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Structure and composition of the titanium oxide layers formed by low-pressure oxidation of the Ni94Ti6(110) surface
surface science ELSEVIER
Surface Science 391 (1997) 216 225
Structure and composition of the titanium oxide layers formed by low-pressure oxidation of the Ni94Ti6( 1 10) surface A. Atrei a.,, U. Bardi b, G. Rovida b a
Dipartamento di Scienze e Tecnologie Chimice e dei Biosistemi. Universith di Siena, Siena, Italy b Dipartimento di Chimica, Universith di Firenze, Firenze, Italy
Received 1 April 1997; accepted for publication 6 June 1997
Abstract
The structure and composition of the titanmm oxide phases obtained by low-pressure oxidation at ca 800 K of the Ni~4Ti~,(110) surface have been investigated by a combination of surface sensitive techniques. We found that upon oxygen exposure in these conditions layers of oxide consisting mainly of TiO 2 grow on the alloy surface. The TiO2 phase formed in the very early stages of oxidation has a quasi-hexagonal unit cell, as shown by low-energy electron diffraction, and its structure has no equivalent in any of the stable phases of TiO> The structure of this ultrathin oxide film consists of a layer of titanium atoms between two atomic phmes of oxygen atoms. Upon increasing the oxygen exposure, islands of epitaxial futile (110) form on the surface, as shown by the X-ray photoelectron diffraction data. ~-~ 1997 Elsevier Science B.V. Ke)'words: Alloys; Epitaxy; Low-energy electron diffraction; Low-energy ion scattering: Nickel: Oxidation; Surface structure; Titanium; Titanium oxide; X-ray photoelectron diffraction; X-ray photoelectron spectroscopy'
1. Introduction
The o x i d a t i o n o f nickel t i t a n i u m alloys is relevant b o t h from the p o i n t o f view o f the functional p r o p e r t i e s o f the substrate as well as o f the oxide layer. The n a t u r e o f the oxide films which form in the early stages o f o x i d a t i o n determines the resistance o f the alloy to c o r r o s i o n in oxidizing environments and the structure o f the oxide alloy interface is i m p o r t a n t in o r d e r to design a ceramic c o a t i n g o f the m a t e r i a l in structural a p p l i c a t i o n s [1]. F u r t h e r m o r e , the o x i d a t i o n processes in the fabrication o f shape m e m o r y alloy NiTi thin films m a y alter the c o m p o s i t i o n o f the alloy and render * Corresponding author. Fax: {+ 39) 55 219802: e-maih [email protected] 0039-6028/97/$17.00 ~:~1997 Elsevier Science B.V. All rights reserved. PII S0039-6028 (97)00485-8
difficult to c o n t r o l the final c o m p o s i t i o n o f the deposited films, a crucial p a r a m e t e r in d e t e r m i n i n g the t r a n s f o r m a t i o n t e m p e r a t u r e [2]. F r o m the p o i n t o f view o f the oxide o v e r l a y e r properties, it has been shown that t i t a n i u m suboxides d e p o s i t e d on nickel surfaces enhance the activity o f the metal for C O h y d r o g e n a t i o n [3]. The thin oxide films p r o d u c e d by the o x i d a t i o n o f this N i T i alloy m a y have a structure ( b o t h a t o m i c a n d electronic) and reactivity different from those o f bulk t i t a n i u m oxide phases. A few studies by means o f A u g e r electron s p e c t r o s c o p y ( A E S ) and ultraviolet p h o t o e l e c t r o n s p e c t r o s c o p y ( U P S ) a b o u t the oxid a t i o n a n d reactivity o f N i T i alloys have been r e p o r t e d in the literature [4,5]. However, since polycrystalline samples were investigated no inform a t i o n a b o u t the structure o f the oxide layers
A. A trei et al. / Su
could be derived. In the present work, we used a combination of surface science techniques to investigate the composition and structure of the oxide layers which form on the (110) surface of the Ni94Ti 6 alloy upon exposure to oxygen at a temperature of ca 800 K in the 10 7 Torr range. N i T i is a quite complex system and its phase diagram shows the existence of several intermetallic phases [6]. The Ni94Ti6 alloy is a solid solution of titanium atoms in the nickel lattice. We chose to study such a diluted Ti alloy in order to be sure to have a single phase since separation of Ni-Ti intermetallic compounds occurs for a Ti concentration higher than ca 10at% [7]. The (110) orientation of the crystal has been chosen to determine the effect of such anisotropic and corrugated substrate on the epitaxial growth of the oxide films.
2. Experimental The NiTi alloy single crystal was prepared as described in Ref. [7]. The crystal was oriented and cut along the (110) plane with an accuracy of +0.5 ° . The sample was then polished in order to get a mirror-like surface. The Ti concentration was found to be 6_+ 0.1 at% by means of X-ray fluorescence and X-ray diffraction [7]. The lattice parameter for this composition was 3.548 A [7]. The experiments were performed in a ultra-high vacuum ( U H V ) chamber with a base pressure of I ×10-1°Torr. The sample was mounted on a manipulator with two rotation axes. The chamber was equipped with hemispherical analyser, used for low-energy ion scattering ( LEIS ), X-ray photoelectron spectroscopy (XPS) and X-ray photoelectron diffraction (XPD) measurements. The A1 K~ radiation ( 1486.6 eV ) was used for the XPS and XPD experiments. The angle between the X-ray source and the analyser was 55 ° and the semicone acceptance angle of the analyser was 5 t The LEIS spectra were obtained using 1 keV He ion impinging on the sample surface with an angle of 45 ° and a scattering angle of 135 °. The XPS binding energy scale was calibrated on the position of the Ag 3d5/2 photoemission peak (368.2 eV) [8] and was accurate within __0.1 eV. The XPS spectra were measured at an emission angle of 60 ° with respect to the surface normal. The atomic sensitiv-
217
ity factors for the O l s and Ti 2p photoemission peaks, used in the quantification of the XPS data, were derived from the spectra of a bulk TiO 2 sample. The XPS spectra were smoothed and the background subtracted according to the Shirley method [9]. The Ti 2p XPS spectra were analysed by means of a curve-fitting procedure. The XPD measurements were performed by monitoring the intensity of the Ti 2p3/2 (kinetic energy 1025eV), Ni2p3/2 (633eV) and O l s (948 eV) peaks as a function of the polar and azimuthal emission angles. The polar angle, 0, is defined with respect to the surface normal and the azimuthal angle, q~, with respect to the [001] direction of the substrate. The anisotropy in the XPD curves is defined as (/max--Imin)/Imax %, where the Imi, and /max are the minimum and maximum intensity values in the XPD curves. The twodimensional patterns representing the XPD intensities were obtained from polar scans measured every 5 ~' of azimuthal angle and covering 180 °. The full 27r plots were built by repeating the measured data exploiting the two-fold symmetry of the surface. The XPD results were interpreted on the basis of the single scattering cluster-spherical wave (SSC-SW) theory [10]. The crystal structure was described as a cluster whose dimension was limited setting a maximum electron path from the emitting atom to the scatterer atoms of 8 A. Larger values of this parameter did not produce any significant change in the calculated XPD curves. The phase shifts for titanium and oxygen were derived from the muffin-potential calculated for TiO. The surface was cleaned by means of cycles of argon ion bombardment (2 keV) and annealing up to 850 K. The surface of this N i T i alloy turns out to be very reactive towards the residual atmosphere in the vacuum chamber and a small quantity of oxygen was always observed in the LEIS spectra.
