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Detailed spectroscopic studies of oxygen on metal carbide surfaces
Surface Science 412/413 (1998) 384–396
Detailed spectroscopic studies of oxygen on metal carbide surfaces Peter Frantz, Stephen V. Didziulis * Surface Science Department, Mechanics and Materials Technology Center, The Aerospace Corporation, El Segundo, CA 90245, USA Received 5 January 1998; accepted for publication 26 May 1998
Abstract Transition metal carbides and nitrides possess remarkable physical properties due to the unusual chemical bonding phenomena that result from the inclusion of a nonmetal atom within the metallic matrix. It has been widely recognized that this unusual situation, which so profoundly affects the physical properties, should also influence the surface chemical behavior. This paper explores the surface chemistry of TiC(100) and VC(100) with oxygen, and the potential impact of surface electronic structure on this chemistry. X-ray photoelectron spectroscopy ( XPS) and high resolution electron energy loss spectroscopy (HREELS ) were employed to observe the chemical state and specific bonds formed between resultant materials in the interfacial region. We find that oxygen adsorbed dissociatively on both surfaces, but discriminated between them by reacting with TiC to form an oxide while forming a metastable overlayer of VNO species on VC. It is proposed that the reaction on TiC was initiated by the preference of oxygen for the carbon atom, due to the predominantly C 2p character of the highest occupied energy level, resulting in evolution of CO and oxidation of the surface. This reaction was hindered on VC, where the additional electron resides in a predominantly x V 3d level. Limited oxidation proceeded on VC with high temperature anneals following room temperature exposure. © 1998 Elsevier Science B.V. All rights reserved. Keywords: Carbides; Electron energy loss spectroscopy; Oxidation; Oxygen; Photoelectron spectroscopy; Scanning tunneling microscopy; Single crystal surfaces; Solid–gas interfaces; Surface chemical reaction; Titanium carbide; Vanadium carbide; Water
1. Introduction Transition metal carbides (TMC ) and nitrides have attracted a great deal of attention in a variety of communities due to their remarkable physical properties. For example, TiC and TiN have found use as extremely hard, wear resistant coatings on mechanical device components [1–4]. In such applications, they protect the underlying steel surfaces from plastic deformation and gouging, and they are generally believed to be less reactive than steels, limiting adhesion and lubricant degradation. While these advantages have prompted successful * Corresponding author. Fax: +1 310 3361636.
implementation in many situations, very little work has been performed to understand and document the surface chemistry of TMCs and its relationship to electronic structure properties. This paper details experimental investigation into the surface chemical properties and reactivity of TiC and VC using high resolution electron energy loss spectroscopy (HREELS) and X-ray photoelectron spectroscopy ( XPS). Oxygen is used as an electron withdrawing adsorbate to probe the effects of surface electronic structure on bonding and reactivity. Previous investigation into the surface electronic structure of these materials has provided a map of the occupied density of states [5–7]. Resonant
0039-6028/98/$ – see front matter © 1998 Elsevier Science B.V. All rights reserved. PII: S0 0 39 - 6 0 28 ( 98 ) 0 04 5 6 -7
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valence band photoelectron spectra of TiC and TiN revealed that for TiC the highest occupied states are located predominantly on the carbon atom, while for TiN (and by analogy for the isoelectronic VC ) they are located on the metal atom [5]. The occupation of these orbitals will influence the preferential sites for electron accepting adsorbates, such as oxygen. Several studies have investigated the surface chemical composition and oxidation behavior of the (100) [8–15] and (111) [8–10,16 ] faces of TMCs. In one study [8], XPS and ion scattering spectroscopy (ISS ) were used to determine the composition of the (100) surfaces of TiC, VC, ZrC, and TaC after oxygen exposure. It was concluded from ISS that the oxygen interacted preferentially with the metal atoms on VC and TaC, and with the carbon atoms on TiC and ZrC, although spectroscopic proof of these chemical interactions has not been reported. These results suggest bulk oxidation may initiate with the reaction of O with the carbon atom, 2 leading to the evolution of CO. Although the reaction of oxygen and TMC surfaces has been studied, the specific reaction products and bonding mechanisms have not been explored with detailed spectroscopic work. To develop a complete understanding of the surface chemistry during interaction with oxygen, it is necessary to clarify the reaction pathways and to identify chemical intermediates leading to the equilibrated surface. The work presented here is to serve as a foundation for further controlled studies of surface electronic properties of three closely related materials, TiC, VC, and TiN. These materials share the rock salt structure, enabling a direct comparison of their structure and chemistry. The (100) surface is nonpolar – ideally containing equal numbers of metal and carbon atoms, enabling an assessment of the chemical differences between the two constituents. Through systematic variation of the metal and nonmetal components, we hope to gain insight into the surface of these materials. Here we show that oxygen may discriminate between potential binding sites because of the different electronic structures of these surfaces. This conclusion may potentially influence the choices of coatings, lubricants, and additives in future tribological applications.
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2. Experimental The TiC(100) sample was obtained as a gift from UBE Industries of Japan, and the growth techniques are described in the literature [17]. The crystal growers estimated the stoichiometry to be TiC , although our quantitative XPS results 0.9–0.92 showed the surface to have 1:1 stoichiometry (within the error of the technique). The VC(100) sample was obtained from the Linfield Research Institute, Linfield College, McMinnville, OR, and had similar stoichiometry by XPS. Both crystals were polished with diamond paste down to a grit size of 0.05 mm and solvent cleaned with ethanol and acetone. Sample cleaning in ultra-high vacuum entailed argon ion sputter/anneal cycles until the oxygen XPS signal was minimized. The accelerating potential of the Ar ion beam was 500 V, and the samples were bombarded while hot (870 K ). This treatment was followed by electron beam heating to 1670 K for 10 s. Gas exposures were performed by placing the sample in the path of a gas dosing system which consisted of a turbo-pumped gas supply, a leak valve, and a 1/4 in dosing tube terminated by a channel plate. This system provided an enhancement in the mass adsorbed of about 25 times the mass deposited from a backfilled gas. Oxygen was supplied by a cylinder of 99.995% purity (Air Products and Chemicals, Tamaqua, PA). Sample heating was performed with a resistively heated tungsten filament mounted directly behind the sample stage on the manipulator. Cooling was provided through a copper braid attached to a liquid nitrogen cooled heat sink. The temperature was monitored with a type K thermocouple on the sample stage. High resolution electron energy loss spectra (HREELS) were taken with a double-pass spectrometer (LK 2000, LK Technologies, Inc.) with a sample current of 10−10 A, resolution of 11 meV, and an elastic peak intensity of 3×105 counts/s. The spectra were obtained with incident beam energies of 7 eV, they were collected digitally and typically took 1 h each. The pass energy of the analyzer was 2 eV. The angle of incidence of the electrons was fixed at 60° with respect to the
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surface normal, and the signal was collected from the specular reflection. XPS analyses were performed with a Surface Science Instruments (SSI ) S-Probe 100 instrument with a monochromated Al Ka source. Survey scans were collected with an 800 mm spot and a 150 eV pass energy, while more detailed analyses were performed with the same spot size and 50 eV pass energy. The spectrometer and vacuum system have been described in detail elsewhere [18].
Table 1 Major core level binding energies (eV ) for clean VC and TiC Material
C 1s
M 2p
VC TiC
282.6 282.6
513.0 454.7
3/2
O 1s — 530.8
3. Results and analysis 3.1. Clean substrates The HREEL spectrum of clean VC, shown in Fig. 1, is dominated by a single loss feature centered at 510 cm−1. No other peaks were found. The feature at 510 cm−1 has previously been identified as the metal–carbon stretching vibration [19]. An XPS survey spectrum of the same sample revealed only peaks due to vanadium and carbon at the expected binding energies, as summarized in Table 1. The HREEL spectrum of TiC, shown in Fig. 2, exhibited a similar metal–carbon stretching vibration (510 cm−1), but an additional loss feature
Fig. 1. HREEL spectrum of the VC surface after repeated Ar ion sputtering and flashing to 1500 K in an electron beam heater. In addition to elastically scattered electrons, a single loss feature due to the vanadium–carbon stretching vibration is found at 510 cm−1.
