Si interface formation studied by photoemission

Surface Science 210 (1989) 99-113 North-Holland, Amsterdam

99

MECHANISM OF THE Y /Si INTERFACE FOIWIATION STUDIED BY PHOTOEMISSION A. PELLISSIER,

R. BAPTIST

Division LETIf DOPT, CENG, 85X, F-38041 Grenoble Cedex, France

and G. CHAUVET DRF, G/Service de Physique, PSC, CENG, 85X9 F-38041 Grenobie Cedex, France Received 29 June 1988; accepted for publication 31 October 1988 XPS and UPS analyses have been used to study the Y/Si interface, obtained by deposition of yttrium atoms on a clean silicon surface. They reveal a three-step-mech~ism of formation, at room temperature. First, a pure yttrium thin layer forms. Then, when the Y thickness reaches roughly two monolayers, the abrupt interface becomes unstable and Si atoms diffuse through the yttrium. The resulting Si concentration is roughly 5-7 at% if the Si atoms are distributed in a homogeneous way over the whole yttrium overlayer. Finally, above four monolayers, metalhc yttrium grows.

For about twenty years, an increasing interest in surface and interface phenomena has developed in ~croelectro~cs and optoel~tro~cs. The reason for this interest is the central role played by surfaces and interfaces in device performance [1,2]; consequently, new devices, which require a thin transition region between two different materials have been conceived [3,4]. The metal-semiconductor contact is widely used in microelectronic circuits as an ohmic contact or Schottky barrier [5-111, Whereas the interfaces between silicon and silicides, noble, near noble and refractory metals have been studied for several years [12-171, the Y-Si system, like the rare earth-Si systems, constitutes a new field of investigation and forms the subject of very few publications [5,18-241. In this study, we focus on the formation of the interface between a thin film ( - l-40 A) of yttrium and a clean silicon (111) oriented surface. Using ultraviolet and X-ray photoemission spectroscopy (UPS and XPS), we present information about the early stages of formation of the Y-Si interface at room temperature, its electrical properties and its chemical nature. 0039~6028/89/$03.50 Q Elsevier Science Publishers B.V. (North-Holland Physics Publishing Division)

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2. Experimental Our aim was to deposit thin, pure yttrium films on clean Si(ll1) surfaces and to analyse them in situ. Initial Si substrates (n-type, P-doped, p = 2.7 !J cm) were cleaned in situ by standard techniques to make reproducible Si(lll)-7 x 7 surfaces, before any yttrium deposition. This surface geometry was obtained by Ar+ ion bombardment cycles (2 x 10m5 mbar, 600 eV, lo-20 PA cm-*), followed by annealing at 850°C and cooling down to room temperature. The existence of the 7 X 7 surface reconstruction was established by observing its characteristic LEED pattern. We undertook a series of pure yttrium (99.9%) depositions on Si substrates held at room temperature, using evaporation by electronic bombardment of an Y ingot at about 1200” C. Before deposition, the yttrium was purified by annealing for several days, in order to eliminate the fluorine contaminant and to reduce the oxygen and carbon ones to roughly 1%. The interaction between yttrium and residual hydrogen could not be entirely suppressed [25]. Twelve samples were prepared successively, with a constant Y deposition rate but increasing deposition times. For the first four depositions, we took as hypothesis that no attenuation of signals occurred through the thin yttrium overlayers. The amount of yttrium was estimated from the Y 3d and Si 2s peak areas and by taking into account the Al Ken photoemission cross sections (Si 2s: u = 0.955; Y 3d: u = 5.98 [26]), and the electron mean free path in silicon (X = 25 A [27]). This allowed an estimation of the Y deposition rate (- 1 A/min + 25%). The extrapolated thicknesses of the thicker films were then assumed to be directly proportional to the evaporation time. This method proved to be as reliable as the quartz microbalance method, which was also installed in our preparation chamber. Furthermore, this calibration procedure was also verified by the decrease of the area ratio of the Si LW (E, = 92.5 eV) peak to the Auger Y (EC = 79.5 eV) peak from 6 = 0.6 ML to B = 12 ML, the measurements of which take advantage of the low electron mean free path in this energy range ( - 5 A). The deposited yttrium overlayers are probably amorphous or microcrystalline, but in order to simplify the calculations, we adopted a model in which the thin film is considered to be a regular stack of monolayers (ML), successively separated by 3 A (this is roughly the lattice parameter in the hcp structure of yttrium). Atoms which form a monolayer are situated in the same plane, parallel to the surface. Since the Y monolayer is, in this model, referred to the monocrystalline Si(lll), it contains 7.8 x 1014 atoms cme2, that is to say the Si(III) atomic surface density. We verified that the resulting density of this phase is lower than that of the real hcp phase of Y (the densities are 3.8 and The amount of Y deposited ranged from 0.2 to 12.4 4.3 g cme3 respectively). ML. This hypothesis of a layer-by-layer mechanism of Y growth is supported by the “regular” decrease of the intensity ratio Si 2s/Y 3d which leads at

