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Surface segregation of Sb in doped TiO2 rutile
apolteo
surface @eief~e Applied Surface Science90 (1995) 289-295
ELSEVIER
Surface segregation of Sb in doped TiO 2 rutile Antonino Gulino a,*, Guglielmo G. Condorelli a, Ignazio Fragalh a, *, Russell G. Egdell b., a Dipartimento di Scienze Chiraiche Universith di Catania, V.le A. Doria 6, 95125 Catania, Italy b Inorganic Chemistry Laboratory. South Parks Road, Oxford OXI 3QR, UK
Received 8 March 1995; acceptedfor publication 2 June 1995
Abstract The surface concentration of Sb in doped TiO2 mille ceramics (Til_5/4xSbxO 2 0 < x < 0.1), has been measured by means of angle-resolved core level X-ray photoelectron spectroscopy (AR-XPS). Depth profiles have been obtained by alternating Ar+-ion bombardment with core level measurements. At low doping levels Sb segregates by substitutional replacement of Ti in a large number of ionic planes whilst at higher Sb doping levels there is evidence of a new Sb-Ti-O amorphous surface phase whose thickness involves about five ionic planes. A rationalization of the monotonic decrease of the work function throughout the doping range studied has been proposed.
1. Introduction It is now widely recognized that the electrical, mechanical and other technological properties of ceramic oxides are controlled to a large extent by segregation of impurities and dopants to surface and grain boundaries [1]. In these cases the structure and stability of the surface are governed by the particular dopant chemical environment. Knowledge of surface composition is of course fundamental for a correct understanding of the surface electronic structure and catalytic properties. The technological applications of TiO 2 as a white pigment and in the field of catalysis and electrocatalysis have led to widespread interest in the surface properties of the material: it is the case that TiO 2 is
* Corresponding authors. Tel.: + 39-95-336578; Fax: + 39-95580138.
among the most studied oxides by the contemporary techniques of surface science [2]. The measurement of compositional distribution (or depth profile) in the near-surface layer is more critical and meaningful than the average composition. To determine the composition as a function of distance from the surface with reasonable high resolution, the combination of sputter sectioning and chemical analysis is the most universally applicable technique [3,4]. A pioneering study involving the use of surface probes to investigate ion bombardment modification of TiO 2 involved use of 500 eV Ar + ion bombardment to generate surface O-vacancy defects. The defective surfaces were studied by ultraviolet photoelectron spectroscopy (UPS), electron loss spectroscopy (ELS) and low-energy electron diffraction (LEED) [5]. UPS spectra of a UHV oxygen annealed (110) TiO 2 single crystal surface show evidence of a
0169-4332/95/$09.50 © 1995 Elsevier Science B.V. All rights reserved SSDI 0169 -43 32(95)00160-3
290
A. Gulino et al. / Applied Surface Science 90 (1995) 289-295
weak surface state at about 1 eV binding energy. This is associated with electrons associated with oxygen deficiency trapped on Ti 3+ sites. After 500 eV Ar ÷ ion bombardment there is loss of long-range order and the LEED pattern disappears, suggesting amorphization of the surface. Moreover, the energy of the extrinsic surface state shows a steady shift and there is a concomitant loss of oxygen. The altered surface state appeared to involve the pairing of Ti 3+ ions. As the density of Ti 3÷ pairs increases, small areas of T i / O 3 begin to form on the surface [5-8]. The UPS and ELS spectra resemble those of vacuum-fractured Ti20 3 very closely [4]. In our own previous work we have studied the effect of the Sb-doping on TiO 2 by photoemission spectroscopy [9]. Atomically clean surfaces were prepared by annealing samples in UHV. Pronounced segregation of Sb to polycrystalline TiO 2 surfaces was found [9]. At the same time the 1 eV state associated with Ti 3÷ is replaced by a bandgap structure much closer to the valence-band edge. In the present investigation we try to extend the understanding of the Sb accommodated on the doped-TiO 2 rutile oxide surface by means of angleresolved core level X-ray photoelectron spectroscopy (AR-XPS) and depth profiles obtained by alternating Ar ÷ ion bombardment with core level measurements. The samples were not annealed in UHV but were subjected to ex-situ annealing in air. Thus in combination with the earlier work we are able to learn about the influence of the annealing atmosphere on surface composition.
2. Experimental details 1%, 4%, 5%, 7% and 10% Sb-doped TiO 2 samples were prepared as already described [9]. The sample powders were pressed into pellets between tungsten carbide dies and fired for several days at 1000°C in recrystallized alumina boats with intermediate regrinding and repelletization. The final samples were robust, pale yellow ceramic discs. X-ray powder diffraction data were identical to those previously obtained [9]. The absolute Sb-doping levels, established by inductively coupled plasma emission spectroscopy (hereafter ICPES), were always around 70% + 10% of the nominal doping level. This is
presumably due to volatilization of Sb40 6 during pellet sintering [9]. AR-XPS, at many different angles (90 °, 80 °, 60 °, 45 °, 30 °, 15°, 11 °, 9 °, 7 °, 5 °, 3°), were measured in a PHI 5600 Multi Technique System, which offers a good control of the electron take-off angle. The spectrometer is equipped with a dual anode X-ray source; a spherical capacitor analyser (SCA) with a mean diameter of 279.4 mm; an electrostatic lens system Omni Focus III. Depth profiles were obtained by alternating XPS core level measurements with mild (1 kV, beam current 0.80 /xA) Ar ÷ ion sputter etching every 6 s. The ion gun was positioned 60 ° relative to the sample surface plane. The argon ion beam was rastered over areas of 2 mm × 2 mm. The photoelectron take-off angle, relative to the surface plane, was 45 °. The nominal analyser resolution was set to 400 meV. Samples were mounted on Mo stubs and held in position with Mo wires. Spectra were excited with AI Kc~ radiation. Structure due to satellite radiation has been subtracted from the spectra before the data processing. The XPS peak intensities were obtained after Shirley background removal [10]. Since the major objective of the present study is to probe the surface Sb distribution, care was taken to introduce hot, furnace-annealed, ceramic discs in order to avoid procedures for in situ cleaning as ion bombardment. The remaining adventitious carbon contamination of the surfaces (284.6 eV binding energy) gave a measured C l s / O Is XPS intensity ratio of around 4/100. The O I s core peak showed a single component, with no evidence of high binding energy shoulders due to water or hydroxide contamination. Owing to the semiconducting nature of the Sb-doped TiO 2 ceramics [11], sample charging was not a major problem. Additional analyses of sample surfaces were carried out by Scanning Auger Microscopy (SAM) and Scanning Electron Microscopy (SEM). SAM analyses were performed using a 5 keV primary electron beam, an electron beam current of 5 nA and a resolution better than 0.2 /~m whilst SEM results were obtained using a 9 keV electron beam and a resolution better than 0.5 /xm. SAM measurement provided information on the spatial uniformity of Sb distribution as well as on the absence of Sb islands on the surface in all the studied samples. The SEM micrography of the 10% Sb-doped ceramic sample
A. Gulino et al. / Applied Surface Science 90 (1995) 289-295
291
(Fig. 1) shows the presence of spherical, particles having sizes ranging from 2000-4000 A on the surface.
