in UHV favors the formation of defects on the surface.

Figure 4: STM images of the Cu2O(100) clean surface (2.4 V, 0.12 nA) with increasing density of defects after a) annealing at 650 °C for 10 min, b) at 650 °C for 40 min and c) at 700°C for 20 min. Annealing the sample can potentially promote segregation of species to the surface (either defects or impurities in the bulk) and a tentative attribution for these defects is the presence of a contaminant that diffuses to the surface during annealing, as for example potassium atoms. The sample used in the current experiments is, as mentioned in the experimental section, a natural crystal, and as such contain trace levels of potassium that tentatively may segregate and enrich at the surface by repeated cycles of sputtering and annealing. However, no potassium enrichment was observed by XPS (not shown here). On the TiO2(110) surface, the density of oxygen vacancies can be controlled by using low energy electron bombardment and such preparation procedure have been used to study the electronic structure of the defect and to 11 ACS Paragon Plus Environment


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understand their reactivity.32,66,67 To experimentally confirm our assignment of the surface defects on the Cu2O(100) surface as oxygen vacancies, we performed similar low energy electron bombardment experiments as this process will favor formation of localized surface defects. There is no plausible processes for electron bombardment induced surface enrichment through segregation from the bulk. STM images of the Cu2O(100) surface for different exposure times to the electron beam are presented Figure 5. The electron beam has spot at the sample with a diameter of approximately 1 mm and the sample drain current was 5 µA. Figure 5a represents the as prepared clean surface. As observed previously, defects are already present after the preparation procedure. From Fig 5b to 5d the exposure time to the electron beam was increased leading to an increase in the surface defect concentration until almost full surface coverage is reached in Figure 5d. This observation confirms that defects do not originate from bulk impurities and shows, as observed on TiO2, that the defects are oxygen vacancies. For each preparation, the LEED pattern remains the similar and exhibit a (3,0;1,1) reconstruction, in agreement with STM observation. Consequently, the observed defects do not modify the symmetry and the dimensions of the surface unit cell. It is also interesting to note the close to complete coverage of the surface by defects, as observed Figure 5d, suggest that each surface unit cell contains one defect. The theoretical complete coverage if one oxygen vacancy is assumed to be formed in each surface unit cell would be 1.8 defects per nm2. This observation points out a significant difference compared to the Cu2O(111)16 surface for which the maximum coverage of surface oxygen vacancies represents one-third of the outermost oxygen ions and induces a different LEED pattern associated to a (√3 × √3)R30° surface reconstruction.


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Figure 5: a) STM images of the as prepared Cu2O(100) clean surface (2.4 V, 0.12 nA), b) after 2 min (electron energy = 80 V, sample current = 5 µA), c) 4 min (80 V, 5 µA), and d) 6 min electron bombardment (80 V, 5 µA). Figure 6 shows two high resolution images for an intermediate and close to complete coverage of defects obtained after electron bombardment of the surfaces. Similarly to what was observed previously, the protrusions are placed in-between two oxygen rows, centered in the trough. Analysis at low defect coverage allows determination of the dimensions of a single protrusion. They are found to be ~ 7 Å both along the oxygen rows and perpendicular to the rows, the same as what was found for the isolated protrusion in Figure 2. This indicate that the nature of observed defects is the same. Figure 6a also supports the hypothesis of non-random coverage and of a preferential distribution forming rows and islands of defects. In order to confirm the non-random distribution, the sample was annealed in UHV until 600°C but no difference was observed on the number or the distribution of defects (see Supporting Information Figure S2) suggesting that ordering may take place at room temperature. In Figure 6b, the structure of the unperturbed surface can be observed between the defects as well as a few darker spots on the rows indicating formation of another type of defect.