A. Pellissier et al. / mechanism of Y/S vacuum

;

siticon

vacuum

101

interface ~orrnat~~n iyttrium:

slllcon

Si2P Si2s

Fig. 1. Energy-band diagrams of n-type Si(lll)-7 x 7: (a) before Y deposition; (b) in contact with the Y overlayer. A consistent energy scale is used in (a) and (b).

higher coverage to the same number of deposited monolayers as supposed in the calibration procedure. The residual pressure in the preparation chamber was 2 x 10-” mbar and we observed an increase to (5-8) X lO_” mbar during Y evaporation. Just after Y deposition in the preparation chamber, samples were directly transferred into the analysis chamber and analysed by UPS and XPS. For our photoemission studies we used a VG ESCALAB with Al I&Y radiation (1486.6 eV) and 0a UV resonance lamp (He I, hv = 21.2 eV). Analysed depths are - 75 and 30 A respectively, and resolutions - 1 eV and 50 meV. Binding energies ( Eb) were measured relative to the Fermi level energy E,, calibrated itself by means of a gold sample (E,(Au 4f) = 84 eV).

3. Results and discussion We first establish the initial energy-band diagram of our silicon substrate, in order to know the influence of the yttrium overlayer on it (fig. la). The presence of surface states, specific to the 7 X 7 geometry causes a depletion zone near the Si surface. This leads to a band bending e#,, and a which are obtained from the Si work function surface potential barrier e&, e+ = 4.55 eV for the Si-7 x 7 surface (this work), the electron affinity ex = 4.01

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of Y/Si

interjace formation

eV [28] and the resistivity value p = 2.7 Q cm. Their values are e+, = 0.25 eV and eGbn = 0.54 eV. The depletion width W calculated from Poisson’s equation is: W= 1700 A. Thus, only a weak portion of the total depletion width, W, is analysed by photoemission spectroscopy (UPS He I: 2% and XPS: 4.5%).

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ML

06

ML

02

ML

ENERGY

(eV)

Fig. 2. Si 2p energy distribution curves measured by XPS for a Si(lll)-7 ~7 surface with increasing Y coverage. The shoulder is marked with a vertical bar. The spectra are normalized to the maximum peak intensity. Note that the upper curve corresponds to an exit angle of 35 O.

A. Pellissier et al. / Mechanism

3.1.

of Y/ Si interface formation

103

Results

We now consider the XPS analysis of yttrium thin films, deposited on silicon substrates. For each of the Si 2s, Si 2p (fig. 2) and Si Auger KLL (fig.

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ENERGY

Fig. 3, Si Auger KLL energy distribution curves measured with increasing Y coverage. The spectra are normalized

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A. Pellissier et al. / Mechanism

104

3) levels, we see a principal yttrium coverage:

of Y/Si interface formation

peak, whose binding energy is constant

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But from 0 = 2.3 ML to 6 = 6.2 ML, each of the Si 2p and Si Auger KLL levels presents a shoulder, located respectively at 1.4 eV lower binding energy and 1.7 eV larger kinetic energy. Each shoulder grows relative to the principal peak from 8 = 2.3 ML to 8 = 4.1 ML and then stays in the same proportion up to 0 = 6.2 ML. At B = 6.2 ML, the area ratio in the Si 2p spectra is evaluated by decomposition of the total curve in two components (shoulder and principal peak). The intensity ratio of shoulder to principal peak is roughly 2.5%. In Si 2s the corresponding shoulder is probably hidden by a satellite due to the Y 3d level excited by the Al Kq,, radiation. The low-angle detection favours the relative importance of the shoulders in comparison to the principal peaks, signifying that the corresponding Si atoms are located above the Si bulk. Beyond 6.2 ML, for example in our 12.4 ML Y thick sample, the silicon XPS signals are weak because of the great absorption through Y and the relatively low silicon photoemission cross section. Fig. 4 represents the shape of the Si 2s peak and the Y 3d,,, 5,2 doublet at 8 = 4.1 ML and fig. 5, the evolution of the Y 3d,,,, 5,2 core level binding energies as a function of yttrium coverage (0). The Y 3d,,,, 5,2 spin-orbit separation is 2 eV. We observe a shift by -2.2 t_ 0.2 eV from E,(Y 3d,,,) = 158.2 + 0.1 eV (6 = 0.2 ML) to E,(Y 3d,,,) = 156.0 f 0.1 eV (8 = 3 ML) and a stagnation at the Y bulk value. Although no shoulder can be detected in the doublet, as it is clearly observed in the Si 2p peak, the valley between the

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BINDING ENERGY (eV) Fig. 4. XPS Si 2s and Y 3d core level spectrum at 6 = 4.1 ML

A. Peliissier et al. / Mechanism

0.5 Y COVERAGE

of Y/S

1.0

interface formation

5.0

105

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0 (ML)

Fig. 5. Y 3d,,,, s,* binding energies (Eb) as functions of the Y coverage (0). Horizontal bars are E, for bulk Y (A E f: + 0.1 eV is indicated by vertical bars).

presents a variable depth which could be and Y 3d,,, components Y 3d,,, due to a small additional structure. But the resolution is not high enough to separate this hypothetical doublet. UPS spectra (Fig. 6) show how the valence band changes when Y is deposited on n-type Si(ll1). The surface state near the Fermi level, which characterizes the metallic nature of the Si(lll)-7 X 7 surface, rapidly disappears with the first Y depositions (0 I 8 I 0.8 ML). Up to 13= 1 ML, there is no metallic character at the Fermi level, but above this thickness, the development of a metallic band is expressed by the more and more abrupt Fermi level edge (1.4 I (YI 12.4 ML). With increasing yttrium coverage, some new peaks appear. We first notice a bump at 8 = 0.2 ML, whose binding energy is about 3.8 eV (peak A). It becomes higher and more refined up to 1 ML and then weaker, until it vanishes (see the last sample at fl= 12.4 ML). Its growth is accompanied by a shift of 0.4 eV to higher binding energies. The second peak (peak B) arises at about 1 ML with a binding energy of roughly 5.4 eV and becomes higher for thicker yttrium depositions. Peak B has already been underscored in UPS spectra of massive Y samples and of thick yttrium layers [25]. It is attributed not to a contamination by oxygen but to a contamination with hydrogen (yttrium, like titanium, is a good getter for hydrogen), which results in the formation of a surface yttrium hydride or of a thin film with hydrogen in solution. From the UPS measurements, we can deduce the evolution of the work function e+ with yttrium coverage (fig. 7). From The Si(lll)-7 x 7 surface value (4.55 ev), it rapidly drops to 3.9 eV at 8 = 0.2 ML. This decrease continues more slowly from 8 = 0.2 ML to 0 = 1.3 ML, where e+ = 3.6 eV. Between 1.3 and 2.5 ML, the slope is steeper and the curve reaches the bulk

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A. Pellissier et al. / Mechanism of Y/Si

interface formation

12i ML

6.2 MI1.1 ML 31 ML 2.3 ML 18 ML

BINDING

ENERGY

k.‘)

Fig. 6. UPS (He I) valence band spectra at increasing Y coverage. Dashed lines indicate peaks A and B.

value (3.1 + 0.1 eV [25]) at 8 = 2.5 ML. Since the variation is rapid for low Y coverages, we made a supplementary measurement at 8 = 0.15 ML and verified that this variation is regular. 3.2. Discussion We notice that the yttrium deposition has no influence on the initial potential barrier value of n-type Si(lll), henceforth denoted as Si(lll)-n. In effect, since metal induced changes in Si band bending are reflected by shifts