3. Results and discussion
3.1. Angle resolved measurements Normal-emission core level photoelectron spectra of Sb-doped TiO 2 ceramic samples, measured in a binding energy region containing both O ls and Sb 3d peaks are shown in Fig. 2. The Sb 3d5/2 peak overlaps the O ls peak almost perfectly, so the Sb 3d3/2 peak only can be used in quantification of the surface Sb content. The intensity ratio between Sb 3d3/2 and Ti 2p peaks is therefore used to derive values for the effective surface Sb/Ti ratio, taking into account the relevant atomic sensitivity factors [12]. It is reasonable to assume that analyser transmission functions for Sb3d and Ti 2p peaks are the same because they have similar kinetic energies. Values obtained in this way, for different emission angles relative to the surface normal, are given in Table 1. It becomes evident that the intensity of the Sb 3d peaks is very much greater than it would be expected on the basis of the nominal doping levels and that the Sb/Ti surface ratios vary less than variations in the nominal doping levels. Moreover, the enhanced
520
I
I
I
l
I
525
550
535
540
545
550
Binding Energy / eV Fig. 2. AlKot excited XPS of (top) 4% Sb-doped TiO 2 and (bottom) 10% Sb-doped TiO 2 measured in the O l s and Sb3d binding energy region. The Sb3ds/2 peak overlaps the O Is peak but the Sb3d3/2 peak is clearly seen at about 540 eV binding energy. Structure due to satellite radiation has been subtracted from the spectra.
Sb intensity at grazing emission angles indicates different Sb concentration in the ionic layers probed by XPS. Consider a series of cation containing planes with interplanar separation d. Assuming that the top n layers are enriched in a dopant it is possible to write a general expression for the angular variation of the XPS intensity ratio l(Sb 3d)/l(Ti 2p) corrected for the relevant atomic sensitivity factors [13] I(Sb 3 d ) / l ( T i 2p) =(O+xH)/[(1-O)+(1-x)H],
(1)
were H = K/(1 - K) and K = e x p ( - n d / A s i n a ) . Table 1 Apparent atomic ratios from Sb3d and Ti2p core intensities in relevant AR-XPS
Fig. I. Scanning electron micrograph of the 10% Sb-doped TiO 2 ceramic sample.
Nominal doping
Experimental [Sb/Ti] a from AR-XPS
Surface Sb from
level %
90'
30 °
9°
3°
Eq. (1)
1.0 4.0 5.0 7.0 10.0
0.05 0.14 0.15 0.16 0.25
0.06 0.17 0.18 0.19 0.26
0.07 0.19 0.21 0.20 0.29
0.08 0.20 0.21 0.23 0.30
0 0 0 0 0
a Estimated uncertainties are _4-10%.
= 0.07 =0.16 =0.17 = 0.19 = 0.23
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A. Gulino et al. / Applied Surface Science 90 (1995) 289-295
Here A represents the photoelectron inelastic mean free path; x represents the nominal doping level; t~ is the photoelectron take-off angle relative to the surface plane; 0 is the Sb occupancy of the Sb-enriched layers. We take A = 12.5 A for both Sb3d and Ti2p core level peaks [12] and d = 3.24 (which corresponds to the separation between adjacent Ti-containing planes in the (110) direction [9]). The intensity ratio I(Sb3d)/l(Ti2p) pattern observed for the 4% Sb-doped TiO 2 at different emission angles is consistent with a 0 value = 0.16 if the Sb enrichment is confined to the four top-most ionic layers. Assuming n = 4 the increases of the S b / T i intensity ratio on going from 90 ° to 3° are in good agreement with the angular variation predicted by Eq. (1). By contrast, if we assume enrichment in the topmost ionic layer we obtain a higher occupancy of Sb which is inconsistent with the observed intensity ratio at grazing emission angles. Here, using the S b / T i = 0.14 value, obtained for the 90 ° emission peak and assuming n = 1, Eq. (1) predicts a 0 value of 0.40. At the most grazing measured emission angle (3 °) this value implies an S b / T i ratio of 0.66. This calculated intensity value is totally inconsistent with the observed value of 0.20. Similar results have been obtained for the 5% Sb-doped sample (Table 1). More muted variation in intensity with take-off angle is found for the 10% Sb-doped sample. In fact, little variation of the S b / T i intensity ratio is observed on going from 90 ° to 30° emission peaks and, even at 3° grazing emission angle, the intensity increase is very much less pronounced, than for the 4% doped system. In this case, the calculated value of 0 = 0.23 is consistent with an enrichment involving five ionic planes. This picture is consistent with Transmission Electron Microscope (TEM) results which show an Sb-enriched layer approximately 15 ~, thick [9]. These results are different to those previously reported for Sb-doped SnO 2, another stable oxide which adopts the rutile structure [2]. Here Egdell and coworkers rationalized the high XPS Sb intensity in terms of replacement of Sn by Sb in the topmost ionic layer of the sample with full saturation of surface cation sites at 3% Sb bulk doping [14-17]. Obviously one limitation of the present approach is the use of a model which assumes a smooth (110) crystal surface to analyse experimental data obtained from a polycrystalline ceramic material.
