Surface oxygen chemistry of a gas-sensing material: SnO2(101)

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Jan 1, 2004 - Experimental techniques and density-functional theory have been employed to identify the surface composition and structure of SnO2(101).
EUROPHYSICS LETTERS

1 January 2004

Europhys. Lett., 65 (1), pp. 61–67 (2004) DOI: 10.1209/epl/i2003-10044-0

Surface oxygen chemistry of a gas-sensing material: SnO2 (101) M. Batzill 1 , A. M. Chaka 2 and U. Diebold 1 (∗ ) 1 Department of Physics, Tulane University - New Orleans, LA 70118, USA 2 Chemical Science and Technology Laboratory, National Institute of Standards Gaithersburg, MD 20899-8380, USA (received 29 July 2003; accepted in final form 14 October 2003) PACS. 68.47.Gh – Oxide surfaces. PACS. 82.65.+r – Surface and interface chemistry; heterogeneous catalysis at surfaces. PACS. 68.37.Ef – Scanning tunneling microscopy (including chemistry induced with STM).

Abstract. – Experimental techniques and density-functional theory have been employed to identify the surface composition and structure of SnO2 (101). The stoichiometric Sn4+ O2 2− surface is only stable at high oxygen chemical potential. For lower oxidizing potential of the gas phase a Sn2+ O2− bulk termination is favored. These two surfaces convert into each other without reconstruction by occupying and vacating bridging oxygen sites. This variability of the surface composition is possible because of the dual valency of Sn and may be one of the fundamental mechanisms responsible for the performance of this material in gas-sensing devices.

Stannic oxide (SnO2 ) is used as an oxidation catalyst and is sensitive to the adsorption of oxidizing and reducing gases, which makes it a widely used material for metal-oxide–based gas-sensing devices. In gas sensors a change in the electrical resistance of polycrystalline SnO2 films is detected if the composition of the ambient gas changes. The sensitivity of a material to the adsorption of different gases is likely to be manifold. One often discussed contribution to the gas-sensing mechanism is the variation of Schottky barriers at contacts between grains in polycrystalline films [1, 2]. The barrier height determines the transmission probability of conduction electrons and thus the conductivity of the film. The Schottky barrier, i.e. the band bending at the surface of SnO2 grains, is a consequence of surface states and charged adsorbates at the surface and is thus influenced by the environment. Here the dependence of the lattice oxygen concentration at the SnO2 (101) surface on the oxygen chemical potential of the gas phase is examined. This is an important first step that eventually may lead to a better fundamental understanding of gas-sensing materials. The single-crystal SnO2 (101) surface is considered here as a model system for gas-sensing materials. Although the (101) surface is the second most stable SnO2 surface and thus is expected to be of high abundance in polycrystalline films, the relative importance of different crystal faces for the gas-sensing response has not yet been established. Nevertheless, the reduced complexity of single-crystal surfaces allows focusing on the oxygen chemistry and (∗ ) E-mail: [email protected] c EDP Sciences 