A. Pellissier et al. / Mechanism of Y/S

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1.0

y COVERAGE

107

interface formation

10

8 (ML)

Fig. 7. Work function (e+) evolution versus Y coverage. The Y bulk value of e+ is indicated by a horizontal bar [25]. The samples are biased with a negative voltage (4.57 V) in order to have an abrupt photoemission threshold (see insert).

in the Si bulk substrate core level binding energies, we conclude that no change occurs in Si band bending as no shift is observed. In other words, the Schottky barrier height (SBH) of the yttrium/silicon (n-type) contact has the same value as the initial surface potential barrier height: e+,,(Y/Si(lll)-n) = 0.54 f 0.1 eV. This value is higher than that found for the YSi,/Si-n contact which is 0.3 eV [20,29]. Although there is a change in chemical bonding at the Si interface, the Fermi level is pinned at the same value relative to the situation without yttrium deposition (fig. lb). According to Norde et al. [29], the sum of eGbn and e&r (the value of the SBH for p-type Si) is the band gap of Si (1.12 eV). Their value of eGbp is 0.69 eV for the Y/Si interface. They should have a corresponding e$,,” value of 0.43 eV. Since our measured eGbn value is 0.54 k 0.1 eV, our disagreement with their result is roughly 10% of the gap. In the following paragraphs, we propose a mechanism of Y growth in which three steps are emphasized. First yttrium atoms settle on a clean Si(ll1) surface, to produce a very thin overlayer (fig. 8a). The presence of the UPS peak A is interpreted as a characteristic of the chemical bond between the first Y monolayer and the upper Si atoms (and possibly hydrogen atoms), because this peak does not exist in massive Y samples [25] and reaches its maximum intensity at fl= 1 ML. We suppose that it reflects a Si 3p-Y 4d hybridization as is encountered in the YSi,,, silicide, although its binding energy is - 3.8 eV instead of 2.7 eV in YSi,., and the Y 3d,,, is located at 156.6 eV instead of 155.6 eV. LEED measurement of the 1 ML thick Y sample shows hardly discrete reflections, indicating that the Si-7 X 7 surface is completely destroyed by the metal, and that the first monolayer of yttrium covers all of the Si surface and no cluster forms.

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A. Pellissier et al. / Mechanism of Y/ Si interface formation

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Fig. 8. Schematic representation of Y/Si interface formatton. Steps (a), (b) and (c) correspond to the first proposed mechanism of Y growth and (a), (c) and (e) to the second.

Secondly, when the yttrium thickness reaches roughly 2 ML, some silicon atoms diffuse into the yttrium. This phenomenon has already been mentioned by Hiraki [30] for some transition metals. The presence of a critical metal film thickness (2 ML in our case) is necessary to break the strongly covalent Si lattice. These upper diffused Si atoms are responsible for the appearance of the low-binding-energy shoulder in the Si 2p and the Si Auger KLL spectra, because they are in an atomic environment different from the bulk. The evolution of the Si 2p and Auger KLL shoulder intensities expresses that the amount of diffused Si atoms increases progressively and stops at roughly four monolayers. Unlike the photoemission results for refractory metal/silicon contacts, the different measurements obtained in this work (Si 2p, Y 3d binding energies, work function and UPS spectra evolutions) do not seem to be directly correlated. For example, between B = 0.2 ML and B = 1 ML, e# is constant while the other parameters are not. Likewise, only e# changes suddenly from B = 1 ML to 2 ML. In this way, the Y/Si interface reveals a different behaviour from the refractory metal/silicon one. This leads us to ask two questions: __ through how many yttrium monolayers do Si atoms diffuse? - what is the resulting Si density distribution?