Fig. 2. HREEL spectrum of the TiC surface after repeated Ar ion sputtering and flashing to 1670 K in an electron beam heater. A titanium–carbon stretch is found at 510 cm−1, with a shoulder centered at about 670 cm−1 due to a titanium–oxygen vibration.
appeared as a shoulder on the M–C stretch, centered at about 670 cm−1. An XPS survey spectrum of the clean TiC surface, summarized in Table 1, revealed a small amount of residual oxygen (530.8 eV ) on the surface, despite repeated attempts to remove it through Ar ion sputtering and higher temperature flashes. The Ti 2p region of the XPS spectrum showed no peaks in addition to the carbidic titanium peaks (454.7 eV for the 2p ) that would indicate the presence of a second 3/2 phase due to oxidation. Oxygen represented about 5% of the atomic concentration within the analyzed volume. This evidence of oxygen contamination, and previous assignment of Ti–O vibrations in Ti(OCH CH ) at 630 cm−1 [20], and the Ti–O 2 34 phonon at 750 cm−1 in TiO [21], lead us to assign 2 the 670 cm−1 peak to the Ti–O stretching vibration of the lattice substituted oxygen impurity. Detailed XPS spectra were used to determine the ratio of metal to carbon composition of both materials. To determine the relative quantities of
P. Frantz, S.V. Didziulis / Surface Science 412/413 (1998) 384–396 Table 2 Ratio of vanadium metal to nonmetal components calculated from the integrated intensities of the peaks listed in Table 1 Ratio
Clean
32 L O 2
100 L O 2
Flash to 670°K
V/C V/O
0.93 —
0.98 6.3
0.89 5.6
1.09 6.6
each atomic species, integrated peak intensities were corrected for differences in sensitivity by Scofield values and mean free path [22]. These values, shown in the first column of Tables 2 and 3, are significantly below 1:1, suggesting that the surfaces are somewhat carbon-rich. However, uncertainties in the sensitivity factors and electron scattering lengths may be present due to the carbidic character of these materials. We surmise that the actual stoichiometry of both materials is closer to 1:1. Two techniques have been used to assess the surface geometry and order of these materials. When effectively cleaned, the VC and TiC surfaces both exhibit very sharp, square, 1×1 LEED patterns, indicating that the surfaces are well-ordered and terminate in a bulk structure [23]. In previous work, ultra-high vacuum STM measurements of TiC(100) prepared in a similar fashion to this ˚ work revealed a well-ordered surface with 300 A wide terraces, separated by single and double step heights, evenly spaced across the surface. STM images obtained within these terraces revealed the expected square crystal lattice [23]. The observed ˚ ) is consistent with the obserlattice spacing (3.1 A vation of sites of electron density associated with either the Ti or C atoms. Numerous surface defects (most likely surface vacancies) were also observed at this scale. Additional work is being performed to elucidate the details of these features. Table 3 Ratio of titanium metal to nonmetal components calculated from the integrated intensities of the peaks listed in Table 1 Ratio Clean
25 L O 2
75 L O 2
150 L O 2
Flash to 670°K
Ti/C Ti/O
0.97 5.3
1.15 2.9
1.10 2.6
1.05 2.9
0.94 10
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3.2. Clean VC exposed to oxygen Plotted in Fig. 3 is a series of HREEL spectra collected following sequential exposures of oxygen to the clean VC(100) surface. Each spectrum was normalized by the intensity of the elastic peak. The oxygen exposures shown here were performed at 145 K; similar results were obtained when the experiment was performed at room temperature (see, for example, spectrum A of Fig. 12 – here, a 2 L oxygen dose at room temperature produced a HREEL spectrum that is nearly identical to the corresponding low temperature experiment shown in Fig. 3D). As the dose was increased from 0.5 to 4 L, we observed the growth of a prominent loss feature at 980 cm−1. This peak may be assigned by comparison with infrared spectra of oxovanadium ions and related compounds [24– 26 ]. In these studies, stretching frequencies of the vanadium–oxygen double bond were commonly reported to be approximately 997 cm−1. Thus, we attribute the 980 cm−1 feature to an oxygen atom at a vanadium atop site. This peak saturated at an exposure of between 4 and 10 L, and no further change was detected up to 100 L exposure [23]. Similar results were obtained when exposures were conducted at room temperature (300 K ). Fig. 3F
Fig. 3. HREEL spectra of VC(100) collected after dosing with sequentially increasing amounts of oxygen at 145 K. Shown here are the clean surface (A) and after doses of 0.5 L (B), 1.0 L (C ), 2.0 L (D) and 4.0 L ( E ), respectively. Spectrum F shows the HREEL spectrum of VC(100) collected after dosing with 100 L of oxygen at 300 K.
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shows that even after large exposures (100 L) at 300 K, there are only slight differences from the results after saturation coverage at low temperature (145 K ). At saturation, the integrated intensity was approximately equal to the intensity of the metal–carbon vibrational feature, and we estimate this to be a coverage of approximately one monolayer. HREELS measurements after exposures of greater than 100 L were not possible due to our inability to recover the signal. Other possible assignments of the 980 cm−1 loss feature are the O–O stretch of molecular oxygen, with the bond softened considerably from the gas phase frequency of 1600 cm−1 by electron donation from the surface [27,28], and C–O single bond stretch [29]. The first mechanism was proposed for the adsorption of oxygen on an analogous carbide, NbC [28]. We consider these possibilities to be unlikely in our case, as a strong signal is evident at room temperature and above, where it is very unlikely that any molecular oxygen would remain adsorbed under UHV conditions. The C–O species conflict with our XPS data, presented below, and with ion scattering data in the literature [8]. After exposures of greater than 1 L, another loss feature was detected in the region of 700 cm−1. No further increase of this peak has been observed beyond 4 L up to the largest exposure of 100 L. By analogy with the assignment of the oxygen contaminant on the TiC surface, we attribute this feature to the stretching of the vanadium–oxygen single bond. We conjecture that this is oxygen bound to vanadium ( V–O–V ) at carbon vacancies, either naturally present or formed as a result of limited CO evolution. x Other spectral features found after oxygen exposure are located at 1400, 2900, and 3650 cm−1. These are identified as the bending (1400 cm−1) and stretching modes (2900 cm−1) of carbon– hydrogen species, and the stretching of the oxygen– hydrogen bond (3650 cm−1). They are likely due to adsorption of background contaminants (hydrocarbons, water, and hydrogen) from the environment during these experiments, which typically elapse 10 h of time, although time and exposure have not been monitored as independent variables. The presence of oxygen may enhance
Fig. 4. V 2p X-ray photoelectron spectra of the VC surface before (solid line) and after exposure to 100 L of oxygen (dashed line) at room temperature.
the adsorption of water and hydrogen, and the oxygen gas supply may contribute as a source of contaminants. It is significant to note that no peaks were observed in the vicinity of 1800 cm−1, where one would expect to find a carbonyl vibrational mode. XPS data collected after room temperature oxygen exposures showed broadening of the V 2p core level peaks, while the C 1s remained unperturbed. Fig. 4 shows the vanadium 2p peaks from the clean sample (solid line) and after a dose of 100 L (dashed line). We find that the peaks have broadened measurably to deeper binding energy, increasing from 2.0 to 2.3 eV. We interpret this change as indicating that the oxygen interacts with the vanadium by withdrawing electron density. This broadening was evident at a lower exposure of 32 L (the smallest used for XPS study). Difference spectra collected during sequential exposures (Fig. 5) show more clearly the shift to deeper binding energy. These spectra were obtained by subtracting the normalized spectrum of clean VC from the spectrum corresponding to each oxygen dose, thereby removing the contribution from dominant bulk VC. After dosing 32 L at room temperature (Fig. 5A), we find the growth of an asymmetric 2p peak located significantly 3/2 deeper in binding energy than the V 2p peaks of VC, which are themselves shifted a small amount
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Fig. 5. Difference spectra of the V 2p peaks during sequentially increasing doses of oxygen at room temperature. In each spectrum, the clean sample data have been subtracted from the data obtained after each dose. Spectra A and B were collected after the 32 L and 100 L doses, respectively.
from the reported metallic vanadium binding energy (512.3 eV for the 2p ) [30]. The binding 3/2 energy of the most intense portion of this feature is lower than 515.8 eV expected for VO [31] and 2 517.0 eV for V O [30]. The difference spectrum 2 5 after 100 L exposure (Fig. 5B) shows no additional increase in the deeper binding energy regions of the V 2p difference spectrum indicative of further oxidation. Fig. 6 shows the C 1s region of the VC photoelectron spectrum before (solid line) and after (dashed line) exposure to 100 L of oxygen. The intensity scale has been expanded to magnify potential changes in spectral features, and the full scale figure is shown in the inset. This figure shows no change in the C 1s spectrum due to oxygen exposure, demonstrating that the oxygen has no detectable interaction with the carbon atom of the VC(100) surface. There was no evidence of perturbation of the C 1s peak throughout sequential doses from 32 to 200 L. This is consistent with the HREELS results which showed no peaks in the regions where one would expect carbon–oxygen vibrational features. The immutability of the C 1s spectrum during oxygen adsorption, coupled with changes observed in the V 2p spectra, supports the previous assignment of the 980 cm−1 feature of
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Fig. 6. Expanded C 1s photoelectron spectra of the VC surface before (solid line) and after (dashed line) exposure to 100 L of oxygen at room temperature. No changes in this spectral feature are evident after oxygen exposure. The inset shows this figure scaled to display the entire data range.
Fig. 7. O 1s photoelectron spectra of the VC surface during sequentially increasing doses of oxygen. Spectrum A was collected from the clean surface, and spectra B and C were collected after exposure to 32 and 100 L of oxygen, respectively. Spectrum D was collected after flashing the sample to 670 K after 100 L exposure.
Fig. 3 to the oxygen vibration at a vanadium atop site. The O 1s peak of oxygen exposed VC, shown in Fig. 7, is located at about 530.0 eV, with a small tail that extends to deeper binding energy. In this figure, the spectrum of the clean sample (Fig. 7A) is shown with the sample after doses of 32 L
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(Fig. 7B) and 100 L (Fig. 7C ). We find that the peak maintained a constant binding energy, and the ratios of vanadium to oxygen determined by integrated peak intensities, listed in Table 2, show that the amount of oxygen changed little as the dose was increased beyond 32 L. This observation supports the previous conclusion that no significant chemical changes occurred beyond the HREELS saturation exposure of 4 to 10 L. Spectrum 7D will be discussed below. From XPS spectral intensities, the ratios of metal to carbon composition were determined to provide a measure of the relative quantity of carbon present at the surface after exposure to oxygen and subsequent heating. Shown in Table 2, these quantities were determined as described above. We have assumed that the absolute quantity of vanadium remained constant throughout the experiment. The V:C ratios, shown in Table 2, showed no systematic changes within our experimental uncertainty, indicating that no carbon loss was detectable during exposure to oxygen at 300 K. The final component of this table, the metal/carbon ratio after heating to 670 K, will be discussed below.