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structure under well-defined conditions. Here, we demonstrate that the SnO2 (101) surface is particularly appealing for such studies because it is stable not only in its stoichiometric termination but also in an oxygen-depleted termination with a SnO surface layer, without formation of complex surface reconstruction that are, for example, encountered on the lowerenergy (110) surface. This implies that the composition of the SnO2 (101) surface is variable over a large range, depending on the equilibrium conditions with the gas phase. Recently, density-functional theory calculations have shown similar variations of the surface composition for other oxide materials [3]. These findings show that a change of the oxidation potential of the gas phase can alter the composition of surfaces. The surface stoichiometry determines the charge distribution (surface states) and thus the Schottky barrier height and therefore has a strong impact on the gas-sensing mechanism. Previous studies of the SnO2 (101) surface are rare. First-principles density-functional theory using the generalized gradient approximation found that the energies of stoichiometric surfaces increase in the sequence (110) (1.04 Jm−2 ), (100) (1.14 Jm−2 ), (101) (1.33 Jm−2 ), and (001) (1.72 Jm−2 ) under vacuum conditions, but provided no information as to relative stabilities in response to changes in environmental conditions [4]. There exist relatively few experimental surface studies on SnO2 single crystals. Most of these were performed on the (110) surface (see [5,6] and references therein), which exhibits complex surface reconstructions. These are still not fully understood, but a consensus exists that this surface is oxygen depleted if annealed in vacuum [7]. The only previous experimental study performed on the (101) (or (011) surface, which is equivalent for the rutile structure) showed that the only ordered LEED pattern has a (1 × 1) symmetry [8], thus indicating that this surface is more stable against reconstruction than the (110) surface. Here the stability of the (101) surface is explained by the dual valency of Sn that allows formation of stable bulk truncations for surfaces with stoichiometric and reduced compositions. The experiments were performed in an UHV apparatus described elsewhere [6]. Additionally, a high-pressure cell was attached for exposure of samples to high-pressure oxygen. Scanning tunneling microscopy (STM) images were obtained under empty-state imaging conditions with bias voltages in a range from 1 to 2 V. Helium ion scattering spectroscopy (ISS) was performed with a primary ion energy of 1200 eV and a scattering angle of 137◦ . The ion current on the sample was ∼ 15 nA. The azimuth of the incident ion beam was closely aligned along the [0 −1 0] direction of the (101) surface. For temperature-programmed ISS a 0.2 K/s temperature ramp was applied to the sample and a spectrum was obtained every ∼ 10 K. Total energy calculations and full relaxations were performed using all-electron firstprinciples density-functional theory, the PBE [9] generalized gradient approximation to the exchange-correlation functional and atom-centered, polarized, split-valence numerical basis sets (DNP) as implemented in the Dmol3 program [10]. Spin-restricted calculations were run on slabs containing fifteen atomic layers for the O-terminated surface and thirteen atomic layers for the Sn-terminated surface with 10 ˚ A vacuum separation in the supercells. A uniform k-point mesh with 18 points was used for the entire surface Brillouin zone. Mulliken population analysis, although not an absolute measure of the exact charge distribution, does yield a consistent measure of relative charge and was used to determine relative oxidation states of atomic species. The 0 K total energy calculations are linked to the finite temperatures and pressures encountered in the experiments through the recently developed method of ab initio thermodynamics [3]. The change in chemical potential ∆µi for species i as a function of temperature for Sn metal, SnO2 (c), and O2 (g) were taken from thermodynamic reference tables [11,12]. The change in chemical potential for the slabs above 0 K was assumed to be comparable to bulk SnO2 . The difference in vibrational entropy between the surface and the bulk phonons for a metal oxide (RuO2 ) when no hydrogen is present has been esti-

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Fig. 1 – (a) 1 × 1 LEED pattern of the SnO2 (101) surface. The missing (±1, 0) spots indicate the glide plane symmetry of this surface. (b) Large-scale STM image (140 × 65 nm2 ). The step edges have a preferential orientation along the crystallographic [−101] direction. (c) STM image of a two-unitcell-wide terrace. The step edges along [−101] exhibit a jagged structure with protrusions separated by the unit cell length. (d) Atomically resolved STM image (3.4 × 3.4 nm2 ). The unit cell is indicated by the white rectangle.

˚2 to the surface free energy in absolute terms, mated to contribute less than +/−10 meV/A and considerably less in relative surface free energies [3]. Samples prepared by sputtering and annealing to 900 K in UHV show sharp (1 × 1) LEED patterns, similar to the one shown in fig. 1(a). Missing (2n − 1, 0) spots indicate glide plane symmetry. STM images of such surfaces are shown in fig. 1(b)-(d). In high-resolution STM images the (1×1) unit cell can be easily discerned (fig. 1(d)). The main features are five bright protrusions; four protrusions in the corner of a rectangle of the expected unit cell dimensions and one protrusion in the center. A slight contrast difference between the left and right side of this unit cell, however, indicates that this is not a centered unit cell but represents a primitive unit cell. On a large scale this surface exhibits a high density of step edges (fig. 1(b)). The step edges are preferentially oriented along the crystallographic [−101] direction, making the surface strongly anisotropic. Long and narrow terraces are frequently observed on this surface. Some of these terraces are only two unit cells wide along the [0 −1 0] direction and extend over several tens of nanometers along the [−101] direction (fig. 1(c)). The step edges along the [−101] direction have a jagged appearance with protrusions separated by the unit cell length. Assuming that Sn atoms are imaged bright under empty-state tunneling conditions [13], it is concluded that the step edges are terminated by protruding Sn atoms. The two possible bulk terminations with a (1 × 1)-pg unit cell are shown in fig. 2. The stoichiometric SnO2 (101) surface is shown in fig. 2(a) and (b). It consists of fivefold coordinated Sn atoms and of alternating rows of twofold coordinated bridging O atoms and threefold