3. Results 3.1. X P S and L E I S results
The XPS spectra for the clean surface of the alloy show that the position of Ti 2p3.2 is shifted
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by 0.4eV towards higher binding energy with respect to the value of the pure metal (454.0 eV) [8]. No chemical shift, in the limit of the accuracy of the measurements, could be observed for the Ni 2p3/2 peak in the alloy compared with pure Ni. The XPS and LEIS data indicate that there is no preferential sputtering of either components and a stable surface composition was reached after cycles of ion sputtering and annealing at 850 K. The oxidation procedure of the alloy surface consisted in exposure to an oxygen pressure in the 10 7 Torr range at 780 K and subsequent cooling (down to 500 K ) in oxygen atmosphere. The Ni 2p XPS spectra show that nickel was not oxidized in these conditions. The Ti 2p XPS spectra measured for various oxygen exposures are shown in Fig. 1.
r-
d
_c
C
b
470
465
460
455
450
Binding energy (eV) Fig. 1. Ti 2p XPS spectra measured after exposures to an oxygen pressure in the 10 v torr range of the alloy surface at 780 K. (a) Clean surface; (b) 30 L plus 30 L during cooling; (c) 60 L plus 30 L during cooling; (d) 150 L plus 100 L during cooling down to ca 500 K. For the last spectrum the components corresponding to Ti ~v (binding energies Ti 2p3/2 and Ti 2pi/2 459.0 and 464.6 eV, respectively) and to Ti m (binding energies Ti 2p3/2 and Ti 2pl/2 456.9 and 461.4, respectively), determined from the curve-fitting analysis, are shown as dashed lines.
Already for relatively low oxygen exposures, the oxidation of the alloy leads to a relevant increase of the surface concentration of titanium (in the form of oxide), as indicated by the growth of Ti 2p peaks at higher binding energies compared with the clean surface. The binding energy of the Ti 2p3/2 peak in the oxide film formed after an oxygen exposure of 30 Langmuir (L) (1 L a n g m u i r = l x l 0 6 T o r r . s ) at high temperature plus 30 L during cooling was 458.6 eV; this value is lower than that in bulk TiO 2 (459.0 eV) but coincident with the Ti 2p3/2 binding energies measured for the TiO 2 layers formed upon oxidation of Pt3Ti alloy surfaces or deposited on Pt surfaces [11]. The O ls binding energy was 530.2 eV. The oxide layer formed under these conditions covers completely the surface of the alloy, since no Ni signal is observed in the LEIS spectra (Fig. 2). The thickness of the oxide layers was estimated from the attenuation of the Ni 2p3/2 signal produced by the oxide film covering surface using the attenuation length for the Ni 2p3/2 photoelectrons calculated according to the Seah and Dench method [12]. From the XPS data, we found that the oxide layer was ca 3 A thick. The Ti/O atomic ratio determined for this phase turns out to be 1:2.1_+0.1. The T i 2 p spectrum measured after an oxygen exposure of 60 L at 780 K and plus 30 L during cooling is shown in Fig. 1, curve c. A shoulder appears on the low binding energy side of the Ti 2p3/2 peak while the main peak is shifted towards higher binding energy with respect to spectrum in Fig. 1, curve b. The curve fitting analysis showed that the Ti 2p3/2 peak consisted of two components at binding energies of 459.0 and 456.6eV. The former component corresponds to TiO2, whereas the latter is typical of titanium suboxides, TiOx, with x ranging from 1.0 to 1.5 [11]. Beyond an exposure of ca 100 L and up to the maximum oxygen dose of 600 L used in the present work, no changes in the Ti 2p XPS spectra were observed and the oxidation process reached a saturation after an exposure of 150 L. The thickness of the titanium oxide film formed under these conditions was estimated to be i l k . We found that the amount of reduced titanium oxide with respect to the TiO2 quantity depends on the exposure to oxygen during cooling, being
A. Atrei et al. ,/Surjitce Science 391 (1997) 216 225
was 530.5 eV. The annealing at ca 700 K for a few minutes of the oxide films prepared under the conditions described above, produced an increase in the amount of titanium suboxides TiOx coming from the reaction between TiO 2 and metallic titanium diffusing from the bulk of the alloy.
He+ 1keV
/
219
Ni
3.2. L E E D observations Ti
O
.__>, e-
t-
.,A,__
.L ,
0.3
I
0.4
I
.,,%__ I
0.5
m
I
,
0.6
I
0.7
,
I
0.8
,
I
0.9
E/E o Fig. 2. LEIS spectra measured for the clean surface (curve a); after 30 L of oxygen at 780 K (curve b); after 150 L of oxygen at 780 K (curve c).