-1.2-
-1.6-
~ -2.0.
-2.4-
-5.5
,
,
,
-5.0
-4.5
.-4.0
, - .5
; - .0
,
,
-2.5
-2.0
log([Sb]/[Wi]) Fig. 3. SurfaceSb3d to Ti2p intensityratio plotted as a function of bulk doping level.
An experimental estimate of the segregation enthalpy of Sb in Sb-doped TiO 2 rutile may be obtained by using the following expression [18]: ( S b / T i ) s = (Sb/Ti)b e x p ( - AHsegr/RT ).
(2)
A plot of the natural logarithms of these ratios, in the present doping range, is given in Fig. 3. The cleaning procedure of the samples involved prolonged annealing at 1000°C followed by rapid cooling to room temperature. Hence we assume that in XPS the equilibrium surface obtained at 1000°C is being probed. The value obtained for the heat of segregation is - 2 . 3 8 0 kJ mol -J. This value is less negative than the corresponding - 2 9 . 4 kJ tool-~ calculated for the Sb-doped SnO z parent system [14]. Of course, the heat of segregation depends on the bulk atomic ratio, so the value obtained is not surprising since two of the present Sb-doped TiO 2 samples (7%, 10%) exceed the solubility limit of antimony in TiO 2 rutile [9,11].
3.2. Ar + ion bombardment of Sb-doped TiO 2 It has been shown that oxygen vacancies are the dominant type of defect formed under ion bombardment of oxides [19,20]. This implies that such bombardment preferentially reduces oxide surfaces. In particular the (110) surface of TiO 2 reduces under a few hundred eV Ar + ion bombardment to a stoichiometry close to that of Ti203 [5-8]. The present
A. Gulino et al. /Applied Surface Science 90 (1995) 289-295
293
Ti f
~u ~
40.
20I
I
I
450
460 Binding energy / eV
470 0
Fig. 4. AIKo~ excited XPS spectra of Ti2p core level for 4% Sb-doped TiO 2 before (bottom) and after (top) sputtering.
10
20
30
40
50
60
Sputter time (min) Fig. 6. Depth profile of 4% Sb-doped TiO 2.
XPS spectra of Ti 2p core levels, obtained after sputtering, show shoulders on the low-binding energy side of the doublet evident for the TiO 2 itself (Fig. 4). These shoulders are due to the presence of Ti 3+ ions on the slightly reduced TiO 2 surfaces [5-8] which assume a pale grey colouration [9]. Finally, steady-state S b / T i compositions close to the ICPES values have been attained after long sputtering of all the analyzed systems showing that under our conditions the preferential sputtering of Ti or Sb is not a major problem as in the case of tin-antimony mixed oxide [21]. Figs. 5 and 6 show depth profiles of 10% and 4% Sb-doped TiO 2 respectively. Ti
80-
60O
40-
< 20-
Sb ' 0
i 1
,
i 2
,
i 3
•
i 4
'
i 5
,
i 6
,
i 7
,
l
•
S
Sputter time (rain) Fig. 5. Depth profile of 10% Sb-doped TiO 2.
i 9
'
i 10
It may be seen that surface compositions differ noticeably from the bulk values. All the samples exhibit the following features: the surface is enriched with Sb; the greater the nominal doping level the greater the surface S b / T i ratio; all the samples, at zero sputter time, show surface atomic compositions in tune with the 0 values calculated from the angleresolved measurements. The main difference between the differently doped samples is the thickness of the enriched layer. In fact, in the case of the 10% Sb-doped sample (Fig. 5) the Sb rapidly decreases with the increase of the sputter time and, after 2.5-3 min, the Sb percentage is close to the ICPES value. By contrast, the 4%-Sb-doped sample (Fig. 6) shows a different behaviour. Here the Sb content decreases slowly with the sputter time to reach the ICPES value after 50-60 min. These two extreme behaviours are obviously related to different thickness of the Sb-enriched layers in the two samples and suggest that it becomes possible to distinguish two limiting models for Sb-doped TiO 2 on the basis of the Sb doping level which, in turn, may be below or above the solubility limit of Sb in TiO 2. It has been seen in fact, by TEM measurements, that at doping levels higher than 5% the Sb-doped TiO 2 shows clear evidence of an amorphous S b - T i - O phase at surface and grain boundaries which, in the 10% doped sample, is approximately 15 ~, thick thus corresponding to 4 - 5 ionic layers [9]. Therefore, the rapid decrease of the Sb atomic percentage observed
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A. Gulino et al. / Applied Surface Science 90 (1995) 289-295
in the 10% Sb-doped T i t z rutile, upon Ar ÷ ion bombardment is due to the sputtering of this amorphous surface layer. Conversely, the behaviour of the 4% Sb-doped sample suggests that the number of ionic planes involved in the Sb-enrichment is underestimated by AR-XPS. In fact, 50-60 min sputtering time is needed to reach the ICPES S b / T i ratio thus suggesting that segregation involves a large number of ionic planes with the Sb percentage diminishing smoothly on going from the topmost ionic layer to the bulk. Obviously the present behaviour could be the results of a radiation-enhanced diffusion mechanism which depends on vacancies and other lattice defects generated by (heavy particle) irradiation [22] and could cause the Sb to move deeper into the substrate. Nevertheless, there is no explanation for the fact that this mechanism is operative for the 4% Sb-doped sample and not for the 10% doped sample, while identical operative conditions are always adopted. All these observations provide a sensible rationale for the monotonic decrease of the work function (as previously determined from the cut-off in secondary electron emission in UPS) throughout the doping range of samples presently studied [9]. As a matter of fact, Sb(III) species, having a formal 5s 2 electron configuration, are accommodated at the surface [9]. In these conditions, the non-centrosymmetric crystal eV 6
\
4
r
0
2
i
I
i
i
f
4 6 8 Doping Level / %
10
Fig. 7. Variationin work functionq~ with nominaldoping levelin Sb-doped Tit 2 ceramics as determined from the cut-off in secondary electronemission in UPS.