Fig. 8. HREEL spectra of TiC(100) collected after dosing with sequentially increasing amounts of oxygen at 300 K. Shown here are the clean surface (A) and after doses of 2.5 L (B), 20 L (C ), 40 L (D) and 80 L ( E ), respectively.
3.3. Clean TiC exposed to oxygen Plotted in Fig. 8 is a series of HREEL spectra collected after sequential exposures of oxygen to the TiC(100) surface at 300 K temperature. As the dose was increased, we observed the initial appearance of a small loss feature at 980 cm−1, which appeared to saturate between exposures of 10 and 20 L. In analogy to the oxygen exposed VC surface, and with the previous report of a 1008 cm−1 stretching frequency for the diatomic TiNO [24], we assign this feature to the TiNO stretching vibration of oxygen at a titanium atop site. The most prominent loss feature following greater oxygen exposure of the TiC surface was located at about 730 cm−1. This broad peak was observed only after relatively large doses of oxygen (>40 L), although its presence may be masked by existing Ti–O features at low coverage. This band is similar to the Ti–O stretching frequency in TiO [21], an 2 assignment which suggests a significant degree of
Fig. 9. Difference spectra of the Ti 2p peaks during sequentially increasing doses of oxygen at room temperature. Spectra A through C were collected after the 25 L, 75 L, and 150 L doses, respectively.
oxidation. Low temperature (150 K ) oxygen exposures produced similar results. XPS data collected after oxygen exposure at 300 K showed that Ti 2p features appeared at higher binding energy than the carbidic Ti 2p peaks, concurrent with growth on the high binding energy side of the C 1s peak. Fig. 9 shows the difference spectra of the Ti 2p peaks after sequentially increasing oxygen doses. After exposure of
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Fig. 10. Expanded C 1s photoelectron spectra of the TiC surface before (solid line) and after (dashed line) exposure to 300 L of oxygen at room temperature. The inset shows this figure scaled to display the entire data range.
25 L ( Fig. 11A) we find growth of features at deeper binding energy than the 2p peaks of TiC near 456 eV on the 2p . The 2p shift of about 3/2 3/2 1.3 eV is similar in magnitude to the shift of the V 2p peaks after the oxygen exposures discussed above. However, in contrast to VC, the Ti core levels appeared to grow features even deeper in binding energy as the exposure was increased. After the 75 and 150 L doses, the 2p difference 3/2 peak center was significantly higher than the binding energy of 454.8 eV expected for TiO [32], yet less than the value of 458.3 eV expected for TiO 2 [33]. The breadth of this Ti 2p feature, stretching 3/2 from 455 to 459 eV, suggests that after large exposures the surface is composed of a variety of TiO species. x Fig. 10 shows the C 1s peak before (solid line) and after (dashed line) exposure to 150 L oxygen at 300 K. We find the growth of a small shoulder on the carbon peak, centered at 283.2 eV, representing a shift of 1.7 eV to deeper binding energy due to a carbon–oxygen interaction. This feature was evident after exposures as low as 25 L. The relatively low binding energy of this feature allows us to rule out any contribution due to carbonaceous contaminants (284.6 eV ). The appearance of this feature indicates that, in addition to the titanium–oxygen interaction demonstrated above,
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Fig. 11. O 1s photoelectron spectra of the TiC surface after sequentially increasing doses of oxygen. Spectrum A, collected from the clean surface, shows a small amount of oxygen (5% of the surface composition) due to the bulk impurity. Spectra B through D were collected after exposure to 25 L, 75 L, and 150 L of oxygen, respectively.
the oxygen interacts with carbon on the TiC surface. The O 1s peak due to the bulk oxygen impurity on clean TiC, shown in Fig. 11A, appeared at 530.5 eV. The intensity of this peak grew and shifted to lower binding energy with oxygen doses of 25 L ( Fig. 11B), 75 L ( Fig. 11C ), and 150 L ( Fig. 11D). After a dose of 150 L, the O 1s peak was located at 530.2 eV, equal to the value expected for TiO [32]. The Ti/C ratio ( Table 3) 2 shows an abrupt increase after the 75 L dose, correlating to a decrease in the Ti/O ratio, and the growth of the 730 cm−1 Ti–O stretch ( Fig. 8) after similar doses. The Ti/C ratio apparently changed little, while the Ti/O ratio decreased only a small amount, after doubling the exposure to 150 L. Metal to carbon ratios, determined as described above and listed in Table 3, indicate that carbon was lost during exposure to oxygen. The data shown here are averaged over two separate experiments. The metal to carbon ratio increased an insignificant amount with the 25 L oxygen dose, but jumped dramatically with the 75 L dose. No further increase was found with larger doses up to 300 L. This is believed to reflect the loss of carbon as CO species from the surface. The most signifix cant loss of carbon occurred after exposures similar
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to those that produced the 730 cm−1 HREELS vibration (Fig. 8E) due to Ti–O, and the pronounced shift in binding energy of the Ti 2p peaks (Figs. 9B and C ). This marks the onset of the reaction to oxidize the TiC surface. 3.4. Effects of increased temperature HREEL spectra of VC after exposure to 2 L of oxygen ( Fig. 12) show that the intensity of the VNO stretch fell to zero as the temperature was increased from 300 to 670 K. In these data, the temperature was held at 470 K ( Fig. 12B) and 670 K ( Fig. 12C ) for 30 s before returning to 300 K for analysis. After the 670 K flash, the 510 cm−1 feature remained slightly more intense, and broader to the high frequency side, than observed in the clean VC sample ( Fig. 12A). No further changes in the spectrum were observed up to the highest temperature flash of 1300 K. Despite the dramatic decrease in the intensity of loss features associated with oxygen adsorption after heating, the XPS spectrum provided evidence of significant oxygen retention within the interfacial region. Fig. 7D shows the O 1s peak on the 100 L dosed VC surface after the sample was heated to 670 K. The center of this peak shifted 0.7 eV to
deeper binding energy and retained approximately 70% of its intensity. Concurrent with these spectral changes, the VC surface lost some of its carbon atoms upon heating. The final column of Table 2 shows that the ratio of metal to carbon composition increased by 0.15 beyond the average ratio before heating. This demonstrates that the surface region lost carbon atoms, likely in the form of evolved CO species. x HREEL spectra of oxygen covered TiC after heating also show a decrease in the intensity of oxygen features. The spectrum collected after flashing to 670 K, not shown, was marked by the absence of the TiNO feature at 980 cm−1 and the reduction of the 730 cm−1 to a broad trail on the high frequency side of the 510 cm−1 Ti–C feature. The O 1s peak showed only a 6% decrease in the integrated intensity, but no measurable change in position as a result of similar treatment. This large change in HREEL spectra, concurrent with small changes in XPS data, indicates that although a small amount of oxygen may have evolved as CO species, most of it had been driven subsurface x in TiC. In contrast to the VC surface, metal/carbon ratios for TiC (final column of Table 3) did not increase during heating within the experimental uncertainty of the technique.
4. Discussion
Fig. 12. HREEL spectra of oxygen exposed VC after increased temperature flash cycles. Spectrum A shows the surface after a 2 L exposure at 300 K. Spectra B and C were collected after raising the temperature to 70 and 670 K before returning to room temperature for data collection.