larger for shorter oxygen exposures. From the curve fitting analysis of the Ti 2p XPS spectrum for the oxide film formed after an oxygen exposure of 150L at 7 8 0 K plus 100L during cooling (Fig. 1, curve d) it turns out that in the Ti 2p3/2 there is the component corresponding to Ti TM, located at 459.0eV, and a component at lower binding energy (456.9 eV) due to titanium in lower oxidation states. The Ti 2p3/2 binding energy of this suboxide is close to the values measured for TiOx thin layers formed upon oxidation of titanium deposits on Ni (457.0 eV) [13] and Pt surfaces (456.2 eV) [14], and upon oxidation of the Pt3Ti alloy (456.8) [15]. In all these cases this component was attributed to Ti m, though the values of the binding energy are in between those characteristic of bulk TiO and Ti203 [16]. The O ls binding energy measured for this oxide layer
The clean surface of the alloy showed a (1 x 1) low-energy electron diffraction (LEED) pattern. Depending on the oxidation conditions different LEED patterns could be observed. The oxide phase formed at 7 8 0 K after 3 0 L of oxygen showed the LEED pattern reported in Fig. 3a. At first glance, taking into account the most intense spots, the superstructure seems to have a hexagonal unit cell. However, a closer inspection of the pattern reveals the existence of a c(2 x 6) coincidence mesh between the unit cell of the substrate and that one of the overlayer. Hence, the unit cell of the oxide overlayer cannot be hexagonal since it cannot fit the rectangular periodicity of the substrate surface. The surface unit cell of this oxide phase can be determined exactly from the c(2 x 6) coincidence net. The matrix notation for this surface unit cell is: -3 1
0
7 6 7
The sides of the unit cell are 2.93 ,~ and 3.04 and the angle between them is 121.2 ° . Since the surface unit cell of this oxide film is very close to a hexagonal mesh, we call it a "quasi-hexagonal" unit cell. Upon increasing the oxygen exposure, the quasi-hexagonal superstructure gradually becomes weaker, disappears and eventually the pattern shown in Fig. 3b can be observed. This pattern is characterized by extraspots along the [110] direction and strikes along the [001] direction. Only the first order extra spots are visible in the diffraction diagram and the distance between these reflexes and the integral order spots is ca one-fifth of the side of substrate unit cell along the
220
A. Atrei et al. / SurJace Science 391 ( 1997j 216 225
(a)
(b) Fig. 3. (a) LEED pattern observed after oxygen exposure of 30 L at 780 K of the alloy surface. Normal incidence of the electron beam, energy 90 eV. (b) LEED pattern produced by the oxide layers obtained after an oxygen exposure of 150 L at 780 K plus 100 L during cooling. Normal incidence of the electron beam, energy 68 eV.
A. Atrei et al. / Sur[itce Science 391 (1997) 216 225
[ 110] direction. These extra spots may be attributed to double diffraction through the periodicity of the substrate and of the overlayer, with the higher order spots being too weak to be observed. Along the [001] the TiO2 phase is disordered as shown by the presence of diffuse intensity. 3.3. X P D results
The XPD patterns for the Ti 2p3/2 and O l s photoemission peaks of the quasi-hexagonal TiO 2
Ols
phase are shown in Fig. 4. Although, the unit cell of the oxide film as determined from the LEED pattern is oblique, the XPD patterns appear sixfold symmetric because the deviations of the cell from a hexagonal mesh are so small than cannot be detected in our experimental conditions. The lack of features at polar emission angles below 40 50 ° is an indication that the oxide layer has a thickness of a few atomic layers. In order to interpret the XPD data we performed SSC calculations on the basis of a structural model having the quasi-hexagonal unit cell determined from the LEED pattern. The dimensions of the overlayer unit cell are such that it is not possible to fit in more than one atom (either Ti or O). Hence, to reproduce the TiO2 stoichiometry we considered a layer of titanium atoms sandwiched between two oxygen planes (Fig. 5). The titanium
I [] i0l
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Ti 2p3/2
221
O atoms (2nd layer)
e
Ti atoms ° .
•
,
,
,
j
Substrate atoms
tu01]
Min.
Max.
Fig. 4. Stereographic projections of the O ls and Ti 2p3, 2 X P D curves measured for the quasi-hexagonal TiOz phase. The intensities are represented as grey-scale levels. Along the radial directions, from the centre to the circle, the polar emission angle varies from normal emission 0 = 0 ' (normal emission) to 0 = 90'. The azimuthal directions of the substrate are reported. The region near normal emission was not measured for the Ti 2p photoemission peak since, after collecting some polar scan, we found that the curves were featureless for 0 below 40 5 0 .
[001]
f
,
4 .
.k.
[ [ITol
Fig. 5. Schematic model for the quasi-hexagonal TiO 2 layer. Both the quasi-hexagonal unit cell of the overlayer and the c(2 x 6) coincidence mesh are shown. In this figure also the substrate atoms are drawn to show the epitaxial relationship between the overlayer and the substrate.
222
A. Atrei et al.
Sur{~tce Science 391 (1997) 216 225
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atoms are located in the centre of the triangles delimited by the nearest oxygen atoms. This model for the TiO2 overlayer structure has a lower symmetry than the substrate and domains, produced by the symmetry operations of the substrate structure, must exist in order to have an overall twofold symmetry. Therefore, in the calculations we considered in addition to the structure shown in Fig. 5 its mirror domain (obtained by reflection through the vertical plane containing the [001] direction). On the basis of this model a good level of agreement with the experimental data was obtained (Fig. 6). The best fit corresponds to an interlayer spacing of 1.1 _+0.1 A. This value is very close to that (1.0A,) calculated assuming the average T i - O nearest neighbour distance in bulk rutile (1.96 A) [17]. The residual disagreement between experimental and calculated curves might be due to a buckling of the oxide film. Because of the large number of non-equivalent atoms in the unit cell and the possible registries of the oxide layer with respect to the substrate, no attempt has been done to introduce in the calculations a rippling of the quasi-hexagonal layer. The structural changes of the titanium oxide layer upon further oxidation of the alloy at high temperature are clearly revealed by XPD. The X P D patterns measured for the Ti 2p3/2 and O l s photoemission peaks of the TiO2 film formed after an oxygen exposure of 600 L at 780 K are shown in Fig. 7. The modulations observed in the XPD patterns indicate that the thicker TiO 2 film grows epitaxially on the alloy surface. The Ti2p3:2 pattern is characterized by intense peaks at polar angles 0 = 70" along the azimuth corresponding to the [001] direction of the substrate (q~=0) and 0 = 5 4 'J along the azimuth corresponding to the [ 11-0] direction of the substrate (~b=90). In the O l s pattern the most intense features are located at 0 = 7 0 ° in the ~b=0 ° azimuth and 0 = 4 5 in the q~= 90 ~azimuth. The intensity enhancement of the
Fig. 6. Comparison between experimental(thick line) and calculated azimuthal XPD curves measured the polar angles 0= 70'and 0=60 for the quasi-hexagonal TiO2 phase. (a) O ls peak, (b) Ti 2p3,2 peak. For each curve the maximum value of the anisotropy defined as A1/lm,x%, where Al=Imax--lm~n is reported.
223
A. Atrei et al. / Sur/ace Science 391 (1997) 216-225
O ls
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AI/Imax%
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.
33 135 180 i
,
,
,
Azimuthal angle (degrees) (a)
Ti2P3/2 Min.
AI/Irnax%
Max.
Fig. 7. Stereographic projection of the O ls and Ti 2p3/2 XPD intensities measured for the TiO2 phase obtained after oxidation in the conditions described in the text. The azimuthal directions of the substrate are reported.
70° ~
27
60°
27
¢R t-
_=
diffraction features at higher polar angles is due to the relatively low thickness of the oxide layer [18]. The models for the SSC simulations of the experimental XPD curves were based on the structure of bulk T i O 2. T i O 2 exists in three polymorphic Fig. 8. Comparison between experimental (thick line) and calculated azimuthal XPD curves measured at various polar angles for the thicker TiO2 phase. (a) O ls peak. (b) Ti 2p3:2 peak. The calculations were performed for epitaxial islands having the futile (110) structure oriented, with respect to the NiTi alloy surface, in such a way that the short side of the oxide overlayer unit cell is parallel to the direction [1 TO] of the substrate.