field at the surface sites leads to 5s-5p mixing. The 5s-5p hybrids have electron densities pointing out of the surface plane, but the surface dipole associated with these electrons is expected to be less than when bridging oxygen bonds to two Ti(IV) ions [9]. Thus the reduced surface dipole lowers the sample work function. The present surface Sb concentrations in the topmost ionic layer of the differently doped samples (0) correlate well with the monotonic decrease of the work function (Fig. 7). Finally the decrease in the intensity of the bandgap emission at 1.8 eV binding energy in the 7% and 10% Sb-doped samples (already observed with UPS measurements) [9] can be safely related to the different geometric structure of the second amorphous S b - T i - O phase present in the surfaces and grain boundaries. These states are now hidden under the more intense O2p based valence band as has been previously observed for Sb-doped SnO 2 [ 15,17].
4. Conclusion The combined application of angle-resolved XPS and depth profile measurements provides a powerful tool of probing surface composition of doped metal oxides. The present study highlights a pronounced surface segregation of Sb to polycrystalline doped T i t z rutile sample. In all the present samples the calculated Sb concentrations in the enriched topmost planes are in tune with Sb percentages obtained at zero sputter time. In the case of 4 - 5 % Sb-doped sample, Sb segregation involves a large number of planes with the Sb concentration varying smoothly on going from the top most ionic layer to the bulk. By contrast, in the case of heavily Sb-doped sample (7% 10%), in which the Sb exceeds the substitutional solubility limit, a new Sb-enriched S b - T i - O amorphous surface phase emerges whose thickness involves about five ionic planes with a rapid return of the bulk [Sb/Ti] ratio. As far as the influence of the annealing atmosphere on surface composition is concerned, we believe that the oxygen surface deficiency, due to in-vacuo annealing [9], favours the increase of the Sb surface concentration since, as found in an earlier study [9], the [D]-Sb(III)20 2] stable (110) termination is required to maintain the electrical neutrality.
A. Gulino et al. / Applied Surface Science 90 (1995) 289-295
Here [] indicates the removal of bridging oxygen atoms. This observation is tuned well with the fact that the 1 eV state, associated with Ti 3÷, is replaced by a bandgap structure much closer to the valenceband edge (1.8 eV) on passing from undoped TiO 2 to Sb-doped TiO 2 [9]. The lower Sb surface concentration found in the present air annealed sample is therefore in tune with the absence of oxygen deficiency. One of the major goals of the present investigation is the rationalization of the monotonic decrease of the work function with the nominal doping level in Sb-doped TiO 2 ceramics.
References [1] R.G. Egdell and S.C. Parker, in: Science of Ceramic Interfaces, Ed. J. Nowotny, Material Science Monographs, Vol. 75 (Elsevier, Amsterdam, 1991) pp. 41-78. [2] V.E. Henrich and P.A. Cox, The Surface Science of Metal Oxides (Cambridge University Press, Cambridge, 1994). [3] N.Q. Lam, Surf. Interf. Anal. 12 (1988) 65. [4] O. Auciello and R. Kelly, Ion Bombardment Modification of Surfaces (Elsevier, Amsterdam, 1984). [5] V.E. Henrich, G. Dresselhaus and H.J. Zeiger, Phys. Rev. Lett. 36 (1976) 1335.
295
[6] R. Sadeghi and V.E. Henrich, J. Catal. 109 (1988) I. [7] H. Mohamed, H.R. Sadeghi and V.E. Henrich, Phys. Rev. B 37 (1988) 8417. [8] J.L. Sullivan, S.O. Saied and I. Bertoti, Vacuum 42 (1991) 1203. [9] A. Gulino, A.E. Taverner, S. Warren, P. Harris and R.G. Egdell, Surf. Sci. 315 (1994) 351. [10] M. Repoux, Surf. Interf. Anal. 18 (1992) 567. [1 l] N. Morita, T. Endo, T. Sato and M. Shimada, J. Solid State Chem. 68 (1987) 106. [12] D. Briggs and M.P. Seah, Eds., Practical Surface Analysis (Wiley, Chichester, 1983). [13] M.P. Seah, Surf. lnterf. Anal. 2 (1980) 222. [14] R.G. EgdeU, W.R. Flavell and P. Tavemer, J. Solid State Chem. 51 0984) 345. [15] P.A. Cox, R.G. Egdell, C. Harding, W.R. Patterson and P. Tavemer, Surf. Sci. 123 0982) 179. [16] P.A. Cox, R.G. Egdell, C. Harding, A.F. Orchard, W.R. Patterson and P. Tavemer, Solid State Commun. 44 (1982) 837. [17] R.G. Egdell, in: Science of Ceramic Interfaces 1I, Ed. J. Nowotny (Elsevier, Amsterdam, 1994) pp. 527-565. [18] L. Cao, R.G. Egdell, W.R. Flavell, K.F. Mok and W.C. Mackrodt, J. Mater. Chem. (1991) 785. [19] H.M. Naguib and R. Kelly, Rad. Eft. 25 (1975) I. [20] J.B. Malherbe, S. Hofmann and J.M. Sanz, Appl. Surf. Sci. 27 (1986) 355. [21] Y. Cross and D.R. Pyke, J. Catal. 58 (1979) 61. [22] A. Benninghoven, F.G. Riidenauer and H.W. Wemer, Secondary Ion Mass Spectrometry, Basic Concepts, Instrumental Aspects, Applications and Trends (Wiley, New York, 1987).