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Fig. 2 – Ball-and-stick models for a stoichiometric ((a) and (b)) and an oxygen-depleted ((c) and (d)) SnO2 (101)-1×1 surface (small white balls indicate Sn, and large dark balls indicate O atoms). (a) and (c) show pseudo-3D views of the surface, and (b) and (d) show a top view of these two surfaces. The unit cells (solid line) and the glide plane symmetry axes (dashed line) are indicated. The numbers in (a) and (b) give calculated relaxations between subsequent layers, i.e. O-Sn, Sn-O, O-O, O-Sn, etc. for the stoichiometric and reduced surfaces with respect to bulk values.

coordinated O atoms situated slightly below the plane of the Sn atoms. A reduced surface with all the bridging oxygen atoms removed results in a Sn-terminated surface as is illustrated in fig. 2(c) and (d). It is difficult to discriminate these two surface terminations by qualitative LEED and in atomic resolution STM images. This is because the symmetry of the unit cell is unchanged and generally the oxygen atoms do not contribute largely to the contrast in empty-state STM images of SnO2 [13]. The preferential step edge orientation along the [−101] direction observed in STM images may however help to discern the surface termination. A simple counting of the number of Sn to O bonds that need to be broken to form step edges along the two lowindex crystallographic directions for a stoichiometric and an oxygen-depleted surface allows an estimate of the step edge energy. On a stoichiometric surface two bonds per unit cell need to be broken to form step edges along either the [−101] or the [0 −1 0] direction. On a surface that is oxygen depleted, on the other hand, one bond needs to be broken for steps along the [0 −1 0] direction but no bonds are broken normal to the [−101] direction if the lower terrace is also tin-terminated. Therefore, we expect [−101] step edges only to be preferred for tin-terminated surfaces. An accurate determination of the surface composition is done by utilizing the extreme surface sensitivity of He-ISS. Figure 3 shows two ISS spectra, one for a sample prepared in UHV and another after exposing the same sample to 10 mbar O2 . Although the same LEED patterns are observed for both preparation conditions, a drastic change in the [O]/[Sn] ratio is observed in ISS. This indicates that the surface converts from a reduced to a stoichiometric surface upon oxygen exposure. Additionally to the single-scattering Sn peak, a second peak is observed at 956 eV kinetic energy for UHV-prepared samples. This peak disappears after

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Fig. 3 – (a) Two spectra taken after UHV preparation (lower spectrum) and after exposure to 10 mbar O2 (upper spectrum) are shown. The peaks are due to single scattering from oxygen and tin. For the UHV-prepared surface a double-scattering peak appears at 956 eV. The inset shows the scattering path of such an event. (b) Evolution of the peak intensities with annealing temperature of the sample. The decrease in the oxygen signal and concurrent increase in tin-signal indicates desorption of bridging oxygen at ∼ 680 K. (c) Free energies of SnO2 (101) surface terminations as a function of the oxygen chemical potential µO . The dashed vertical lines indicate the allowed range of µO . The value of µO terms of temperature at the experimentally employed UHV conditions (∼ 10−14 mbar O2 pressure) is indicated at the top of the figure. At a chemical potential of −1.8 eV (650 K for UHV conditions) the reduced surface becomes favored, in good agreement with the experiments shown in (b).