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A. Atrei et aL / Suffdce Science 391 (1997) 216 225
224
crystalline structures: rutile, anatase and brookite [19]. We started the model calculations with the (110) surface of futile since this crystallographic plane is the most stable termination of futile. We tried different topmost layer terminations of the futile (110) islands and the interlayer spacing as well as the in-plane parameters were varied in the calculations. As shown in Fig. 8 a fairly good level of agreement between experimental and calculated XPD curves is obtained for the (110) bulk truncated structure of rutile oriented in such a way that the short side of the oxide unit cell is parallel to the [110] direction of the substrate. The model is also consistent with the observed LEED pattern for this phase, since every 5 unit cells of the substrate in [110] direction there is a coincidence with the periodicity of the rutile (110) surface, assuming a 6% expansion of the oxide overlayer. We found that the XPD calculations were sensitive to the termination of the surface and the best agreement was obtained for the termination of TiO2(ll0) (Fig. 9) proposed in Ref.[20] and recently confirmed by surface X-ray diffraction results [21]. However, as it may be expected, there is only a small effect of the variations of the topmost layer spacings on the calculated curves. In addition to the (110) orientation we tried also the (100) surface of rutile and the (100) surface of anatase. These surfaces were chosen because the rectangular unit cell of these oxides could fit the rectangular unit mesh of the substrate. None
Oxygen atoms
topmostlayer
Titanium atoms
O
0"° "" ""
2nd layer
bulk layer
Fig. 9. Schematic structure of the rutile(ll0) surface showing the oxygen-terminatedsurfacewith six-foldand five-foldcoordinated titanium atoms.
of these models gave a theory-experiment agreement comparable with that obtained using the rutile (110) surface.
4. Discussion
The XPS results for the oxidation at high temperature of the Ni94Tidl10) surface indicate that in the very early stages TiO2 islands grow whereas some amounts of titanium suboxides form upon further oxidation. This behaviour may be explained if the oxidation rate decreases upon increasing the thickness of the oxide layer. When the oxide phase is a few layers thick all the titanium diffusing from the bulk is oxidized to Ti TM. On the other hand, if the oxidation rate decreases when the oxide layer becomes thicker titanium atoms can react at the alloy oxide interface leading to the formation of reduced titanium oxides. This titanium suboxide contains Ti m, as revealed by the low binding energy shoulder of the Ti 2p3. 2 peak corresponding to Ti TM. The increase (ca 0.4 eV ) of the Ti 2p peaks for Ti TM and O ls binding energies observed when TiO~ is formed can be explained by a shift of the Fermi level (due to the partial occupation of the titanium oxide conduction band) towards the vacuum level. This shift results in an apparent increase of the core level binding energies, as suggested by Dwyer et al. for titanium oxide layers on platinum [22]. From the surface crystallographic point of view the most relevant result of this study is the demonstration of the possibility to grow a TiO 2 phase which has no equivalent in bulk crystals. In the early stages of oxidation of the alloy surface, a structure is formed (with a quasi-hexagonal unit cell) consisting of a layer of titanium atoms between two planes of oxygen atoms instead of the growth of oxide patches having one of the crystalline structures (perhaps slightly distorted) of bulk TiO 2. To our knowledge this is the first time that the structure of an epitaxial oxide layers produced by the oxidation of an alloy surface cannot be derived from that of some bulk oxide. For instance, in the early stages of oxidation of Ni3AI(001), Ni3AI( 111 ) [23] and NiAI(110) [24] a fraction of the surface was covered by islands of
A. Atrei et al. / Surlace Seience 391 (1997) 216 225
7-alumina having a thickness comparable to the length of unit cell side in the direction perpendicular to surface. Also the LEED pattern produced by the oxide layers formed upon oxidation of the (111) and (100) surfaces of the Pt3Ti alloy were interpreted in terms of the structure of closed packed planes of bulk titanium oxides [25]. The reason why in the earlier stages of oxidation the quasi-hexagonal TiO2 phase forms instead of islands having the structure of one of the bulk phases of TiO2 appears to be the corrugation of the substrate surface. It is conceivable that a closed packed structure like rutile (110) can hardly fit the corrugation of a fcc(110) surface (in particular in the [001] direction, perpendicular to the grooves). On the other hand, bulk TiO 2 phases should form upon oxidation of the more flat (100) and ( 111 ) surfaces of the NiTi alloy.
5. Conclusions We found that upon exposure to oxygen of the Ni94Ti6(110) surface at high temperature layers of TiO2 form at the surface. The LEED and XPD data showed that while relatively thick oxide films consist of epitaxial islands of rutile oriented along the (110) surface, the quasi-hexagonal TiO2 phase has its own structure with no equivalent in bulk titanium dioxide phases. The formation of this two-dimensional oxide phase may be related to the corrugation of the substrate which hinders the growth with the structure of one of the bulk TiO 2 crystalline phases.
Acknowledgements
This work was partially supported by MURST. The authors are grateful to Mr E. Fisher for the
225
preparation of the NiTi sample and to Professor G. Kostorz for the useful discussion.