surface @eief~e Applied Surface Science90 (1995) 289-295
ELSEVIER
Surface segregation of Sb in doped TiO 2 rutile Antonino Gulino a,*, Guglielmo G. Condorelli a, Ignazio Fragalh a, *, Russell G. Egdell b., a Dipartimento di Scienze Chiraiche Universith di Catania, V.le A. Doria 6, 95125 Catania, Italy b Inorganic Chemistry Laboratory. South Parks Road, Oxford OXI 3QR, UK
Received 8 March 1995; acceptedfor publication 2 June 1995
Abstract The surface concentration of Sb in doped TiO2 mille ceramics (Til_5/4xSbxO 2 0 < x < 0.1), has been measured by means of angle-resolved core level X-ray photoelectron spectroscopy (AR-XPS). Depth profiles have been obtained by alternating Ar+-ion bombardment with core level measurements. At low doping levels Sb segregates by substitutional replacement of Ti in a large number of ionic planes whilst at higher Sb doping levels there is evidence of a new Sb-Ti-O amorphous surface phase whose thickness involves about five ionic planes. A rationalization of the monotonic decrease of the work function throughout the doping range studied has been proposed.
1. Introduction It is now widely recognized that the electrical, mechanical and other technological properties of ceramic oxides are controlled to a large extent by segregation of impurities and dopants to surface and grain boundaries [1]. In these cases the structure and stability of the surface are governed by the particular dopant chemical environment. Knowledge of surface composition is of course fundamental for a correct understanding of the surface electronic structure and catalytic properties. The technological applications of TiO 2 as a white pigment and in the field of catalysis and electrocatalysis have led to widespread interest in the surface properties of the material: it is the case that TiO 2 is
* Corresponding authors. Tel.: + 39-95-336578; Fax: + 39-95580138.
among the most studied oxides by the contemporary techniques of surface science [2]. The measurement of compositional distribution (or depth profile) in the near-surface layer is more critical and meaningful than the average composition. To determine the composition as a function of distance from the surface with reasonable high resolution, the combination of sputter sectioning and chemical analysis is the most universally applicable technique [3,4]. A pioneering study involving the use of surface probes to investigate ion bombardment modification of TiO 2 involved use of 500 eV Ar + ion bombardment to generate surface O-vacancy defects. The defective surfaces were studied by ultraviolet photoelectron spectroscopy (UPS), electron loss spectroscopy (ELS) and low-energy electron diffraction (LEED) [5]. UPS spectra of a UHV oxygen annealed (110) TiO 2 single crystal surface show evidence of a
0169-4332/95/$09.50 © 1995 Elsevier Science B.V. All rights reserved SSDI 0169 -43 32(95)00160-3
290
A. Gulino et al. / Applied Surface Science 90 (1995) 289-295
weak surface state at about 1 eV binding energy. This is associated with electrons associated with oxygen deficiency trapped on Ti 3+ sites. After 500 eV Ar ÷ ion bombardment there is loss of long-range order and the LEED pattern disappears, suggesting amorphization of the surface. Moreover, the energy of the extrinsic surface state shows a steady shift and there is a concomitant loss of oxygen. The altered surface state appeared to involve the pairing of Ti 3+ ions. As the density of Ti 3÷ pairs increases, small areas of T i / O 3 begin to form on the surface [5-8]. The UPS and ELS spectra resemble those of vacuum-fractured Ti20 3 very closely [4]. In our own previous work we have studied the effect of the Sb-doping on TiO 2 by photoemission spectroscopy [9]. Atomically clean surfaces were prepared by annealing samples in UHV. Pronounced segregation of Sb to polycrystalline TiO 2 surfaces was found [9]. At the same time the 1 eV state associated with Ti 3÷ is replaced by a bandgap structure much closer to the valence-band edge. In the present investigation we try to extend the understanding of the Sb accommodated on the doped-TiO 2 rutile oxide surface by means of angleresolved core level X-ray photoelectron spectroscopy (AR-XPS) and depth profiles obtained by alternating Ar ÷ ion bombardment with core level measurements. The samples were not annealed in UHV but were subjected to ex-situ annealing in air. Thus in combination with the earlier work we are able to learn about the influence of the annealing atmosphere on surface composition.