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Figure 6: High resolution STM images of the Cu2O(100) surface (2.4 V, 0.12 nA) after low electron bombardment for a) 2 min (80 V, 4.2 µA) and b) 6 min (80 V, 5 µA). In an attempt to “heal” the oxygen vacancies formed by electron bombardment, we annealed the surface in oxygen (p(O2) = 10-6 mbar, 10 min at 600°C). The number of oxygen vacancies decreased after the O2 exposure. Figure 7a presents a STM image prior to annealing in O2 with a defect coverage of 0.6 defects per nm2 (~ 0.3 defects per surface unit cell). After annealing in O2 the surface defect coverage has decreased with one third to a surface coverage of 0.4 defects per nm2 (~0.2 defects per surface unit cell). However, at p(O2) = 1 × 10-6 mbar, the annealing does not lead to a complete healing of all oxygen vacancies (Figure 7b). A significantly higher pressure of O2 is required for a complete healing, e.g. on surfaces exposed to the ambient air and only one cycle of sputtering-annealing defects are rare on the surface (≤ 0.025 defects per nm2 or 0.012 per unit cell). On the Cu2O(111), annealing in ultra-high vacuum form a distinct (√3 × √3) R30° reconstruction while annealing in O2 partially restores the stoichiometric (1 × 1) oxygen terminated surface, however, after extensive use the crystal cannot be prepared with a complete (1 × 1) termination at O2 pressures accessible in an UHV chamber. Islands of (√3 × √3) R30° will always be present.16


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Figure 7: High resolution STM images of the Cu2O(100) surface (2.4 V, 0.12 nA). a) after low electron bombardment for 6 min (80 V, 4.8 µA), defect coverage of 0.6 defects per nm2 (~ 0.3 defect per surface unit cell ((3,0;1,1) cell)) and b) after annealing in O2 at 10-6 mbar during 10 min at 600°C, defect coverage of 0.4 defect per nm2 (~ 0.2 defect per (3,0;1,1) surface unit cell. In order to gain further understanding of the oxygen vacancies, we have also employed DFT calculations. There are three unique oxygen atoms in each (3,0;1,1) unit cell, see Figure 1c. The energies for formation of oxygen vacancies at these different positions are included in Table 1. We first note that the vacancy formation energy is ordered as B>A>C for the various oxygen positions. Thus we find, somewhat surprisingly, the most favored position for formation of a vacancy is site B located at a valley site. This site is favored by 0.4 eV over the A site where an oxygen sits exposed at the top of the ridge (site A). Site A would be the naïve guess as the most favored site, similar to the exposed unsaturated O atoms (OCUS) of the Cu2O(111) surface. The finding that the B site is favored is, however, in line with the experimental observation that the vacancies appear in between the ridges. The least favored position for an O vacancy is at the C site, which is largely concealed under copper atoms at the base of the ridge. Simulated STM images at high and low defect concentrations are included in Figure 8 at +2 V bias (other biases are found in Figure S1 of the SI). The simulated images show a clear increase of intensity (i.e. accumulation of charge density relative other sites) in between the valleys in good agreement with the experimental images.



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Table 1. Oxygen vacancy positions, and electronic (∆Evac) and Gibbs free energies (∆Gvac) in eV. Site

Location

∆EVac

∆GVac

A

On ridge

2.39

1.05

B

Valley (cell center)

2.04

0.69

C

Base of ridge

2.63

1.28

DISCUSSION In the present study, the atomic structure of the Cu2O(100) surface and the nature of defect formed on its surface are studied by LEED, STM, and DFT. The influence of the sample preparation conditions as well as electron bombardment and O2 exposure on the surface atomic structure is investigated in order to establish an understanding of the nature of the defects and of the conditions that favor their formation. The defects can, in principle, have many different origins; they can be due to water adsorption and OH formation (from the residual gas in the chamber), as well as oxygen or copper vacancies. In the case of water adsorption, Deng et al. have shown that hydroxyl formation on copper oxide occurs for H2O pressures higher than 10-7 mbar at room temperature.20 Stenlid et al. observed hydroxyls groups on the clean Cu2O(100) surface by XPS after exposure to a dose of 3 L H2O at 117 K and step-wise annealing to room temperature. However, no peak corresponding to hydroxyl groups (at 531.5 eV) was observed in UHV at room temperature.19 Consequently, the defects observed here do not come from residual water adsorption. A more likely explanation is oxygen vacancies. Such defects have long been thought to dominate the reactivity of oxide surfaces and therefore much research is directed towards understanding their properties. Comparison with results obtained on Cu2O(111), titanium, and cerium oxide leads to the conclusion that these defects are oxygen vacancies. This implies that defect protrusions in the STM images are due to coordinatively unsaturated copper ions, which is quite reasonable since broken bonds leave additional charge on these ions resulting in a larger tunneling current from occupied states in the valence band. This hypothesis is further corroborated by the DFT results presented herein. The results can also be compared with the Cu2O(111) surface, where the (√3 × √3)R30° reconstructed surface is obtained after a few cycles 16 ACS Paragon Plus Environment