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R.J. Lad, Surf. Rev. Lett. 2 (1995) 109. J.D. Bush, A.D. Johnson, J. Appl. Phys. 68 (1990) 6224. Y.W. Chung, G. Xiong, C.C. Kao, J. Catal. 85 (1984) 237. P.H. McBreen, M. Polak, Surf. Sci. 179 (1987) 483. A. Petri, A. Neumann, J. Kiippers, J. Vac. Sci Technol. A 8 (1990) 2576. [6] T.B. Massalski, Binary Alloy Phase Diagrams, ASM, Metals Park, OH, 1986. [7] T. Ungar, M. Lambrigger, G. Kostorz, Mater. Sci. Eng. A 138 (1991) 147. [8] J. Chastain (Ed.), Handbook of X-ray Photolectron Spectroscopy. Perkin Elmer, Eden Prairie, MN, 1992. [9] D. Briggs, M.P. Seah (Eds.), Practical Surface Analysis, Wiley, New York, 1983. [10] C.S. Fadley, in: R.Z. Bachrach (Ed.), Synchrotron Radiation Research: Advances in Surface in Surface Science. Plenum, New York, 1991. [11] U. Bardi, Catal. Lett. 5 (1990) 81. [12] M.P. Seah, W.A. Dench, Surf. Interface Anal. 1 (1979) 2. [13] J.P. Espinos, A. Fernandez, A.R. Gonzalez-Elipe, Surf. Sci. (1993) 402. [14] C.M. GreenlieL J.M. White, C.S. Ko, R.J. Gorte, J. Phys. Chem. 89 (1985) 5025. [15] A. Dausher, L. Hilaire, J.C. Spirlet, W. Muller, G. Maire. Surf. Sci. 204 (1978) 161. [16] A.F. Carley, P.R. Chalker, J.C. Riviere, M.W. Roberts. J. Chem. Soc. Faraday Trans. 83 (1987) 351. [17] R.W.G. Wyckoff, Crystal Structures, Wiley, New York. 1964. [18] C.S. Fadley, Prog. Surf. Sci. 16 (1984) 275. [19] A.F. Wells. Structural Inorganic Chemistry, Clarendon Press, Oxford, 1985. [20] V.E. Henrich, R.L. Kurtz. Phys. Rev. B 23 (1981) 6280. [21] G. Charlton, P.B. Howes, C.L. Nicklin, P. Steadman, J.S.G. Taylor, C.A. Muryn, S.P. Harte, J. Mercer, R. Mc Grath, D. Norman, T.S. Turner, G. Thorntom Phys. Rev. Len. 78 (1997)495. [22] D.J. Dwyer, S.D. Cameron, J. Gland, Surf. Sci. ( 1985 ) 430. [23] U. Bardi, A. Atrei, G. Rovida, Surf. Sci. 268 (1992) 87. [24] R.M. Jaeger, H. Kuhlenbeck, H.-J. Freund. M. Wuttig, W. Hoffmann, R. Franchy, H. lbach, Surf. Sci. 259 ( 1991 ) 235. [25] U. Bardi, P.N. Ross, J. Vac. Sci. Technol. 2 (1984) 1461.
Surface Science 391 (1997) 216 225
Structure and composition of the titanium oxide layers formed by low-pressure oxidation of the Ni94Ti6( 1 10) surface A. Atrei a.,, U. Bardi b, G. Rovida b a
Dipartamento di Scienze e Tecnologie Chimice e dei Biosistemi. Universith di Siena, Siena, Italy b Dipartimento di Chimica, Universith di Firenze, Firenze, Italy
Received 1 April 1997; accepted for publication 6 June 1997
Abstract
The structure and composition of the titanmm oxide phases obtained by low-pressure oxidation at ca 800 K of the Ni~4Ti~,(110) surface have been investigated by a combination of surface sensitive techniques. We found that upon oxygen exposure in these conditions layers of oxide consisting mainly of TiO 2 grow on the alloy surface. The TiO2 phase formed in the very early stages of oxidation has a quasi-hexagonal unit cell, as shown by low-energy electron diffraction, and its structure has no equivalent in any of the stable phases of TiO> The structure of this ultrathin oxide film consists of a layer of titanium atoms between two atomic phmes of oxygen atoms. Upon increasing the oxygen exposure, islands of epitaxial futile (110) form on the surface, as shown by the X-ray photoelectron diffraction data. ~-~ 1997 Elsevier Science B.V. Ke)'words: Alloys; Epitaxy; Low-energy electron diffraction; Low-energy ion scattering: Nickel: Oxidation; Surface structure; Titanium; Titanium oxide; X-ray photoelectron diffraction; X-ray photoelectron spectroscopy'
1. Introduction
The o x i d a t i o n o f nickel t i t a n i u m alloys is relevant b o t h from the p o i n t o f view o f the functional p r o p e r t i e s o f the substrate as well as o f the oxide layer. The n a t u r e o f the oxide films which form in the early stages o f o x i d a t i o n determines the resistance o f the alloy to c o r r o s i o n in oxidizing environments and the structure o f the oxide alloy interface is i m p o r t a n t in o r d e r to design a ceramic c o a t i n g o f the m a t e r i a l in structural a p p l i c a t i o n s [1]. F u r t h e r m o r e , the o x i d a t i o n processes in the fabrication o f shape m e m o r y alloy NiTi thin films m a y alter the c o m p o s i t i o n o f the alloy and render * Corresponding author. Fax: {+ 39) 55 219802: e-maih [email protected] 0039-6028/97/$17.00 ~:~1997 Elsevier Science B.V. All rights reserved. PII S0039-6028 (97)00485-8
difficult to c o n t r o l the final c o m p o s i t i o n o f the deposited films, a crucial p a r a m e t e r in d e t e r m i n i n g the t r a n s f o r m a t i o n t e m p e r a t u r e [2]. F r o m the p o i n t o f view o f the oxide o v e r l a y e r properties, it has been shown that t i t a n i u m suboxides d e p o s i t e d on nickel surfaces enhance the activity o f the metal for C O h y d r o g e n a t i o n [3]. The thin oxide films p r o d u c e d by the o x i d a t i o n o f this N i T i alloy m a y have a structure ( b o t h a t o m i c a n d electronic) and reactivity different from those o f bulk t i t a n i u m oxide phases. A few studies by means o f A u g e r electron s p e c t r o s c o p y ( A E S ) and ultraviolet p h o t o e l e c t r o n s p e c t r o s c o p y ( U P S ) a b o u t the oxid a t i o n a n d reactivity o f N i T i alloys have been r e p o r t e d in the literature [4,5]. However, since polycrystalline samples were investigated no inform a t i o n a b o u t the structure o f the oxide layers
A. A trei et al. / Su
could be derived. In the present work, we used a combination of surface science techniques to investigate the composition and structure of the oxide layers which form on the (110) surface of the Ni94Ti 6 alloy upon exposure to oxygen at a temperature of ca 800 K in the 10 7 Torr range. N i T i is a quite complex system and its phase diagram shows the existence of several intermetallic phases [6]. The Ni94Ti6 alloy is a solid solution of titanium atoms in the nickel lattice. We chose to study such a diluted Ti alloy in order to be sure to have a single phase since separation of Ni-Ti intermetallic compounds occurs for a Ti concentration higher than ca 10at% [7]. The (110) orientation of the crystal has been chosen to determine the effect of such anisotropic and corrugated substrate on the epitaxial growth of the oxide films.