2. Experimental details 1%, 4%, 5%, 7% and 10% Sb-doped TiO 2 samples were prepared as already described [9]. The sample powders were pressed into pellets between tungsten carbide dies and fired for several days at 1000°C in recrystallized alumina boats with intermediate regrinding and repelletization. The final samples were robust, pale yellow ceramic discs. X-ray powder diffraction data were identical to those previously obtained [9]. The absolute Sb-doping levels, established by inductively coupled plasma emission spectroscopy (hereafter ICPES), were always around 70% + 10% of the nominal doping level. This is
presumably due to volatilization of Sb40 6 during pellet sintering [9]. AR-XPS, at many different angles (90 °, 80 °, 60 °, 45 °, 30 °, 15°, 11 °, 9 °, 7 °, 5 °, 3°), were measured in a PHI 5600 Multi Technique System, which offers a good control of the electron take-off angle. The spectrometer is equipped with a dual anode X-ray source; a spherical capacitor analyser (SCA) with a mean diameter of 279.4 mm; an electrostatic lens system Omni Focus III. Depth profiles were obtained by alternating XPS core level measurements with mild (1 kV, beam current 0.80 /xA) Ar ÷ ion sputter etching every 6 s. The ion gun was positioned 60 ° relative to the sample surface plane. The argon ion beam was rastered over areas of 2 mm × 2 mm. The photoelectron take-off angle, relative to the surface plane, was 45 °. The nominal analyser resolution was set to 400 meV. Samples were mounted on Mo stubs and held in position with Mo wires. Spectra were excited with AI Kc~ radiation. Structure due to satellite radiation has been subtracted from the spectra before the data processing. The XPS peak intensities were obtained after Shirley background removal [10]. Since the major objective of the present study is to probe the surface Sb distribution, care was taken to introduce hot, furnace-annealed, ceramic discs in order to avoid procedures for in situ cleaning as ion bombardment. The remaining adventitious carbon contamination of the surfaces (284.6 eV binding energy) gave a measured C l s / O Is XPS intensity ratio of around 4/100. The O I s core peak showed a single component, with no evidence of high binding energy shoulders due to water or hydroxide contamination. Owing to the semiconducting nature of the Sb-doped TiO 2 ceramics [11], sample charging was not a major problem. Additional analyses of sample surfaces were carried out by Scanning Auger Microscopy (SAM) and Scanning Electron Microscopy (SEM). SAM analyses were performed using a 5 keV primary electron beam, an electron beam current of 5 nA and a resolution better than 0.2 /~m whilst SEM results were obtained using a 9 keV electron beam and a resolution better than 0.5 /xm. SAM measurement provided information on the spatial uniformity of Sb distribution as well as on the absence of Sb islands on the surface in all the studied samples. The SEM micrography of the 10% Sb-doped ceramic sample
A. Gulino et al. / Applied Surface Science 90 (1995) 289-295
291
(Fig. 1) shows the presence of spherical, particles having sizes ranging from 2000-4000 A on the surface.
3. Results and discussion
3.1. Angle resolved measurements Normal-emission core level photoelectron spectra of Sb-doped TiO 2 ceramic samples, measured in a binding energy region containing both O ls and Sb 3d peaks are shown in Fig. 2. The Sb 3d5/2 peak overlaps the O ls peak almost perfectly, so the Sb 3d3/2 peak only can be used in quantification of the surface Sb content. The intensity ratio between Sb 3d3/2 and Ti 2p peaks is therefore used to derive values for the effective surface Sb/Ti ratio, taking into account the relevant atomic sensitivity factors [12]. It is reasonable to assume that analyser transmission functions for Sb3d and Ti 2p peaks are the same because they have similar kinetic energies. Values obtained in this way, for different emission angles relative to the surface normal, are given in Table 1. It becomes evident that the intensity of the Sb 3d peaks is very much greater than it would be expected on the basis of the nominal doping levels and that the Sb/Ti surface ratios vary less than variations in the nominal doping levels. Moreover, the enhanced
520
I
I
I
l
I
525
550
535
540
545
550
Binding Energy / eV Fig. 2. AlKot excited XPS of (top) 4% Sb-doped TiO 2 and (bottom) 10% Sb-doped TiO 2 measured in the O l s and Sb3d binding energy region. The Sb3ds/2 peak overlaps the O Is peak but the Sb3d3/2 peak is clearly seen at about 540 eV binding energy. Structure due to satellite radiation has been subtracted from the spectra.
Sb intensity at grazing emission angles indicates different Sb concentration in the ionic layers probed by XPS. Consider a series of cation containing planes with interplanar separation d. Assuming that the top n layers are enriched in a dopant it is possible to write a general expression for the angular variation of the XPS intensity ratio l(Sb 3d)/l(Ti 2p) corrected for the relevant atomic sensitivity factors [13] I(Sb 3 d ) / l ( T i 2p) =(O+xH)/[(1-O)+(1-x)H],
(1)
were H = K/(1 - K) and K = e x p ( - n d / A s i n a ) . Table 1 Apparent atomic ratios from Sb3d and Ti2p core intensities in relevant AR-XPS
Fig. I. Scanning electron micrograph of the 10% Sb-doped TiO 2 ceramic sample.
Nominal doping
Experimental [Sb/Ti] a from AR-XPS
Surface Sb from
level %
90'
30 °
9°
3°
Eq. (1)
1.0 4.0 5.0 7.0 10.0
0.05 0.14 0.15 0.16 0.25
0.06 0.17 0.18 0.19 0.26
0.07 0.19 0.21 0.20 0.29
0.08 0.20 0.21 0.23 0.30
0 0 0 0 0
a Estimated uncertainties are _4-10%.
= 0.07 =0.16 =0.17 = 0.19 = 0.23
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Here A represents the photoelectron inelastic mean free path; x represents the nominal doping level; t~ is the photoelectron take-off angle relative to the surface plane; 0 is the Sb occupancy of the Sb-enriched layers. We take A = 12.5 A for both Sb3d and Ti2p core level peaks [12] and d = 3.24 (which corresponds to the separation between adjacent Ti-containing planes in the (110) direction [9]). The intensity ratio I(Sb3d)/l(Ti2p) pattern observed for the 4% Sb-doped TiO 2 at different emission angles is consistent with a 0 value = 0.16 if the Sb enrichment is confined to the four top-most ionic layers. Assuming n = 4 the increases of the S b / T i intensity ratio on going from 90 ° to 3° are in good agreement with the angular variation predicted by Eq. (1). By contrast, if we assume enrichment in the topmost ionic layer we obtain a higher occupancy of Sb which is inconsistent with the observed intensity ratio at grazing emission angles. Here, using the S b / T i = 0.14 value, obtained for the 90 ° emission peak and assuming n = 1, Eq. (1) predicts a 0 value of 0.40. At the most grazing measured emission angle (3 °) this value implies an S b / T i ratio of 0.66. This calculated intensity value is totally inconsistent with the observed value of 0.20. Similar results have been obtained for the 5% Sb-doped sample (Table 1). More muted variation in intensity with take-off angle is found for the 10% Sb-doped sample. In fact, little variation of the S b / T i intensity ratio is observed on going from 90 ° to 30° emission peaks and, even at 3° grazing emission angle, the intensity increase is very much less pronounced, than for the 4% doped system. In this case, the calculated value of 0 = 0.23 is consistent with an enrichment involving five ionic planes. This picture is consistent with Transmission Electron Microscope (TEM) results which show an Sb-enriched layer approximately 15 ~, thick [9]. These results are different to those previously reported for Sb-doped SnO 2, another stable oxide which adopts the rutile structure [2]. Here Egdell and coworkers rationalized the high XPS Sb intensity in terms of replacement of Sn by Sb in the topmost ionic layer of the sample with full saturation of surface cation sites at 3% Sb bulk doping [14-17]. Obviously one limitation of the present approach is the use of a model which assumes a smooth (110) crystal surface to analyse experimental data obtained from a polycrystalline ceramic material.