The combination of HREELS and XPS studies on the adsorption and reaction of relatively low oxygen exposures on the (100) surfaces of VC and TiC has verified a significant difference between the materials and identified surface reaction products. In this discussion, we will first summarize the experimental results, followed by potential explanations and implications of our findings. The adsorption of oxygen on VC(100) essentially saturates at relatively low exposures (~10 L) under the conditions studied. This saturation was indicated by the constant intensity and binding energy of the O 1s XPS peak and by an unchanging HREEL spectrum after exposures ranging from 10 to 100 L. The HREELS data are dominated by the 980 cm−1 peak that we have assigned as the VNO stretch of adsorbed, dissociated oxygen atop
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a vanadium atom. The XPS substrate core level results were consistent with these data, displaying a broadening on the deeper binding energy side of the V 2p peaks and no observable effect on the C 1s peak. The V/C atomic concentration ratios derived from the XPS results show no systematic changes with increasing exposure, implying that the oxygen overlayer is stable and that the substrate is not being oxidized via the loss of carbon oxides under these conditions. An additional small feature was observed in the HREELS near 700 cm−1, which is consistent with a vanadium–oxygen single bond. The low intensity of this feature shows that this reaction product was minor, perhaps arising from surface carbon vacancies or structural defects that favor the formation of a bridging oxygen species V–O–V. This surface species would also be typical of that expected from surface oxidation, indicating that a small number of surface sites may have reacted by losing carbon, as detailed for TiC below. Upon heating to temperatures in the range of 470 to 670 K, the O/VC HREELS data change considerably while most of the oxygen was retained by the substrate as indicated by XPS. Specifically, the VNO and V–O–V vibrations disappeared, while features may have grown under the V–C stretch in the spectrum. The O 1s peak lost some intensity and shifted to higher binding energy, indicating the formation of a different chemical species. The V/C concentration ratio increased, consistent with some loss of surface carbon as expected for the onset of oxidation. The drastic HREELS changes coupled to the XPS data may indicate that much of the remaining oxygen has been driven below the surface by heating. The results on TiC(100) were significantly different, as the XPS oxygen signal did not saturate in the exposure regime studied. The initial exposures (<40 L) populated a site with a small HREELS signal near 980 cm−1, similar to the major feature observed on VC and thus assigned as TiNO. With increasing exposures, however, a lower energy HREELS feature grew to dominate the spectrum. This feature was in the region expected for Ti–O–Ti surface species, such as would be expected in TiO . The XPS core levels 2 showed changes consistent with these two different
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surface products. At lower exposures, the Ti 2p levels were broadened to deeper binding energy and the Ti/C concentration ratio shows little change. With increasing exposure, the Ti/C concentration ratio increased and deeper binding energy features grew on the Ti 2p peaks, demonstrating the loss of carbon and the oxidation of the surface at 300 K. Significant changes were also observed in the C 1s XPS spectrum of O/TiC, showing that oxygen was interacting with surface carbon. This interaction was evidenced by a deeper binding energy shoulder on the C 1s peak. No definitive HREELS peaks characteristic of carbon–oxygen bond frequencies were found as would be expected from the measurable perturbation of the C 1s XPS peak. It is conceivable that oxygen could occupy an adsorption site involving both titanium and carbon atoms, and the stretching frequency of this species could be present in the broad feature observed between 500 and 1000 cm−1 (see Fig. 8E). The measured Ti/O XPS concentration ratio was roughly half that found for VC after the larger exposures, consistent with more extensive surface reaction on TiC. Finally, heating to 670 K did alter the HREEL spectrum, but had little effect on the relative concentration of the surface species, showing that further oxidation was not being thermally driven. For either material, the spontaneous formation of the metal oxide and a carbon oxide is thermodynamically favored by a considerable amount. In fact, if left in laboratory ambient, TiC and VC will form oxide layers. The lack of oxidation of VC(100) under our experimental conditions indicates the presence of a kinetic barrier that is not present for TiC. We surmise that the surface reaction products formed can either enhance or inhibit further reaction. These reaction products are a result of the TiC and VC surface geometric and electronic structures, and the differences between the materials in these regards may explain the experimental results. The difference between pristine, defect-free surfaces of TiC and VC is the identity of the metal atom. Chemically, this difference results in the presence of an additional valence electron in a formula unit of VC. Considering the electronic
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structure from a localized chemical bonding perspective, the highest occupied valence orbitals on TiC are predominantly C 2p in nature, basically representing a filled valence band. The additional electron present in VC must occupy a molecular orbital that is predominantly V 3d in nature, and this has been verified by variable photon energy photoelectron spectroscopy on these materials [5]. Oxygen seeks electrons on a surface, hence the initial interaction of dissociated oxygen with the two materials will occur at sites with the most accessible electron density. Simply, this will be the C atoms on TiC and the V atoms on VC. The thermodynamically stable and physically volatile nature of gas phase carbon oxides facilitates their loss on TiC, and the simultaneous formation of a titanium oxide. However, the VNO species formed on VC is stable enough under the conditions of our experiment to preclude further reaction with this material. Elevated temperature can drive the oxidation of VC, as can extended time under an oxidizing environment such as the laboratory ambient. Theoretical analysis of adsorption on transition metal carbides, from the perspective of the surface electronic structure, supports this picture of oxygen adsorption [6 ]. It is also clear that these surfaces have defects, usually believed to arise from carbon vacancies resulting from the tendency of these materials to be substoichiometric. Although our quantitative XPS results of the clean surfaces do not support this notion, we believe that STM images offer compelling proof of the existence of vacancies that we will assume are missing carbon atoms [23]. Therefore, some of the surface chemistry must account for these sites. We propose that the majority species observed in HREELS arise from oxygen adsorption and chemistry on well-defined terrace sites, while minority species result from the interaction of oxygen with surface defects. We believe this because STM of TiC has shown that most of the surface is well ordered, but a significant number of defect sites are evident [23]. The minority species on TiC data produced the 980 cm−1 feature that was evident at low exposures. The similarity between this feature and that identified as the major species on VC implies a similar adsorption site, i.e. a Ti atom with addi-
tional electron density. This can most likely happen in one of two ways. First, carbon vacancies will cause the surface to appear more metallic, shifting electron density that would have resided on the carbon back to the neighboring titanium atoms. Thus, the titanium sites adjacent to vacancies may form this surface product. A second possibility is defects due to the presence of oxygen contaminants substituting for carbon in the TiC crystal structure. The effect of this substitution is to add additional valence electrons to the surface (similar to the changes between TiC and TiN ), and perhaps creating added electron density to neighboring Ti atoms. Because we were unable to remove all of the surface oxygen, we cannot distinguish between these possibilities. However, ion scattering studies indicate that the initial adsorption of oxygen on TiC may fill defect sites [8,34]. The presence of C vacancies on VC will also create additional electron density on neighboring V atoms and create a site where oxygen can bridge between two V atoms without interference from a carbon atom. The HREELS data indicate that two different sites are occupied at low exposure on VC, with no apparent preference. The minor species has been assigned as a V–O single bond that we believe is more likely to occur where carbon vacancies exist. It is possible that some of the oxygen atoms forming VNO species could also be bonding at defect sites, with a portion forming the bridge bonded species. It is conceivable that oxygen could bridge between two vanadium atoms in a pristine lattice site. Until the concentration of defects on our VC surfaces has been determined absolutely, we will be unable to draw firm conclusions. Our quantitative XPS (metal/carbon ratios) suggests that the number of defects on the two materials should be similar. In general, our results are consistent with previous ion scattering studies reported in the literature on these surfaces [8]. In that work, the bonding of oxygen with carbon on TiC was inferred from the loss of scattering intensity from carbon. However, the relatively facile oxidation of the TiC surface makes it unclear whether any blocking of C sites by adsorbed oxygen is observable, or whether carbon monoxide is formed so quickly that capturing significant numbers of reaction
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intermediates such as surface carbonyls is not possible on the time scale of these experiments. We have shown that oxygen does perturb the C 1s XPS peak, but could find no HREELS features that unambiguously identify C–O species. It is likely that the ISS results showed a combination of site blocking and carbon loss due to surface oxidation. Finally, the use of electronic structure arguments to predict surface chemistry can potentially enable predictions of the behavior of similar materials and other adsorbates. The results of this paper indicate that electron withdrawing adsorbates will initially bond with surface sites having the most accessible electron density, V on VC and C on TiC. As a result, the isoelectronic TiN and VC might well be predicted to behave similarly to oxygen exposure under these conditions. The formation of nitrogen oxides is not as favorable thermodynamically, but the driving force may well be the very stable metal oxide. Future experiments are planned on single crystal TiN and with different adsorbates to determine if these trends can be generalized.
5. Summary and conclusions We have demonstrated methods for obtaining reproducible, clean TiC and VC surfaces, and the properties of these surfaces have been investigated with spectroscopic probes. It has been shown that oxygen behaves differently on these surfaces; oxidizing the TiC surface, after binding to the carbon atom and releasing CO , while preferring to bond x with the metal atom on VC, forming a metastable overlayer. Explanation of these preferences is likely the different electronic structures of the materials. The additional electron in the VC formula unit resides in a predominantly V 3d orbital, while the highest occupied orbital in TiC is predominantly C 2p. These experiments have shown that TMC surfaces can form strong bonds to oxygen containing species, indicating that lubricant additives containing oxygen will likely interact with these materials. We are presently studying these interactions, where more complex and technologically relevant materi-
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als are exposed to these TMC surfaces under similar conditions.
Acknowledgements This work is supported by The Aerospace Corporation IR&D program funded by the Space and Missile Systems Center (SMC ) of the USAF under contract number F04701-93-C-0094.
References [1] H.S. Cheng, T.P. Chang, W.D. Sproul, in: D. Dowson, C.M. Taylor, M. Godet (Eds.), Mechanics of Coatings, Leeds–Lyon 16, Tribology Series 17, Elsevier Science, Amsterdam, 1990. [2] H.J. Boving, H.E. Hintermann, Tribol. Int. 23 (1990) 129. [3] K. Kawata, Surf. Coat. Technol. 54 (1992) 604. [4] L. Toth, Transition Metal Carbides and Nitrides, Academic Press, New York, 1991. [5] S.V. Didziulis, J.R. Lince, T.B. Stewart, E. Eklund, Inorg. Chem. 3 (1994) 1979. [6 ] S.A. Janson, R. Hoffman, Surf. Sci. 197 (1988) 474. [7] L.I. Johansson, Surf. Sci. Rep. 21 (1995) 177. [8] R. Souda, T. Aizawa, S. Otani, Y. Ishizawa, Surf. Sci. 256 (1991) 19. [9] H.L. Hakansson, L.I. Johansson, M. Hammar, M. Gothelid, Phys. Rev. B 47 (1993) 10769. [10] K. Edamoto, T. Anzawa, E. Miyazaki, S. Otani, Surf. Sci. 287 (1993) 667. [11] L.I. Johansson, A. Callenas, Solid State Commun. 42 (1982) 299. [12] L.I. Johansson, H.I. Johansson, K.L. Hakansson, Phys. Rev. B 48 (1993) 14520. [13] C. Oshima, M. Aono, T. Tanaka, S. Kawai, S. Zaima, Y. Shibata, Surf. Sci. 102 (1981) 312. [14] P.A.P. Lindberg, L.I. Johansson, Surf. Sci. 192 (1987) 366. [15] K. Miura, R. Souda, T. Aizawa, C. Oshima, S. Otani, Y. Ishizawa, J. Vac. Sci. Technol. A 7 (1989) 3013. [16 ] S. Zaima, Y. Shibata, H. Adachi, C. Oshima, S. Otani, M. Aono, Y. Ishizawa, Surf. Sci. 157 (1985) 380. [17] S. Otani, S. Honma, T. Tanaka, Y. Ishizawa, J. Cryst. Growth 61 (1983) 1. [18] S.V. Didziulis, P.D. Fleischauer, Langmuir 6 (1990) 621. [19] J.G. Chen, B.D. DeVries, B. Fruhberger, C.M. Kim, Z.-M. Liu, J. Vac. Sci. Technol. 13 (1995) 1600. [20] C.J. Puchert (Ed.), The Aldrich Library of FTIR Spectra, vol. 3, Aldrich Chemical Co., Milwaukee, first edn., 1989. [21] M.A. Henderson, Surf. Sci. 355 (1996) 151. [22] M-Probe Software Manual Version 1.30, Surface Science Instruments, Mountain View, 1990.