2. Experimental The NiTi alloy single crystal was prepared as described in Ref. [7]. The crystal was oriented and cut along the (110) plane with an accuracy of +0.5 ° . The sample was then polished in order to get a mirror-like surface. The Ti concentration was found to be 6_+ 0.1 at% by means of X-ray fluorescence and X-ray diffraction [7]. The lattice parameter for this composition was 3.548 A [7]. The experiments were performed in a ultra-high vacuum ( U H V ) chamber with a base pressure of I ×10-1°Torr. The sample was mounted on a manipulator with two rotation axes. The chamber was equipped with hemispherical analyser, used for low-energy ion scattering ( LEIS ), X-ray photoelectron spectroscopy (XPS) and X-ray photoelectron diffraction (XPD) measurements. The A1 K~ radiation ( 1486.6 eV ) was used for the XPS and XPD experiments. The angle between the X-ray source and the analyser was 55 ° and the semicone acceptance angle of the analyser was 5 t The LEIS spectra were obtained using 1 keV He ion impinging on the sample surface with an angle of 45 ° and a scattering angle of 135 °. The XPS binding energy scale was calibrated on the position of the Ag 3d5/2 photoemission peak (368.2 eV) [8] and was accurate within __0.1 eV. The XPS spectra were measured at an emission angle of 60 ° with respect to the surface normal. The atomic sensitiv-
217
ity factors for the O l s and Ti 2p photoemission peaks, used in the quantification of the XPS data, were derived from the spectra of a bulk TiO 2 sample. The XPS spectra were smoothed and the background subtracted according to the Shirley method [9]. The Ti 2p XPS spectra were analysed by means of a curve-fitting procedure. The XPD measurements were performed by monitoring the intensity of the Ti 2p3/2 (kinetic energy 1025eV), Ni2p3/2 (633eV) and O l s (948 eV) peaks as a function of the polar and azimuthal emission angles. The polar angle, 0, is defined with respect to the surface normal and the azimuthal angle, q~, with respect to the [001] direction of the substrate. The anisotropy in the XPD curves is defined as (/max--Imin)/Imax %, where the Imi, and /max are the minimum and maximum intensity values in the XPD curves. The twodimensional patterns representing the XPD intensities were obtained from polar scans measured every 5 ~' of azimuthal angle and covering 180 °. The full 27r plots were built by repeating the measured data exploiting the two-fold symmetry of the surface. The XPD results were interpreted on the basis of the single scattering cluster-spherical wave (SSC-SW) theory [10]. The crystal structure was described as a cluster whose dimension was limited setting a maximum electron path from the emitting atom to the scatterer atoms of 8 A. Larger values of this parameter did not produce any significant change in the calculated XPD curves. The phase shifts for titanium and oxygen were derived from the muffin-potential calculated for TiO. The surface was cleaned by means of cycles of argon ion bombardment (2 keV) and annealing up to 850 K. The surface of this N i T i alloy turns out to be very reactive towards the residual atmosphere in the vacuum chamber and a small quantity of oxygen was always observed in the LEIS spectra.
3. Results 3.1. X P S and L E I S results
The XPS spectra for the clean surface of the alloy show that the position of Ti 2p3.2 is shifted
218
A. A trei et al. / SurJace Science 391 ( 1997) 216 225
by 0.4eV towards higher binding energy with respect to the value of the pure metal (454.0 eV) [8]. No chemical shift, in the limit of the accuracy of the measurements, could be observed for the Ni 2p3/2 peak in the alloy compared with pure Ni. The XPS and LEIS data indicate that there is no preferential sputtering of either components and a stable surface composition was reached after cycles of ion sputtering and annealing at 850 K. The oxidation procedure of the alloy surface consisted in exposure to an oxygen pressure in the 10 7 Torr range at 780 K and subsequent cooling (down to 500 K ) in oxygen atmosphere. The Ni 2p XPS spectra show that nickel was not oxidized in these conditions. The Ti 2p XPS spectra measured for various oxygen exposures are shown in Fig. 1.
r-
d
_c
C
b
470
465
460
455
450
Binding energy (eV) Fig. 1. Ti 2p XPS spectra measured after exposures to an oxygen pressure in the 10 v torr range of the alloy surface at 780 K. (a) Clean surface; (b) 30 L plus 30 L during cooling; (c) 60 L plus 30 L during cooling; (d) 150 L plus 100 L during cooling down to ca 500 K. For the last spectrum the components corresponding to Ti ~v (binding energies Ti 2p3/2 and Ti 2pi/2 459.0 and 464.6 eV, respectively) and to Ti m (binding energies Ti 2p3/2 and Ti 2pl/2 456.9 and 461.4, respectively), determined from the curve-fitting analysis, are shown as dashed lines.
Already for relatively low oxygen exposures, the oxidation of the alloy leads to a relevant increase of the surface concentration of titanium (in the form of oxide), as indicated by the growth of Ti 2p peaks at higher binding energies compared with the clean surface. The binding energy of the Ti 2p3/2 peak in the oxide film formed after an oxygen exposure of 30 Langmuir (L) (1 L a n g m u i r = l x l 0 6 T o r r . s ) at high temperature plus 30 L during cooling was 458.6 eV; this value is lower than that in bulk TiO 2 (459.0 eV) but coincident with the Ti 2p3/2 binding energies measured for the TiO 2 layers formed upon oxidation of Pt3Ti alloy surfaces or deposited on Pt surfaces [11]. The O ls binding energy was 530.2 eV. The oxide layer formed under these conditions covers completely the surface of the alloy, since no Ni signal is observed in the LEIS spectra (Fig. 2). The thickness of the oxide layers was estimated from the attenuation of the Ni 2p3/2 signal produced by the oxide film covering surface using the attenuation length for the Ni 2p3/2 photoelectrons calculated according to the Seah and Dench method [12]. From the XPS data, we found that the oxide layer was ca 3 A thick. The Ti/O atomic ratio determined for this phase turns out to be 1:2.1_+0.1. The T i 2 p spectrum measured after an oxygen exposure of 60 L at 780 K and plus 30 L during cooling is shown in Fig. 1, curve c. A shoulder appears on the low binding energy side of the Ti 2p3/2 peak while the main peak is shifted towards higher binding energy with respect to spectrum in Fig. 1, curve b. The curve fitting analysis showed that the Ti 2p3/2 peak consisted of two components at binding energies of 459.0 and 456.6eV. The former component corresponds to TiO2, whereas the latter is typical of titanium suboxides, TiOx, with x ranging from 1.0 to 1.5 [11]. Beyond an exposure of ca 100 L and up to the maximum oxygen dose of 600 L used in the present work, no changes in the Ti 2p XPS spectra were observed and the oxidation process reached a saturation after an exposure of 150 L. The thickness of the titanium oxide film formed under these conditions was estimated to be i l k . We found that the amount of reduced titanium oxide with respect to the TiO2 quantity depends on the exposure to oxygen during cooling, being
A. Atrei et al. ,/Surjitce Science 391 (1997) 216 225
was 530.5 eV. The annealing at ca 700 K for a few minutes of the oxide films prepared under the conditions described above, produced an increase in the amount of titanium suboxides TiOx coming from the reaction between TiO 2 and metallic titanium diffusing from the bulk of the alloy.