oxygen exposure. The position of this second peak is consistent with a double-scattering event from Sn atoms (see inset in fig. 3(a)). Occupation of the bridging oxygen rows inhibits this process and no double-scattering peak can be observed for the stoichiometric surface. In order to probe the stability of the oxidized surface under UHV conditions, TP-ISS studies were performed. Figure 3(b) shows the evolution of the three observed peaks with annealing temperature. Decrease in the oxygen peak is concurrent with an increase in the Sn peak and appearance and increase in the double-scattering peak consistent with desorption of bridging oxygen atoms. After reaching an intermediate plateau in a temperature range between ∼ 700 K and 800 K, the Sn peak intensities increase again until at 870 K the intensities reach another plateau. This may indicate a further reduction of the surface and possible removal of second-layer oxygen atoms. A facile reduction mechanism was also indicated for SnO2 (110) [6]. Unlike the (110) surface, however, the (101) surface does not reconstruct upon reduction. The stability of the oxygen-depleted (101) surface can be explained by the dual valency of Sn. Sn is known to exist as Sn2+ or Sn4+ in bulk compounds. The stoichiometric (101) surface (fig. 2(a)) is created by breaking the same number of O-Sn and Sn-O bonds and thus the surface maintains a Sn4+ O2 2−

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stoichiometry. In order to create a Sn-terminated surface, three Sn-O bonds are broken per surface Sn atom, thus leaving the surface Sn atoms threefold coordinated with bonds to threefold coordinated O atoms (fig. 2(c)). This results in a surface termination with a perfect Sn2+ O2− stoichiometry and thus satisfies the valency of Sn. For a reduced (110) surface, no analogous arrangement of atoms exists that can readily satisfy the Sn2+ oxidation state. This explains the stability of the (101) surface, while the (110) surface forms complex reconstructions despite the fact that the surface energy of the stoichiometric (110) is lower than that of (101). All-electron first-principles density-functional theory calculations were performed to verify the experimental observations and to calculate the stability of the proposed surface terminations across the range of oxygen chemical potential µO . The results are shown in fig. 3(c). The left-hand side of the figure corresponds to an oxygen-poor environment in which all of the oxygen in the reference system is bound within the metal oxide and the pressure above the surface is nominal. The maximum µO on the right of fig. 3(c) is equivalent to an oxygen atom bound in an O2 condensate on the surface. The stoichiometric (101) surface is independent of the oxygen pressure and thus exhibits a constant surface energy of 100.3 meV/˚ A2 across the allowed range of µO [14]. The free energy of the reduced surface, however, exhibits a strong dependence on µO . Under UHV conditions, formation of the reduced surface is strongly favored over the stoichiometric surface, with a surface free energy of 13 meV/˚ A2 at the lower limit of µO . As µO is increased, there is a threshold value of approximately −1.8 eV above which the reduced surface becomes thermodynamically less favored than the stoichiometric surface, and is predicted to reoxidize. Conversely, below this threshold the oxidized surface will be readily reduced in the absence of significant kinetic barriers. Assuming a partial oxygen pressure of 10−14 mbar under UHV conditions in our experiments, this threshold corresponds to a temperature of 650 K, which is in fortuitously good quantitative agreement with experimentally observed desorption temperature of oxygen at ∼ 680 K. It should be noted that the −1.8 eV threshold corresponds to the reaction energy between the oxidized and reduced surfaces plus O2 , and thus represents a lower bound to the kinetic barrier. Both the stoichiometric and the reduced surfaces did not exhibit large surface relaxations (see fig. 2), indicating the relatively stable nature of these surfaces and that minimal structural changes are required to convert from one to the other as the oxygen concentration is varied. This is also consistent with our experimental results that show that the (101) surface can be readily oxidized at 300 K, in stark contrast to the (110)-4 × 1 surface that only converts to a stoichiometric surface at elevated temperatures (700 K) and high oxygen pressures [6]. The threefold coordinated Sn atoms of the reduced (101) surface have similarities with the Sn atoms in tin-monoxide. For this material it is well established that “lone electron pairs” form by hybridization of s and p atomic orbitals. These electron pairs form “hats” of charge that screen the Sn2+ ions [15]. Our electronic structure calculations for the reduced (101) surface confirmed the formation of such charge hats. In addition, the surface tin atoms exhibit a net Mulliken population of +0.41e, which is nearly half the bulk value of +0.73e and consistent with the change in oxidation state. These changes in the charge distribution at the surface may cause a band bending at the surface and thus contribute to the Schottky barrier. Changes in barrier heights between grains would explain conductivity changes in polycrystalline SnO2 films when exposed to reducing or oxidizing gases. Alternatively, an easy reduction of SnO2 surfaces may also imply an easy reducibility of the bulk. The formation of an entirely oxygen-depleted surface, as has been demonstrated in this work, may facilitate the formation of bulk oxygen vacancies by providing low-energy sites at the surface for oxygen atoms to jump to from bulk sites. In such a scenario the surface would act as a reservoir of oxygen atoms to diffuse into the bulk under oxidizing conditions and as reservoir of vacancies for oxygen atoms to diffuse to from the bulk under reducing conditions. Oxygen vacancies in