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[23] P.P. Frantz, S.V. Didziulis, P.B. Merrill, S.S. Perry, Tribol. Lett. 4 (1998) 141. [24] G. Herzberg, Molecular Spectra and Molecular Structure I. Spectra of Diatomic Molecules, Van Nostrand Reinhold, New York, 1950. [25] C.E. Mannix, A.P. Zipp, J. Inorg. Chem. 41 (1979) 59. [26 ] R.G. Garvey, R.O. Ragsdale, Inorg. Chim. Acta 2 (1968) 191. [27] C.M. Kim, B.D. DeVries, B. Fruhberger, J.G. Chen, Surf. Sci. 327 (1995) 81. [28] T. Aizawa, W. Hayami, R. Souda, S. Otani, T. Tanaka, Y. Ishizawa, Surf. Sci. 357 (1996) 645.
[29] F.M. Hoffman, R.A. de Paola, Phys. Rev. Lett. 52 (1984) 1697. [30] R. Larsson, B. Folkesson, G. Schon, Chem. Scr. 3 (1973) 88. [31] C.J. Groenenboom, G. Sawatzky, H.J. de L. Meijer, F. Jellinek, J. Organomet. Chem. 76 (1974) C4. [32] H.F. Franzen, M.X. Umana, J.R. McCreary, R.J. Thorn, J. Solid State Chem. 18 (1976) 363. [33] M. Murata, K. Wakino, S. Ikeda, J. Electron Spectrosc. 6 (1975) 459. [34] M. Aono, Y. Hou, R. Souda, C. Oshima, S. Otani, Y. Ishizawa, Phys. Rev. Lett. 50 (1983) 1293.
Detailed spectroscopic studies of oxygen on metal carbide surfaces Peter Frantz, Stephen V. Didziulis * Surface Science Department, Mechanics and Materials Technology Center, The Aerospace Corporation, El Segundo, CA 90245, USA Received 5 January 1998; accepted for publication 26 May 1998
Abstract Transition metal carbides and nitrides possess remarkable physical properties due to the unusual chemical bonding phenomena that result from the inclusion of a nonmetal atom within the metallic matrix. It has been widely recognized that this unusual situation, which so profoundly affects the physical properties, should also influence the surface chemical behavior. This paper explores the surface chemistry of TiC(100) and VC(100) with oxygen, and the potential impact of surface electronic structure on this chemistry. X-ray photoelectron spectroscopy ( XPS) and high resolution electron energy loss spectroscopy (HREELS ) were employed to observe the chemical state and specific bonds formed between resultant materials in the interfacial region. We find that oxygen adsorbed dissociatively on both surfaces, but discriminated between them by reacting with TiC to form an oxide while forming a metastable overlayer of VNO species on VC. It is proposed that the reaction on TiC was initiated by the preference of oxygen for the carbon atom, due to the predominantly C 2p character of the highest occupied energy level, resulting in evolution of CO and oxidation of the surface. This reaction was hindered on VC, where the additional electron resides in a predominantly x V 3d level. Limited oxidation proceeded on VC with high temperature anneals following room temperature exposure. © 1998 Elsevier Science B.V. All rights reserved. Keywords: Carbides; Electron energy loss spectroscopy; Oxidation; Oxygen; Photoelectron spectroscopy; Scanning tunneling microscopy; Single crystal surfaces; Solid–gas interfaces; Surface chemical reaction; Titanium carbide; Vanadium carbide; Water
1. Introduction Transition metal carbides (TMC ) and nitrides have attracted a great deal of attention in a variety of communities due to their remarkable physical properties. For example, TiC and TiN have found use as extremely hard, wear resistant coatings on mechanical device components [1–4]. In such applications, they protect the underlying steel surfaces from plastic deformation and gouging, and they are generally believed to be less reactive than steels, limiting adhesion and lubricant degradation. While these advantages have prompted successful * Corresponding author. Fax: +1 310 3361636.
implementation in many situations, very little work has been performed to understand and document the surface chemistry of TMCs and its relationship to electronic structure properties. This paper details experimental investigation into the surface chemical properties and reactivity of TiC and VC using high resolution electron energy loss spectroscopy (HREELS) and X-ray photoelectron spectroscopy ( XPS). Oxygen is used as an electron withdrawing adsorbate to probe the effects of surface electronic structure on bonding and reactivity. Previous investigation into the surface electronic structure of these materials has provided a map of the occupied density of states [5–7]. Resonant
0039-6028/98/$ – see front matter © 1998 Elsevier Science B.V. All rights reserved. PII: S0 0 39 - 6 0 28 ( 98 ) 0 04 5 6 -7
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valence band photoelectron spectra of TiC and TiN revealed that for TiC the highest occupied states are located predominantly on the carbon atom, while for TiN (and by analogy for the isoelectronic VC ) they are located on the metal atom [5]. The occupation of these orbitals will influence the preferential sites for electron accepting adsorbates, such as oxygen. Several studies have investigated the surface chemical composition and oxidation behavior of the (100) [8–15] and (111) [8–10,16 ] faces of TMCs. In one study [8], XPS and ion scattering spectroscopy (ISS ) were used to determine the composition of the (100) surfaces of TiC, VC, ZrC, and TaC after oxygen exposure. It was concluded from ISS that the oxygen interacted preferentially with the metal atoms on VC and TaC, and with the carbon atoms on TiC and ZrC, although spectroscopic proof of these chemical interactions has not been reported. These results suggest bulk oxidation may initiate with the reaction of O with the carbon atom, 2 leading to the evolution of CO. Although the reaction of oxygen and TMC surfaces has been studied, the specific reaction products and bonding mechanisms have not been explored with detailed spectroscopic work. To develop a complete understanding of the surface chemistry during interaction with oxygen, it is necessary to clarify the reaction pathways and to identify chemical intermediates leading to the equilibrated surface. The work presented here is to serve as a foundation for further controlled studies of surface electronic properties of three closely related materials, TiC, VC, and TiN. These materials share the rock salt structure, enabling a direct comparison of their structure and chemistry. The (100) surface is nonpolar – ideally containing equal numbers of metal and carbon atoms, enabling an assessment of the chemical differences between the two constituents. Through systematic variation of the metal and nonmetal components, we hope to gain insight into the surface of these materials. Here we show that oxygen may discriminate between potential binding sites because of the different electronic structures of these surfaces. This conclusion may potentially influence the choices of coatings, lubricants, and additives in future tribological applications.
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2. Experimental The TiC(100) sample was obtained as a gift from UBE Industries of Japan, and the growth techniques are described in the literature [17]. The crystal growers estimated the stoichiometry to be TiC , although our quantitative XPS results 0.9–0.92 showed the surface to have 1:1 stoichiometry (within the error of the technique). The VC(100) sample was obtained from the Linfield Research Institute, Linfield College, McMinnville, OR, and had similar stoichiometry by XPS. Both crystals were polished with diamond paste down to a grit size of 0.05 mm and solvent cleaned with ethanol and acetone. Sample cleaning in ultra-high vacuum entailed argon ion sputter/anneal cycles until the oxygen XPS signal was minimized. The accelerating potential of the Ar ion beam was 500 V, and the samples were bombarded while hot (870 K ). This treatment was followed by electron beam heating to 1670 K for 10 s. Gas exposures were performed by placing the sample in the path of a gas dosing system which consisted of a turbo-pumped gas supply, a leak valve, and a 1/4 in dosing tube terminated by a channel plate. This system provided an enhancement in the mass adsorbed of about 25 times the mass deposited from a backfilled gas. Oxygen was supplied by a cylinder of 99.995% purity (Air Products and Chemicals, Tamaqua, PA). Sample heating was performed with a resistively heated tungsten filament mounted directly behind the sample stage on the manipulator. Cooling was provided through a copper braid attached to a liquid nitrogen cooled heat sink. The temperature was monitored with a type K thermocouple on the sample stage. High resolution electron energy loss spectra (HREELS) were taken with a double-pass spectrometer (LK 2000, LK Technologies, Inc.) with a sample current of 10−10 A, resolution of 11 meV, and an elastic peak intensity of 3×105 counts/s. The spectra were obtained with incident beam energies of 7 eV, they were collected digitally and typically took 1 h each. The pass energy of the analyzer was 2 eV. The angle of incidence of the electrons was fixed at 60° with respect to the
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surface normal, and the signal was collected from the specular reflection. XPS analyses were performed with a Surface Science Instruments (SSI ) S-Probe 100 instrument with a monochromated Al Ka source. Survey scans were collected with an 800 mm spot and a 150 eV pass energy, while more detailed analyses were performed with the same spot size and 50 eV pass energy. The spectrometer and vacuum system have been described in detail elsewhere [18].