He+ 1keV
/
219
Ni
3.2. L E E D observations Ti
O
.__>, e-
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.,A,__
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0.3
I
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I
.,,%__ I
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,
I
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E/E o Fig. 2. LEIS spectra measured for the clean surface (curve a); after 30 L of oxygen at 780 K (curve b); after 150 L of oxygen at 780 K (curve c).
larger for shorter oxygen exposures. From the curve fitting analysis of the Ti 2p XPS spectrum for the oxide film formed after an oxygen exposure of 150L at 7 8 0 K plus 100L during cooling (Fig. 1, curve d) it turns out that in the Ti 2p3/2 there is the component corresponding to Ti TM, located at 459.0eV, and a component at lower binding energy (456.9 eV) due to titanium in lower oxidation states. The Ti 2p3/2 binding energy of this suboxide is close to the values measured for TiOx thin layers formed upon oxidation of titanium deposits on Ni (457.0 eV) [13] and Pt surfaces (456.2 eV) [14], and upon oxidation of the Pt3Ti alloy (456.8) [15]. In all these cases this component was attributed to Ti m, though the values of the binding energy are in between those characteristic of bulk TiO and Ti203 [16]. The O ls binding energy measured for this oxide layer
The clean surface of the alloy showed a (1 x 1) low-energy electron diffraction (LEED) pattern. Depending on the oxidation conditions different LEED patterns could be observed. The oxide phase formed at 7 8 0 K after 3 0 L of oxygen showed the LEED pattern reported in Fig. 3a. At first glance, taking into account the most intense spots, the superstructure seems to have a hexagonal unit cell. However, a closer inspection of the pattern reveals the existence of a c(2 x 6) coincidence mesh between the unit cell of the substrate and that one of the overlayer. Hence, the unit cell of the oxide overlayer cannot be hexagonal since it cannot fit the rectangular periodicity of the substrate surface. The surface unit cell of this oxide phase can be determined exactly from the c(2 x 6) coincidence net. The matrix notation for this surface unit cell is: -3 1
0
7 6 7
The sides of the unit cell are 2.93 ,~ and 3.04 and the angle between them is 121.2 ° . Since the surface unit cell of this oxide film is very close to a hexagonal mesh, we call it a "quasi-hexagonal" unit cell. Upon increasing the oxygen exposure, the quasi-hexagonal superstructure gradually becomes weaker, disappears and eventually the pattern shown in Fig. 3b can be observed. This pattern is characterized by extraspots along the [110] direction and strikes along the [001] direction. Only the first order extra spots are visible in the diffraction diagram and the distance between these reflexes and the integral order spots is ca one-fifth of the side of substrate unit cell along the
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A. Atrei et al. / SurJace Science 391 ( 1997j 216 225
(a)
(b) Fig. 3. (a) LEED pattern observed after oxygen exposure of 30 L at 780 K of the alloy surface. Normal incidence of the electron beam, energy 90 eV. (b) LEED pattern produced by the oxide layers obtained after an oxygen exposure of 150 L at 780 K plus 100 L during cooling. Normal incidence of the electron beam, energy 68 eV.
A. Atrei et al. / Sur[itce Science 391 (1997) 216 225
[ 110] direction. These extra spots may be attributed to double diffraction through the periodicity of the substrate and of the overlayer, with the higher order spots being too weak to be observed. Along the [001] the TiO2 phase is disordered as shown by the presence of diffuse intensity. 3.3. X P D results
The XPD patterns for the Ti 2p3/2 and O l s photoemission peaks of the quasi-hexagonal TiO 2
Ols
phase are shown in Fig. 4. Although, the unit cell of the oxide film as determined from the LEED pattern is oblique, the XPD patterns appear sixfold symmetric because the deviations of the cell from a hexagonal mesh are so small than cannot be detected in our experimental conditions. The lack of features at polar emission angles below 40 50 ° is an indication that the oxide layer has a thickness of a few atomic layers. In order to interpret the XPD data we performed SSC calculations on the basis of a structural model having the quasi-hexagonal unit cell determined from the LEED pattern. The dimensions of the overlayer unit cell are such that it is not possible to fit in more than one atom (either Ti or O). Hence, to reproduce the TiO2 stoichiometry we considered a layer of titanium atoms sandwiched between two oxygen planes (Fig. 5). The titanium
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221
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,
,
,
j
Substrate atoms
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Fig. 4. Stereographic projections of the O ls and Ti 2p3, 2 X P D curves measured for the quasi-hexagonal TiOz phase. The intensities are represented as grey-scale levels. Along the radial directions, from the centre to the circle, the polar emission angle varies from normal emission 0 = 0 ' (normal emission) to 0 = 90'. The azimuthal directions of the substrate are reported. The region near normal emission was not measured for the Ti 2p photoemission peak since, after collecting some polar scan, we found that the curves were featureless for 0 below 40 5 0 .
[001]
f
,
4 .
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[ [ITol
Fig. 5. Schematic model for the quasi-hexagonal TiO 2 layer. Both the quasi-hexagonal unit cell of the overlayer and the c(2 x 6) coincidence mesh are shown. In this figure also the substrate atoms are drawn to show the epitaxial relationship between the overlayer and the substrate.
222
A. Atrei et al.
Sur{~tce Science 391 (1997) 216 225
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atoms are located in the centre of the triangles delimited by the nearest oxygen atoms. This model for the TiO2 overlayer structure has a lower symmetry than the substrate and domains, produced by the symmetry operations of the substrate structure, must exist in order to have an overall twofold symmetry. Therefore, in the calculations we considered in addition to the structure shown in Fig. 5 its mirror domain (obtained by reflection through the vertical plane containing the [001] direction). On the basis of this model a good level of agreement with the experimental data was obtained (Fig. 6). The best fit corresponds to an interlayer spacing of 1.1 _+0.1 A. This value is very close to that (1.0A,) calculated assuming the average T i - O nearest neighbour distance in bulk rutile (1.96 A) [17]. The residual disagreement between experimental and calculated curves might be due to a buckling of the oxide film. Because of the large number of non-equivalent atoms in the unit cell and the possible registries of the oxide layer with respect to the substrate, no attempt has been done to introduce in the calculations a rippling of the quasi-hexagonal layer. The structural changes of the titanium oxide layer upon further oxidation of the alloy at high temperature are clearly revealed by XPD. The X P D patterns measured for the Ti 2p3/2 and O l s photoemission peaks of the TiO2 film formed after an oxygen exposure of 600 L at 780 K are shown in Fig. 7. The modulations observed in the XPD patterns indicate that the thicker TiO 2 film grows epitaxially on the alloy surface. The Ti2p3:2 pattern is characterized by intense peaks at polar angles 0 = 70" along the azimuth corresponding to the [001] direction of the substrate (q~=0) and 0 = 5 4 'J along the azimuth corresponding to the [ 11-0] direction of the substrate (~b=90). In the O l s pattern the most intense features are located at 0 = 7 0 ° in the ~b=0 ° azimuth and 0 = 4 5 in the q~= 90 ~azimuth. The intensity enhancement of the
Fig. 6. Comparison between experimental(thick line) and calculated azimuthal XPD curves measured the polar angles 0= 70'and 0=60 for the quasi-hexagonal TiO2 phase. (a) O ls peak, (b) Ti 2p3,2 peak. For each curve the maximum value of the anisotropy defined as A1/lm,x%, where Al=Imax--lm~n is reported.