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the bulk are known to intrinsically n-dope SnO2 and thus alter the conductivity of the material. In either case, the variable oxygen concentration at the surface of SnO2 is an important factor to explain the fundamental gas-sensing mechanism of this material. Although a similar oxidation and reduction behavior seems to exist for all SnO2 surfaces, the (101) surface is particularly well suited for such studies because it is stable with respect to reconstruction. ∗∗∗ The authors thank Prof. R. Helbig for providing the SnO2 samples. Financial support from NSF grant CHE-0109804 and NASA is gratefully acknowledged. REFERENCES [1] Barsan N. and Weimar U., J. Electroceramics, 7 (2001) 143; Barsan N., Schweizer¨ pel W., Fresenius J. Anal. Chem., 365 (1999) 287. Berberich M. and Go ` C., Guidi V., Stefancich M., Carotta M. C. and Martinelli G., J. Appl. Phys., [2] Malagu 91 (2002) 808; Lantto V., Rantala T. T. and Rantala T. S., J. Eur. Ceramic Soc., 21 (2001) 1961. [3] Wang X.-G., Weiss W., Shaikhutdinov Sh. K., Ritter M., Petersen M., Wagner F., ¨ gel R. and Scheffler M., Phys. Rev. Lett., 81 (1998) 1038; Wang X.-G., Chaka A. Schlo and Scheffler M., Phys. Rev. Lett., 84 (2000) 3650; Reuter K. and Scheffler M., Phys. Rev. B, 65 (2001) 035406; Dulub O., Diebold U. and Kresse G., Phys. Rev. Lett., 90 (2003) 016102. [4] Oviedo J. and Gillan M. J., Surf. Sci., 463 (2000) 93. [5] Pang C. L., Haycock S. A., Raza H., Møller P. J. and Thornton G., Phys. Rev. B, 62 (2000) R7775. [6] Batzill M., Katsiev K. and Diebold U., Surf. Sci., 529 (2003) 295. [7] Cox D. F., Fryberger T. B. and Semancik S., Phys. Rev. B, 38 (1998) 2072; Surf. Sci., 224 (1989) 121; Cavicchi R., Tarlov M. and Semancik S., J. Vac. Sci. Technol. A, 8 (1990) 2347. ´sart E., Darville J. and Gilles J. M., Surf. Sci., 126 (1983) 518. [8] de Fre [9] Perdew J. P., Burke K. and Ernzerhof M., Phys. Rev. Lett., 77 (1996) 3865. [10] Delley B., J. Chem. Phys., 92 (1990) 508; 113 (2000) 7756. Dmol3 is available from Accelrys, Inc. [11] Chase M. W., NIST-JANAF Thermochemical Tables, 4th edition, J. Phys. Chem. Ref. Data (1998). [12] Gurvich L. V., Veyts I. V. and Alcock C. B., Thermodynamic Properties of Individual Substances, 4th edition, Vol. 2 (Hemisphere Publishing Corporation, New York) 1991. [13] Oviedo J. and Gillan M. J., Surf. Sci., 513 (2002) 26. [14] This value using an all-electron methodology and numerical basis sets is 20% higher than the 83 meV/˚ A2 (1.33 J/m2 ) calculated by Gillan using plane waves and pseudopotentials. [15] Watson G. W., J. Chem. Phys., 114 (2001) 758; Terra J. and Guenzburger D., Phys. Rev. B, 44 (1991) 8584; Meyer M., Onida G., Palummo M. and Reining L., Phys. Rev. B, 64 (2001) 045119.