Table 1 Major core level binding energies (eV ) for clean VC and TiC Material
C 1s
M 2p
VC TiC
282.6 282.6
513.0 454.7
3/2
O 1s — 530.8
3. Results and analysis 3.1. Clean substrates The HREEL spectrum of clean VC, shown in Fig. 1, is dominated by a single loss feature centered at 510 cm−1. No other peaks were found. The feature at 510 cm−1 has previously been identified as the metal–carbon stretching vibration [19]. An XPS survey spectrum of the same sample revealed only peaks due to vanadium and carbon at the expected binding energies, as summarized in Table 1. The HREEL spectrum of TiC, shown in Fig. 2, exhibited a similar metal–carbon stretching vibration (510 cm−1), but an additional loss feature
Fig. 1. HREEL spectrum of the VC surface after repeated Ar ion sputtering and flashing to 1500 K in an electron beam heater. In addition to elastically scattered electrons, a single loss feature due to the vanadium–carbon stretching vibration is found at 510 cm−1.
Fig. 2. HREEL spectrum of the TiC surface after repeated Ar ion sputtering and flashing to 1670 K in an electron beam heater. A titanium–carbon stretch is found at 510 cm−1, with a shoulder centered at about 670 cm−1 due to a titanium–oxygen vibration.
appeared as a shoulder on the M–C stretch, centered at about 670 cm−1. An XPS survey spectrum of the clean TiC surface, summarized in Table 1, revealed a small amount of residual oxygen (530.8 eV ) on the surface, despite repeated attempts to remove it through Ar ion sputtering and higher temperature flashes. The Ti 2p region of the XPS spectrum showed no peaks in addition to the carbidic titanium peaks (454.7 eV for the 2p ) that would indicate the presence of a second 3/2 phase due to oxidation. Oxygen represented about 5% of the atomic concentration within the analyzed volume. This evidence of oxygen contamination, and previous assignment of Ti–O vibrations in Ti(OCH CH ) at 630 cm−1 [20], and the Ti–O 2 34 phonon at 750 cm−1 in TiO [21], lead us to assign 2 the 670 cm−1 peak to the Ti–O stretching vibration of the lattice substituted oxygen impurity. Detailed XPS spectra were used to determine the ratio of metal to carbon composition of both materials. To determine the relative quantities of
P. Frantz, S.V. Didziulis / Surface Science 412/413 (1998) 384–396 Table 2 Ratio of vanadium metal to nonmetal components calculated from the integrated intensities of the peaks listed in Table 1 Ratio
Clean
32 L O 2
100 L O 2
Flash to 670°K
V/C V/O
0.93 —
0.98 6.3
0.89 5.6
1.09 6.6
each atomic species, integrated peak intensities were corrected for differences in sensitivity by Scofield values and mean free path [22]. These values, shown in the first column of Tables 2 and 3, are significantly below 1:1, suggesting that the surfaces are somewhat carbon-rich. However, uncertainties in the sensitivity factors and electron scattering lengths may be present due to the carbidic character of these materials. We surmise that the actual stoichiometry of both materials is closer to 1:1. Two techniques have been used to assess the surface geometry and order of these materials. When effectively cleaned, the VC and TiC surfaces both exhibit very sharp, square, 1×1 LEED patterns, indicating that the surfaces are well-ordered and terminate in a bulk structure [23]. In previous work, ultra-high vacuum STM measurements of TiC(100) prepared in a similar fashion to this ˚ work revealed a well-ordered surface with 300 A wide terraces, separated by single and double step heights, evenly spaced across the surface. STM images obtained within these terraces revealed the expected square crystal lattice [23]. The observed ˚ ) is consistent with the obserlattice spacing (3.1 A vation of sites of electron density associated with either the Ti or C atoms. Numerous surface defects (most likely surface vacancies) were also observed at this scale. Additional work is being performed to elucidate the details of these features. Table 3 Ratio of titanium metal to nonmetal components calculated from the integrated intensities of the peaks listed in Table 1 Ratio Clean
25 L O 2
75 L O 2
150 L O 2
Flash to 670°K
Ti/C Ti/O
0.97 5.3
1.15 2.9
1.10 2.6
1.05 2.9
0.94 10
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3.2. Clean VC exposed to oxygen Plotted in Fig. 3 is a series of HREEL spectra collected following sequential exposures of oxygen to the clean VC(100) surface. Each spectrum was normalized by the intensity of the elastic peak. The oxygen exposures shown here were performed at 145 K; similar results were obtained when the experiment was performed at room temperature (see, for example, spectrum A of Fig. 12 – here, a 2 L oxygen dose at room temperature produced a HREEL spectrum that is nearly identical to the corresponding low temperature experiment shown in Fig. 3D). As the dose was increased from 0.5 to 4 L, we observed the growth of a prominent loss feature at 980 cm−1. This peak may be assigned by comparison with infrared spectra of oxovanadium ions and related compounds [24– 26 ]. In these studies, stretching frequencies of the vanadium–oxygen double bond were commonly reported to be approximately 997 cm−1. Thus, we attribute the 980 cm−1 feature to an oxygen atom at a vanadium atop site. This peak saturated at an exposure of between 4 and 10 L, and no further change was detected up to 100 L exposure [23]. Similar results were obtained when exposures were conducted at room temperature (300 K ). Fig. 3F
Fig. 3. HREEL spectra of VC(100) collected after dosing with sequentially increasing amounts of oxygen at 145 K. Shown here are the clean surface (A) and after doses of 0.5 L (B), 1.0 L (C ), 2.0 L (D) and 4.0 L ( E ), respectively. Spectrum F shows the HREEL spectrum of VC(100) collected after dosing with 100 L of oxygen at 300 K.
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shows that even after large exposures (100 L) at 300 K, there are only slight differences from the results after saturation coverage at low temperature (145 K ). At saturation, the integrated intensity was approximately equal to the intensity of the metal–carbon vibrational feature, and we estimate this to be a coverage of approximately one monolayer. HREELS measurements after exposures of greater than 100 L were not possible due to our inability to recover the signal. Other possible assignments of the 980 cm−1 loss feature are the O–O stretch of molecular oxygen, with the bond softened considerably from the gas phase frequency of 1600 cm−1 by electron donation from the surface [27,28], and C–O single bond stretch [29]. The first mechanism was proposed for the adsorption of oxygen on an analogous carbide, NbC [28]. We consider these possibilities to be unlikely in our case, as a strong signal is evident at room temperature and above, where it is very unlikely that any molecular oxygen would remain adsorbed under UHV conditions. The C–O species conflict with our XPS data, presented below, and with ion scattering data in the literature [8]. After exposures of greater than 1 L, another loss feature was detected in the region of 700 cm−1. No further increase of this peak has been observed beyond 4 L up to the largest exposure of 100 L. By analogy with the assignment of the oxygen contaminant on the TiC surface, we attribute this feature to the stretching of the vanadium–oxygen single bond. We conjecture that this is oxygen bound to vanadium ( V–O–V ) at carbon vacancies, either naturally present or formed as a result of limited CO evolution. x Other spectral features found after oxygen exposure are located at 1400, 2900, and 3650 cm−1. These are identified as the bending (1400 cm−1) and stretching modes (2900 cm−1) of carbon– hydrogen species, and the stretching of the oxygen– hydrogen bond (3650 cm−1). They are likely due to adsorption of background contaminants (hydrocarbons, water, and hydrogen) from the environment during these experiments, which typically elapse 10 h of time, although time and exposure have not been monitored as independent variables. The presence of oxygen may enhance
Fig. 4. V 2p X-ray photoelectron spectra of the VC surface before (solid line) and after exposure to 100 L of oxygen (dashed line) at room temperature.
the adsorption of water and hydrogen, and the oxygen gas supply may contribute as a source of contaminants. It is significant to note that no peaks were observed in the vicinity of 1800 cm−1, where one would expect to find a carbonyl vibrational mode. XPS data collected after room temperature oxygen exposures showed broadening of the V 2p core level peaks, while the C 1s remained unperturbed. Fig. 4 shows the vanadium 2p peaks from the clean sample (solid line) and after a dose of 100 L (dashed line). We find that the peaks have broadened measurably to deeper binding energy, increasing from 2.0 to 2.3 eV. We interpret this change as indicating that the oxygen interacts with the vanadium by withdrawing electron density. This broadening was evident at a lower exposure of 32 L (the smallest used for XPS study). Difference spectra collected during sequential exposures (Fig. 5) show more clearly the shift to deeper binding energy. These spectra were obtained by subtracting the normalized spectrum of clean VC from the spectrum corresponding to each oxygen dose, thereby removing the contribution from dominant bulk VC. After dosing 32 L at room temperature (Fig. 5A), we find the growth of an asymmetric 2p peak located significantly 3/2 deeper in binding energy than the V 2p peaks of VC, which are themselves shifted a small amount
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Fig. 5. Difference spectra of the V 2p peaks during sequentially increasing doses of oxygen at room temperature. In each spectrum, the clean sample data have been subtracted from the data obtained after each dose. Spectra A and B were collected after the 32 L and 100 L doses, respectively.
from the reported metallic vanadium binding energy (512.3 eV for the 2p ) [30]. The binding 3/2 energy of the most intense portion of this feature is lower than 515.8 eV expected for VO [31] and 2 517.0 eV for V O [30]. The difference spectrum 2 5 after 100 L exposure (Fig. 5B) shows no additional increase in the deeper binding energy regions of the V 2p difference spectrum indicative of further oxidation. Fig. 6 shows the C 1s region of the VC photoelectron spectrum before (solid line) and after (dashed line) exposure to 100 L of oxygen. The intensity scale has been expanded to magnify potential changes in spectral features, and the full scale figure is shown in the inset. This figure shows no change in the C 1s spectrum due to oxygen exposure, demonstrating that the oxygen has no detectable interaction with the carbon atom of the VC(100) surface. There was no evidence of perturbation of the C 1s peak throughout sequential doses from 32 to 200 L. This is consistent with the HREELS results which showed no peaks in the regions where one would expect carbon–oxygen vibrational features. The immutability of the C 1s spectrum during oxygen adsorption, coupled with changes observed in the V 2p spectra, supports the previous assignment of the 980 cm−1 feature of
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Fig. 6. Expanded C 1s photoelectron spectra of the VC surface before (solid line) and after (dashed line) exposure to 100 L of oxygen at room temperature. No changes in this spectral feature are evident after oxygen exposure. The inset shows this figure scaled to display the entire data range.