-1.2-
-1.6-
~ -2.0.
-2.4-
-5.5
,
,
,
-5.0
-4.5
.-4.0
, - .5
; - .0
,
,
-2.5
-2.0
log([Sb]/[Wi]) Fig. 3. SurfaceSb3d to Ti2p intensityratio plotted as a function of bulk doping level.
An experimental estimate of the segregation enthalpy of Sb in Sb-doped TiO 2 rutile may be obtained by using the following expression [18]: ( S b / T i ) s = (Sb/Ti)b e x p ( - AHsegr/RT ).
(2)
A plot of the natural logarithms of these ratios, in the present doping range, is given in Fig. 3. The cleaning procedure of the samples involved prolonged annealing at 1000°C followed by rapid cooling to room temperature. Hence we assume that in XPS the equilibrium surface obtained at 1000°C is being probed. The value obtained for the heat of segregation is - 2 . 3 8 0 kJ mol -J. This value is less negative than the corresponding - 2 9 . 4 kJ tool-~ calculated for the Sb-doped SnO z parent system [14]. Of course, the heat of segregation depends on the bulk atomic ratio, so the value obtained is not surprising since two of the present Sb-doped TiO 2 samples (7%, 10%) exceed the solubility limit of antimony in TiO 2 rutile [9,11].
3.2. Ar + ion bombardment of Sb-doped TiO 2 It has been shown that oxygen vacancies are the dominant type of defect formed under ion bombardment of oxides [19,20]. This implies that such bombardment preferentially reduces oxide surfaces. In particular the (110) surface of TiO 2 reduces under a few hundred eV Ar + ion bombardment to a stoichiometry close to that of Ti203 [5-8]. The present
A. Gulino et al. /Applied Surface Science 90 (1995) 289-295
293
Ti f
~u ~
40.
20I
I
I
450
460 Binding energy / eV
470 0
Fig. 4. AIKo~ excited XPS spectra of Ti2p core level for 4% Sb-doped TiO 2 before (bottom) and after (top) sputtering.
10
20
30
40
50
60
Sputter time (min) Fig. 6. Depth profile of 4% Sb-doped TiO 2.
XPS spectra of Ti 2p core levels, obtained after sputtering, show shoulders on the low-binding energy side of the doublet evident for the TiO 2 itself (Fig. 4). These shoulders are due to the presence of Ti 3+ ions on the slightly reduced TiO 2 surfaces [5-8] which assume a pale grey colouration [9]. Finally, steady-state S b / T i compositions close to the ICPES values have been attained after long sputtering of all the analyzed systems showing that under our conditions the preferential sputtering of Ti or Sb is not a major problem as in the case of tin-antimony mixed oxide [21]. Figs. 5 and 6 show depth profiles of 10% and 4% Sb-doped TiO 2 respectively. Ti
80-
60O
40-
< 20-
Sb ' 0
i 1
,
i 2
,
i 3
•
i 4
'
i 5
,
i 6
,
i 7
,
l
•
S
Sputter time (rain) Fig. 5. Depth profile of 10% Sb-doped TiO 2.
i 9
'
i 10
It may be seen that surface compositions differ noticeably from the bulk values. All the samples exhibit the following features: the surface is enriched with Sb; the greater the nominal doping level the greater the surface S b / T i ratio; all the samples, at zero sputter time, show surface atomic compositions in tune with the 0 values calculated from the angleresolved measurements. The main difference between the differently doped samples is the thickness of the enriched layer. In fact, in the case of the 10% Sb-doped sample (Fig. 5) the Sb rapidly decreases with the increase of the sputter time and, after 2.5-3 min, the Sb percentage is close to the ICPES value. By contrast, the 4%-Sb-doped sample (Fig. 6) shows a different behaviour. Here the Sb content decreases slowly with the sputter time to reach the ICPES value after 50-60 min. These two extreme behaviours are obviously related to different thickness of the Sb-enriched layers in the two samples and suggest that it becomes possible to distinguish two limiting models for Sb-doped TiO 2 on the basis of the Sb doping level which, in turn, may be below or above the solubility limit of Sb in TiO 2. It has been seen in fact, by TEM measurements, that at doping levels higher than 5% the Sb-doped TiO 2 shows clear evidence of an amorphous S b - T i - O phase at surface and grain boundaries which, in the 10% doped sample, is approximately 15 ~, thick thus corresponding to 4 - 5 ionic layers [9]. Therefore, the rapid decrease of the Sb atomic percentage observed
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A. Gulino et al. / Applied Surface Science 90 (1995) 289-295
in the 10% Sb-doped T i t z rutile, upon Ar ÷ ion bombardment is due to the sputtering of this amorphous surface layer. Conversely, the behaviour of the 4% Sb-doped sample suggests that the number of ionic planes involved in the Sb-enrichment is underestimated by AR-XPS. In fact, 50-60 min sputtering time is needed to reach the ICPES S b / T i ratio thus suggesting that segregation involves a large number of ionic planes with the Sb percentage diminishing smoothly on going from the topmost ionic layer to the bulk. Obviously the present behaviour could be the results of a radiation-enhanced diffusion mechanism which depends on vacancies and other lattice defects generated by (heavy particle) irradiation [22] and could cause the Sb to move deeper into the substrate. Nevertheless, there is no explanation for the fact that this mechanism is operative for the 4% Sb-doped sample and not for the 10% doped sample, while identical operative conditions are always adopted. All these observations provide a sensible rationale for the monotonic decrease of the work function (as previously determined from the cut-off in secondary electron emission in UPS) throughout the doping range of samples presently studied [9]. As a matter of fact, Sb(III) species, having a formal 5s 2 electron configuration, are accommodated at the surface [9]. In these conditions, the non-centrosymmetric crystal eV 6
\
4
r
0
2
i
I
i
i
f
4 6 8 Doping Level / %
10
Fig. 7. Variationin work functionq~ with nominaldoping levelin Sb-doped Tit 2 ceramics as determined from the cut-off in secondary electronemission in UPS.