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of ion bombardment and annealing in UHV and consists of wide terraces covered by protrusions. These protrusion resemble the shape and size of the protrusion observed on the Cu2O(100) and represent the removal of one-third of the surface oxygen anions.16 Furthermore, the annealing in O2 can partially restore the oxygen-covered surface yielding a (1 × 1) surface structure. Cu2O(100) seem to present a similar behavior, annealing and sputtering cycles induce the formation of surface defects without reaching one-third of the surface oxygen. However, after low energy electron bombardment, the surface can be fully covered by defects and partially recovered by annealing in O2, giving further evidence that the origin of a protrusion is indeed an oxygen vacancy. One final question remains to be considered; the bright spots observed on the STM images are attributed to oxygen vacancies, but which oxygen from the surface unit cell has been removed? As already discussed, the surface unit cell contains three types of surface oxygen atoms, one fully exposed and two partly concealed by copper atoms. The location of the protrusion (in between two rows) as well as the unperturbed (3,0;1,1) LEED pattern of the fully covered surface by defects, seem to exclude the oxygen atom which is located at the edge of the unit cell on top of the rows. However, no experimental argument allows us to distinguish between the other two. Based on the DFT results we can, however, argue that vacancy formation at the center of the unit cell, in between the ridges, is indeed energetically favorable compared to formation of vacancies at the ridges. Specifically we find that formation of a vacancy at position B (Figure 1) is likely. As shown in Figure 8, calculated electronic DOS of Cu2O bulk indicate that O vacancies introduce band-gap states consistent with the experimental observations. Simulated STM images further confirms that this kind of vacancies could lead to bright spots in the positions observed experientially.



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Figure 8: a) Surface model for the reconstructed Cu2O(100) surface. In b) is the simulated STM image for the surface at a coverage of 1/3 oxygen vacancies (OVac) per unit cell at the B site shown. Simulated images at other biases and at full OVac coverage are included in the supporting information. c) Comparison of the simulated density of states for Cu2O (black solid line), oxygen vacant Cu2O (red solid line), experimental STS analysis from the unperturbed surface (black dashed line) and from oxygen vacancies (red dashed line). CONCLUSIONS In order to gain a detailed understanding of the Cu2O(100) structure and reactivity, we have investigated the formation of surface defects and their properties. We investigated the nature of surface defects and their coverage on the surface by a combination of STM and DFT simulations. The most prevalent defects are identified as oxygen vacancies. STS analysis and DFT simulated density of states reveal band gap states induced by the surface defects. Furthermore, the oxygen vacancy coverage depends on preparation conditions, mainly on annealing time and temperature. However, using low energy electron bombardment, we showed that the oxygen vacancy coverage can be controlled and reach a complete surface coverage (one oxygen vacancy per unit cell) without any modification of the surface periodicity. This highlights the importance of using a local probe when investigating oxide surfaces. Formation energies of oxygen vacancies and simulated STM images obtained by DFT calculations allows us to identify the missing oxygen atom placed in position B in between two oxygen rows. On-going studies are conducted to further understand how the oxygen vacancy coverage influence the surface reactivity.

ACKNOWLEDGEMENT This work was funded by the Swedish Research Council (VR), the Knut och Alice Wallenbergs stiftelse, the Ragnar Holm foundation for Heloise Tissot’s fellowship, and the Trygger’s foundation for Chunlei Wang’s fellowship. The Swedish National Infrastructure for Computing (SNIC) is acknowledged for providing computational resources at the National Supercomputer 18 ACS Paragon Plus Environment

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Centre in Linköping University NSC as well as at the PDC Centre for High Performance Computing (PDC-HPC).

SUPPORTING INFORMATION DESCRIPTION The influence of annealing on the organization of surface defects and additional DFT simulations are presented in the supporting information.

REFERENCES (1) (2) (3) (4) (5) (6) (7) (8) (9) (10) (11) (12) (13)

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The Surface Structure of Cu2O(100): Nature of Defects

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