Fig. 7. O 1s photoelectron spectra of the VC surface during sequentially increasing doses of oxygen. Spectrum A was collected from the clean surface, and spectra B and C were collected after exposure to 32 and 100 L of oxygen, respectively. Spectrum D was collected after flashing the sample to 670 K after 100 L exposure.
Fig. 3 to the oxygen vibration at a vanadium atop site. The O 1s peak of oxygen exposed VC, shown in Fig. 7, is located at about 530.0 eV, with a small tail that extends to deeper binding energy. In this figure, the spectrum of the clean sample (Fig. 7A) is shown with the sample after doses of 32 L
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(Fig. 7B) and 100 L (Fig. 7C ). We find that the peak maintained a constant binding energy, and the ratios of vanadium to oxygen determined by integrated peak intensities, listed in Table 2, show that the amount of oxygen changed little as the dose was increased beyond 32 L. This observation supports the previous conclusion that no significant chemical changes occurred beyond the HREELS saturation exposure of 4 to 10 L. Spectrum 7D will be discussed below. From XPS spectral intensities, the ratios of metal to carbon composition were determined to provide a measure of the relative quantity of carbon present at the surface after exposure to oxygen and subsequent heating. Shown in Table 2, these quantities were determined as described above. We have assumed that the absolute quantity of vanadium remained constant throughout the experiment. The V:C ratios, shown in Table 2, showed no systematic changes within our experimental uncertainty, indicating that no carbon loss was detectable during exposure to oxygen at 300 K. The final component of this table, the metal/carbon ratio after heating to 670 K, will be discussed below.
Fig. 8. HREEL spectra of TiC(100) collected after dosing with sequentially increasing amounts of oxygen at 300 K. Shown here are the clean surface (A) and after doses of 2.5 L (B), 20 L (C ), 40 L (D) and 80 L ( E ), respectively.
3.3. Clean TiC exposed to oxygen Plotted in Fig. 8 is a series of HREEL spectra collected after sequential exposures of oxygen to the TiC(100) surface at 300 K temperature. As the dose was increased, we observed the initial appearance of a small loss feature at 980 cm−1, which appeared to saturate between exposures of 10 and 20 L. In analogy to the oxygen exposed VC surface, and with the previous report of a 1008 cm−1 stretching frequency for the diatomic TiNO [24], we assign this feature to the TiNO stretching vibration of oxygen at a titanium atop site. The most prominent loss feature following greater oxygen exposure of the TiC surface was located at about 730 cm−1. This broad peak was observed only after relatively large doses of oxygen (>40 L), although its presence may be masked by existing Ti–O features at low coverage. This band is similar to the Ti–O stretching frequency in TiO [21], an 2 assignment which suggests a significant degree of
Fig. 9. Difference spectra of the Ti 2p peaks during sequentially increasing doses of oxygen at room temperature. Spectra A through C were collected after the 25 L, 75 L, and 150 L doses, respectively.
oxidation. Low temperature (150 K ) oxygen exposures produced similar results. XPS data collected after oxygen exposure at 300 K showed that Ti 2p features appeared at higher binding energy than the carbidic Ti 2p peaks, concurrent with growth on the high binding energy side of the C 1s peak. Fig. 9 shows the difference spectra of the Ti 2p peaks after sequentially increasing oxygen doses. After exposure of
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Fig. 10. Expanded C 1s photoelectron spectra of the TiC surface before (solid line) and after (dashed line) exposure to 300 L of oxygen at room temperature. The inset shows this figure scaled to display the entire data range.
25 L ( Fig. 11A) we find growth of features at deeper binding energy than the 2p peaks of TiC near 456 eV on the 2p . The 2p shift of about 3/2 3/2 1.3 eV is similar in magnitude to the shift of the V 2p peaks after the oxygen exposures discussed above. However, in contrast to VC, the Ti core levels appeared to grow features even deeper in binding energy as the exposure was increased. After the 75 and 150 L doses, the 2p difference 3/2 peak center was significantly higher than the binding energy of 454.8 eV expected for TiO [32], yet less than the value of 458.3 eV expected for TiO 2 [33]. The breadth of this Ti 2p feature, stretching 3/2 from 455 to 459 eV, suggests that after large exposures the surface is composed of a variety of TiO species. x Fig. 10 shows the C 1s peak before (solid line) and after (dashed line) exposure to 150 L oxygen at 300 K. We find the growth of a small shoulder on the carbon peak, centered at 283.2 eV, representing a shift of 1.7 eV to deeper binding energy due to a carbon–oxygen interaction. This feature was evident after exposures as low as 25 L. The relatively low binding energy of this feature allows us to rule out any contribution due to carbonaceous contaminants (284.6 eV ). The appearance of this feature indicates that, in addition to the titanium–oxygen interaction demonstrated above,
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Fig. 11. O 1s photoelectron spectra of the TiC surface after sequentially increasing doses of oxygen. Spectrum A, collected from the clean surface, shows a small amount of oxygen (5% of the surface composition) due to the bulk impurity. Spectra B through D were collected after exposure to 25 L, 75 L, and 150 L of oxygen, respectively.
the oxygen interacts with carbon on the TiC surface. The O 1s peak due to the bulk oxygen impurity on clean TiC, shown in Fig. 11A, appeared at 530.5 eV. The intensity of this peak grew and shifted to lower binding energy with oxygen doses of 25 L ( Fig. 11B), 75 L ( Fig. 11C ), and 150 L ( Fig. 11D). After a dose of 150 L, the O 1s peak was located at 530.2 eV, equal to the value expected for TiO [32]. The Ti/C ratio ( Table 3) 2 shows an abrupt increase after the 75 L dose, correlating to a decrease in the Ti/O ratio, and the growth of the 730 cm−1 Ti–O stretch ( Fig. 8) after similar doses. The Ti/C ratio apparently changed little, while the Ti/O ratio decreased only a small amount, after doubling the exposure to 150 L. Metal to carbon ratios, determined as described above and listed in Table 3, indicate that carbon was lost during exposure to oxygen. The data shown here are averaged over two separate experiments. The metal to carbon ratio increased an insignificant amount with the 25 L oxygen dose, but jumped dramatically with the 75 L dose. No further increase was found with larger doses up to 300 L. This is believed to reflect the loss of carbon as CO species from the surface. The most signifix cant loss of carbon occurred after exposures similar
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to those that produced the 730 cm−1 HREELS vibration (Fig. 8E) due to Ti–O, and the pronounced shift in binding energy of the Ti 2p peaks (Figs. 9B and C ). This marks the onset of the reaction to oxidize the TiC surface. 3.4. Effects of increased temperature HREEL spectra of VC after exposure to 2 L of oxygen ( Fig. 12) show that the intensity of the VNO stretch fell to zero as the temperature was increased from 300 to 670 K. In these data, the temperature was held at 470 K ( Fig. 12B) and 670 K ( Fig. 12C ) for 30 s before returning to 300 K for analysis. After the 670 K flash, the 510 cm−1 feature remained slightly more intense, and broader to the high frequency side, than observed in the clean VC sample ( Fig. 12A). No further changes in the spectrum were observed up to the highest temperature flash of 1300 K. Despite the dramatic decrease in the intensity of loss features associated with oxygen adsorption after heating, the XPS spectrum provided evidence of significant oxygen retention within the interfacial region. Fig. 7D shows the O 1s peak on the 100 L dosed VC surface after the sample was heated to 670 K. The center of this peak shifted 0.7 eV to
deeper binding energy and retained approximately 70% of its intensity. Concurrent with these spectral changes, the VC surface lost some of its carbon atoms upon heating. The final column of Table 2 shows that the ratio of metal to carbon composition increased by 0.15 beyond the average ratio before heating. This demonstrates that the surface region lost carbon atoms, likely in the form of evolved CO species. x HREEL spectra of oxygen covered TiC after heating also show a decrease in the intensity of oxygen features. The spectrum collected after flashing to 670 K, not shown, was marked by the absence of the TiNO feature at 980 cm−1 and the reduction of the 730 cm−1 to a broad trail on the high frequency side of the 510 cm−1 Ti–C feature. The O 1s peak showed only a 6% decrease in the integrated intensity, but no measurable change in position as a result of similar treatment. This large change in HREEL spectra, concurrent with small changes in XPS data, indicates that although a small amount of oxygen may have evolved as CO species, most of it had been driven subsurface x in TiC. In contrast to the VC surface, metal/carbon ratios for TiC (final column of Table 3) did not increase during heating within the experimental uncertainty of the technique.
4. Discussion
Fig. 12. HREEL spectra of oxygen exposed VC after increased temperature flash cycles. Spectrum A shows the surface after a 2 L exposure at 300 K. Spectra B and C were collected after raising the temperature to 70 and 670 K before returning to room temperature for data collection.