223
A. Atrei et al. / Sur/ace Science 391 (1997) 216-225
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Ti2P3/2 Min.
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Fig. 7. Stereographic projection of the O ls and Ti 2p3/2 XPD intensities measured for the TiO2 phase obtained after oxidation in the conditions described in the text. The azimuthal directions of the substrate are reported.
70° ~
27
60°
27
¢R t-
_=
diffraction features at higher polar angles is due to the relatively low thickness of the oxide layer [18]. The models for the SSC simulations of the experimental XPD curves were based on the structure of bulk T i O 2. T i O 2 exists in three polymorphic Fig. 8. Comparison between experimental (thick line) and calculated azimuthal XPD curves measured at various polar angles for the thicker TiO2 phase. (a) O ls peak. (b) Ti 2p3:2 peak. The calculations were performed for epitaxial islands having the futile (110) structure oriented, with respect to the NiTi alloy surface, in such a way that the short side of the oxide overlayer unit cell is parallel to the direction [1 TO] of the substrate.
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A. Atrei et aL / Suffdce Science 391 (1997) 216 225
224
crystalline structures: rutile, anatase and brookite [19]. We started the model calculations with the (110) surface of futile since this crystallographic plane is the most stable termination of futile. We tried different topmost layer terminations of the futile (110) islands and the interlayer spacing as well as the in-plane parameters were varied in the calculations. As shown in Fig. 8 a fairly good level of agreement between experimental and calculated XPD curves is obtained for the (110) bulk truncated structure of rutile oriented in such a way that the short side of the oxide unit cell is parallel to the [110] direction of the substrate. The model is also consistent with the observed LEED pattern for this phase, since every 5 unit cells of the substrate in [110] direction there is a coincidence with the periodicity of the rutile (110) surface, assuming a 6% expansion of the oxide overlayer. We found that the XPD calculations were sensitive to the termination of the surface and the best agreement was obtained for the termination of TiO2(ll0) (Fig. 9) proposed in Ref.[20] and recently confirmed by surface X-ray diffraction results [21]. However, as it may be expected, there is only a small effect of the variations of the topmost layer spacings on the calculated curves. In addition to the (110) orientation we tried also the (100) surface of rutile and the (100) surface of anatase. These surfaces were chosen because the rectangular unit cell of these oxides could fit the rectangular unit mesh of the substrate. None
Oxygen atoms
topmostlayer
Titanium atoms
O
0"° "" ""
2nd layer
bulk layer
Fig. 9. Schematic structure of the rutile(ll0) surface showing the oxygen-terminatedsurfacewith six-foldand five-foldcoordinated titanium atoms.
of these models gave a theory-experiment agreement comparable with that obtained using the rutile (110) surface.
4. Discussion
The XPS results for the oxidation at high temperature of the Ni94Tidl10) surface indicate that in the very early stages TiO2 islands grow whereas some amounts of titanium suboxides form upon further oxidation. This behaviour may be explained if the oxidation rate decreases upon increasing the thickness of the oxide layer. When the oxide phase is a few layers thick all the titanium diffusing from the bulk is oxidized to Ti TM. On the other hand, if the oxidation rate decreases when the oxide layer becomes thicker titanium atoms can react at the alloy oxide interface leading to the formation of reduced titanium oxides. This titanium suboxide contains Ti m, as revealed by the low binding energy shoulder of the Ti 2p3. 2 peak corresponding to Ti TM. The increase (ca 0.4 eV ) of the Ti 2p peaks for Ti TM and O ls binding energies observed when TiO~ is formed can be explained by a shift of the Fermi level (due to the partial occupation of the titanium oxide conduction band) towards the vacuum level. This shift results in an apparent increase of the core level binding energies, as suggested by Dwyer et al. for titanium oxide layers on platinum [22]. From the surface crystallographic point of view the most relevant result of this study is the demonstration of the possibility to grow a TiO 2 phase which has no equivalent in bulk crystals. In the early stages of oxidation of the alloy surface, a structure is formed (with a quasi-hexagonal unit cell) consisting of a layer of titanium atoms between two planes of oxygen atoms instead of the growth of oxide patches having one of the crystalline structures (perhaps slightly distorted) of bulk TiO 2. To our knowledge this is the first time that the structure of an epitaxial oxide layers produced by the oxidation of an alloy surface cannot be derived from that of some bulk oxide. For instance, in the early stages of oxidation of Ni3AI(001), Ni3AI( 111 ) [23] and NiAI(110) [24] a fraction of the surface was covered by islands of
A. Atrei et al. / Surlace Seience 391 (1997) 216 225
7-alumina having a thickness comparable to the length of unit cell side in the direction perpendicular to surface. Also the LEED pattern produced by the oxide layers formed upon oxidation of the (111) and (100) surfaces of the Pt3Ti alloy were interpreted in terms of the structure of closed packed planes of bulk titanium oxides [25]. The reason why in the earlier stages of oxidation the quasi-hexagonal TiO2 phase forms instead of islands having the structure of one of the bulk phases of TiO2 appears to be the corrugation of the substrate surface. It is conceivable that a closed packed structure like rutile (110) can hardly fit the corrugation of a fcc(110) surface (in particular in the [001] direction, perpendicular to the grooves). On the other hand, bulk TiO 2 phases should form upon oxidation of the more flat (100) and ( 111 ) surfaces of the NiTi alloy.
5. Conclusions We found that upon exposure to oxygen of the Ni94Ti6(110) surface at high temperature layers of TiO2 form at the surface. The LEED and XPD data showed that while relatively thick oxide films consist of epitaxial islands of rutile oriented along the (110) surface, the quasi-hexagonal TiO2 phase has its own structure with no equivalent in bulk titanium dioxide phases. The formation of this two-dimensional oxide phase may be related to the corrugation of the substrate which hinders the growth with the structure of one of the bulk TiO 2 crystalline phases.
Acknowledgements
This work was partially supported by MURST. The authors are grateful to Mr E. Fisher for the
225
preparation of the NiTi sample and to Professor G. Kostorz for the useful discussion.
References [1] [2] [3] [4] [5]
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