field at the surface sites leads to 5s-5p mixing. The 5s-5p hybrids have electron densities pointing out of the surface plane, but the surface dipole associated with these electrons is expected to be less than when bridging oxygen bonds to two Ti(IV) ions [9]. Thus the reduced surface dipole lowers the sample work function. The present surface Sb concentrations in the topmost ionic layer of the differently doped samples (0) correlate well with the monotonic decrease of the work function (Fig. 7). Finally the decrease in the intensity of the bandgap emission at 1.8 eV binding energy in the 7% and 10% Sb-doped samples (already observed with UPS measurements) [9] can be safely related to the different geometric structure of the second amorphous S b - T i - O phase present in the surfaces and grain boundaries. These states are now hidden under the more intense O2p based valence band as has been previously observed for Sb-doped SnO 2 [ 15,17].
4. Conclusion The combined application of angle-resolved XPS and depth profile measurements provides a powerful tool of probing surface composition of doped metal oxides. The present study highlights a pronounced surface segregation of Sb to polycrystalline doped T i t z rutile sample. In all the present samples the calculated Sb concentrations in the enriched topmost planes are in tune with Sb percentages obtained at zero sputter time. In the case of 4 - 5 % Sb-doped sample, Sb segregation involves a large number of planes with the Sb concentration varying smoothly on going from the top most ionic layer to the bulk. By contrast, in the case of heavily Sb-doped sample (7% 10%), in which the Sb exceeds the substitutional solubility limit, a new Sb-enriched S b - T i - O amorphous surface phase emerges whose thickness involves about five ionic planes with a rapid return of the bulk [Sb/Ti] ratio. As far as the influence of the annealing atmosphere on surface composition is concerned, we believe that the oxygen surface deficiency, due to in-vacuo annealing [9], favours the increase of the Sb surface concentration since, as found in an earlier study [9], the [D]-Sb(III)20 2] stable (110) termination is required to maintain the electrical neutrality.
A. Gulino et al. / Applied Surface Science 90 (1995) 289-295
Here [] indicates the removal of bridging oxygen atoms. This observation is tuned well with the fact that the 1 eV state, associated with Ti 3÷, is replaced by a bandgap structure much closer to the valenceband edge (1.8 eV) on passing from undoped TiO 2 to Sb-doped TiO 2 [9]. The lower Sb surface concentration found in the present air annealed sample is therefore in tune with the absence of oxygen deficiency. One of the major goals of the present investigation is the rationalization of the monotonic decrease of the work function with the nominal doping level in Sb-doped TiO 2 ceramics.
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[6] R. Sadeghi and V.E. Henrich, J. Catal. 109 (1988) I. [7] H. Mohamed, H.R. Sadeghi and V.E. Henrich, Phys. Rev. B 37 (1988) 8417. [8] J.L. Sullivan, S.O. Saied and I. Bertoti, Vacuum 42 (1991) 1203. [9] A. Gulino, A.E. Taverner, S. Warren, P. Harris and R.G. Egdell, Surf. Sci. 315 (1994) 351. [10] M. Repoux, Surf. Interf. Anal. 18 (1992) 567. [1 l] N. Morita, T. Endo, T. Sato and M. Shimada, J. Solid State Chem. 68 (1987) 106. [12] D. Briggs and M.P. Seah, Eds., Practical Surface Analysis (Wiley, Chichester, 1983). [13] M.P. Seah, Surf. lnterf. Anal. 2 (1980) 222. [14] R.G. EgdeU, W.R. Flavell and P. Tavemer, J. Solid State Chem. 51 0984) 345. [15] P.A. Cox, R.G. Egdell, C. Harding, W.R. Patterson and P. Tavemer, Surf. Sci. 123 0982) 179. [16] P.A. Cox, R.G. Egdell, C. Harding, A.F. Orchard, W.R. Patterson and P. Tavemer, Solid State Commun. 44 (1982) 837. [17] R.G. Egdell, in: Science of Ceramic Interfaces 1I, Ed. J. Nowotny (Elsevier, Amsterdam, 1994) pp. 527-565. [18] L. Cao, R.G. Egdell, W.R. Flavell, K.F. Mok and W.C. Mackrodt, J. Mater. Chem. (1991) 785. [19] H.M. Naguib and R. Kelly, Rad. Eft. 25 (1975) I. [20] J.B. Malherbe, S. Hofmann and J.M. Sanz, Appl. Surf. Sci. 27 (1986) 355. [21] Y. Cross and D.R. Pyke, J. Catal. 58 (1979) 61. [22] A. Benninghoven, F.G. Riidenauer and H.W. Wemer, Secondary Ion Mass Spectrometry, Basic Concepts, Instrumental Aspects, Applications and Trends (Wiley, New York, 1987).