The combination of HREELS and XPS studies on the adsorption and reaction of relatively low oxygen exposures on the (100) surfaces of VC and TiC has verified a significant difference between the materials and identified surface reaction products. In this discussion, we will first summarize the experimental results, followed by potential explanations and implications of our findings. The adsorption of oxygen on VC(100) essentially saturates at relatively low exposures (~10 L) under the conditions studied. This saturation was indicated by the constant intensity and binding energy of the O 1s XPS peak and by an unchanging HREEL spectrum after exposures ranging from 10 to 100 L. The HREELS data are dominated by the 980 cm−1 peak that we have assigned as the VNO stretch of adsorbed, dissociated oxygen atop
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a vanadium atom. The XPS substrate core level results were consistent with these data, displaying a broadening on the deeper binding energy side of the V 2p peaks and no observable effect on the C 1s peak. The V/C atomic concentration ratios derived from the XPS results show no systematic changes with increasing exposure, implying that the oxygen overlayer is stable and that the substrate is not being oxidized via the loss of carbon oxides under these conditions. An additional small feature was observed in the HREELS near 700 cm−1, which is consistent with a vanadium–oxygen single bond. The low intensity of this feature shows that this reaction product was minor, perhaps arising from surface carbon vacancies or structural defects that favor the formation of a bridging oxygen species V–O–V. This surface species would also be typical of that expected from surface oxidation, indicating that a small number of surface sites may have reacted by losing carbon, as detailed for TiC below. Upon heating to temperatures in the range of 470 to 670 K, the O/VC HREELS data change considerably while most of the oxygen was retained by the substrate as indicated by XPS. Specifically, the VNO and V–O–V vibrations disappeared, while features may have grown under the V–C stretch in the spectrum. The O 1s peak lost some intensity and shifted to higher binding energy, indicating the formation of a different chemical species. The V/C concentration ratio increased, consistent with some loss of surface carbon as expected for the onset of oxidation. The drastic HREELS changes coupled to the XPS data may indicate that much of the remaining oxygen has been driven below the surface by heating. The results on TiC(100) were significantly different, as the XPS oxygen signal did not saturate in the exposure regime studied. The initial exposures (<40 L) populated a site with a small HREELS signal near 980 cm−1, similar to the major feature observed on VC and thus assigned as TiNO. With increasing exposures, however, a lower energy HREELS feature grew to dominate the spectrum. This feature was in the region expected for Ti–O–Ti surface species, such as would be expected in TiO . The XPS core levels 2 showed changes consistent with these two different
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surface products. At lower exposures, the Ti 2p levels were broadened to deeper binding energy and the Ti/C concentration ratio shows little change. With increasing exposure, the Ti/C concentration ratio increased and deeper binding energy features grew on the Ti 2p peaks, demonstrating the loss of carbon and the oxidation of the surface at 300 K. Significant changes were also observed in the C 1s XPS spectrum of O/TiC, showing that oxygen was interacting with surface carbon. This interaction was evidenced by a deeper binding energy shoulder on the C 1s peak. No definitive HREELS peaks characteristic of carbon–oxygen bond frequencies were found as would be expected from the measurable perturbation of the C 1s XPS peak. It is conceivable that oxygen could occupy an adsorption site involving both titanium and carbon atoms, and the stretching frequency of this species could be present in the broad feature observed between 500 and 1000 cm−1 (see Fig. 8E). The measured Ti/O XPS concentration ratio was roughly half that found for VC after the larger exposures, consistent with more extensive surface reaction on TiC. Finally, heating to 670 K did alter the HREEL spectrum, but had little effect on the relative concentration of the surface species, showing that further oxidation was not being thermally driven. For either material, the spontaneous formation of the metal oxide and a carbon oxide is thermodynamically favored by a considerable amount. In fact, if left in laboratory ambient, TiC and VC will form oxide layers. The lack of oxidation of VC(100) under our experimental conditions indicates the presence of a kinetic barrier that is not present for TiC. We surmise that the surface reaction products formed can either enhance or inhibit further reaction. These reaction products are a result of the TiC and VC surface geometric and electronic structures, and the differences between the materials in these regards may explain the experimental results. The difference between pristine, defect-free surfaces of TiC and VC is the identity of the metal atom. Chemically, this difference results in the presence of an additional valence electron in a formula unit of VC. Considering the electronic
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structure from a localized chemical bonding perspective, the highest occupied valence orbitals on TiC are predominantly C 2p in nature, basically representing a filled valence band. The additional electron present in VC must occupy a molecular orbital that is predominantly V 3d in nature, and this has been verified by variable photon energy photoelectron spectroscopy on these materials [5]. Oxygen seeks electrons on a surface, hence the initial interaction of dissociated oxygen with the two materials will occur at sites with the most accessible electron density. Simply, this will be the C atoms on TiC and the V atoms on VC. The thermodynamically stable and physically volatile nature of gas phase carbon oxides facilitates their loss on TiC, and the simultaneous formation of a titanium oxide. However, the VNO species formed on VC is stable enough under the conditions of our experiment to preclude further reaction with this material. Elevated temperature can drive the oxidation of VC, as can extended time under an oxidizing environment such as the laboratory ambient. Theoretical analysis of adsorption on transition metal carbides, from the perspective of the surface electronic structure, supports this picture of oxygen adsorption [6 ]. It is also clear that these surfaces have defects, usually believed to arise from carbon vacancies resulting from the tendency of these materials to be substoichiometric. Although our quantitative XPS results of the clean surfaces do not support this notion, we believe that STM images offer compelling proof of the existence of vacancies that we will assume are missing carbon atoms [23]. Therefore, some of the surface chemistry must account for these sites. We propose that the majority species observed in HREELS arise from oxygen adsorption and chemistry on well-defined terrace sites, while minority species result from the interaction of oxygen with surface defects. We believe this because STM of TiC has shown that most of the surface is well ordered, but a significant number of defect sites are evident [23]. The minority species on TiC data produced the 980 cm−1 feature that was evident at low exposures. The similarity between this feature and that identified as the major species on VC implies a similar adsorption site, i.e. a Ti atom with addi-
tional electron density. This can most likely happen in one of two ways. First, carbon vacancies will cause the surface to appear more metallic, shifting electron density that would have resided on the carbon back to the neighboring titanium atoms. Thus, the titanium sites adjacent to vacancies may form this surface product. A second possibility is defects due to the presence of oxygen contaminants substituting for carbon in the TiC crystal structure. The effect of this substitution is to add additional valence electrons to the surface (similar to the changes between TiC and TiN ), and perhaps creating added electron density to neighboring Ti atoms. Because we were unable to remove all of the surface oxygen, we cannot distinguish between these possibilities. However, ion scattering studies indicate that the initial adsorption of oxygen on TiC may fill defect sites [8,34]. The presence of C vacancies on VC will also create additional electron density on neighboring V atoms and create a site where oxygen can bridge between two V atoms without interference from a carbon atom. The HREELS data indicate that two different sites are occupied at low exposure on VC, with no apparent preference. The minor species has been assigned as a V–O single bond that we believe is more likely to occur where carbon vacancies exist. It is possible that some of the oxygen atoms forming VNO species could also be bonding at defect sites, with a portion forming the bridge bonded species. It is conceivable that oxygen could bridge between two vanadium atoms in a pristine lattice site. Until the concentration of defects on our VC surfaces has been determined absolutely, we will be unable to draw firm conclusions. Our quantitative XPS (metal/carbon ratios) suggests that the number of defects on the two materials should be similar. In general, our results are consistent with previous ion scattering studies reported in the literature on these surfaces [8]. In that work, the bonding of oxygen with carbon on TiC was inferred from the loss of scattering intensity from carbon. However, the relatively facile oxidation of the TiC surface makes it unclear whether any blocking of C sites by adsorbed oxygen is observable, or whether carbon monoxide is formed so quickly that capturing significant numbers of reaction
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intermediates such as surface carbonyls is not possible on the time scale of these experiments. We have shown that oxygen does perturb the C 1s XPS peak, but could find no HREELS features that unambiguously identify C–O species. It is likely that the ISS results showed a combination of site blocking and carbon loss due to surface oxidation. Finally, the use of electronic structure arguments to predict surface chemistry can potentially enable predictions of the behavior of similar materials and other adsorbates. The results of this paper indicate that electron withdrawing adsorbates will initially bond with surface sites having the most accessible electron density, V on VC and C on TiC. As a result, the isoelectronic TiN and VC might well be predicted to behave similarly to oxygen exposure under these conditions. The formation of nitrogen oxides is not as favorable thermodynamically, but the driving force may well be the very stable metal oxide. Future experiments are planned on single crystal TiN and with different adsorbates to determine if these trends can be generalized.
5. Summary and conclusions We have demonstrated methods for obtaining reproducible, clean TiC and VC surfaces, and the properties of these surfaces have been investigated with spectroscopic probes. It has been shown that oxygen behaves differently on these surfaces; oxidizing the TiC surface, after binding to the carbon atom and releasing CO , while preferring to bond x with the metal atom on VC, forming a metastable overlayer. Explanation of these preferences is likely the different electronic structures of the materials. The additional electron in the VC formula unit resides in a predominantly V 3d orbital, while the highest occupied orbital in TiC is predominantly C 2p. These experiments have shown that TMC surfaces can form strong bonds to oxygen containing species, indicating that lubricant additives containing oxygen will likely interact with these materials. We are presently studying these interactions, where more complex and technologically relevant materi-
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als are exposed to these TMC surfaces under similar conditions.
Acknowledgements This work is supported by The Aerospace Corporation IR&D program funded by the Space and Missile Systems Center (SMC ) of the USAF under contract number F04701-93-C-0094.
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