A. Peftissieret al. / Mechanismof Y / Si interfaceformation

109

In order to answer these questions, two mechanisms of growth are proposed. In both of them, the amount of diffused silicon will be evaluated, using the area ratio of the Si 2p shoulder to the Si 2p principal peak, the area ratio of the Y 3d to the Si 2p core level peaks and also by taking into account the electron mean free path in yttrium which is - 25 A for E, = 1330 eV [31]. In the two models, the metallic yttrium thickness grows until 4 ML, while Si atoms are progressively accumulating very close to the interface (first model, fig. 8b) or distributing in a homogeneous way over the whole growing yttrium thickness (second model, fig. SC). The resulting Si densities are respectively 30% and 5%-7%. When the Si diffusion stops, the corresponding yttrium thickness is four monolayers. The second mechanism seems to be the more probable: the yttrium atoms are not perturbed by the low silicon density and stay approximately in the same electronic configuration as the pure yttrium overlayer and no additional doublet appears. We note that effectively the Y 3d binding energy curves (fig. 5) do not have any plateau. Otherwise, if the overlayer was silicon-rich, the work function would not reach the Y thickness value at 8 = 2-3 ML. By observing Si 2p energy distribution curves of yttrium/silicon and other metal/silicon contacts, de1 Giudice et al. [32] had observed, with synchrotron radiation (hv = 112 ev), different Si 2p spectra. Such an energy choice allowed them, better than with X-ray photoemission, to analyse the surface, relative to the bulk. In their experiment, when Y was deposited, two new structures appeared successively in the Si 2p curves: they were located at A E, = - 0.8 eV and AE,, = - 1.25 eV from the bulk value and are interpreted respectively as representative of a 50 at% Si compound and then of a solid solution. The respective thicknesses were not mentioned. Our experimental results could not bear out this mechanism. If our samples presented the formation of two successive compounds with such great concentrations, at the interface, we would have observed the two corresponding Si 2p shoulders and principal peak. In fact we detect only one shoulder, whose intensity is representative of a much lower amount of diffused Si atoms. In the third step of the mechanism, at 6 = 4 ML, the Si 2p and Si Auger KLL shoulders stop increasing, showing that Si diffusion ceases, then pure Y metal grows (figs. 8d and Se). At 8 = 12.4 ML, no trace of silicon is seen in XPS spectra and the UPS valence band is that of a metallic Y thick film [25]. In these two proposed models, the strong Y-Si interfacial chemical bond, which is responsible, in our opinion, for the UPS peak A, is not destroyed by the small amount of diffused atoms close to the interface. That is why this peak does not disappear but becomes weaker as the interface is more and more bidden by the growing yttrium thickness. The penetration of hydrogen atoms into the yttrium overlayer and the resulting chemical bonding is the origin of peak B in the UPS spectra, corresponding to the affinity level of the H- ion [25].

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A. Pellissier et al. / Mechanism of Y/ Si inierface formation

We do not explain why the work function falls off very slowly between 0.2 and 1 ML, showing then a plateau which is not connected to an almost constant binding energy for the Y 3d core level over this same range. As proposed by one of the referees, it is possible that the change in e$ between 1 and 2 ML is caused by either (i) a lower atomic density in the second and subsequent layers, or (ii) increased surface disorder (e.g. the starts of simultaneous multilayer growth). Let us now observe the evolution of the Y 3d binding energy curves: for very low Y coverage (9 = 1 ML), XPS spectra reflect the bonding properties between the yttrium overlayer, the silicon substrate and possibly hydrogen. An electron charge transfer from the yttrium atoms to the silicon ones causes an increase of the binding energy of the Y 3d core levels in comparison with the metal bulk level value. As the amount of Y becomes greater, this effect weakens for all the Y atoms, thus the Y 3d binding energy decreases progressively. We do not observe any shift for the Si 2s and Si 2p peaks, as it could be deduced from a simple charge transfer mechanism. We suppose that this effect is hidden by the XPS Si bulk contribution. The Si diffusion induces also an electron charge transfer between the diffused Si atoms and the Y atoms which are in their neighbourhood. This charge transfer does not affect the Y 3d core level because of the low density of silicon. It is evaluated from the positions of the shoulders in Si 2p and Si Auger KLL spectra, using the following relations [33]: AR”,“(Si) = +0.15

eV,

AV+ 1.25 eV,

where ARg(Si) is the extra-atomic relaxation energy and Av, the potential variation which is proportional to the charge transfer. The positive value of AV signifies that electrons are transferred from yttrium to silicon. If the value of the chemical shift per lost electron for Si 2p is + 2.2 eV, as estimated by Ley [34], each Si atom receives 0.57 electron from its all yttrium neighbours. This electron transfer is more important than in bulk disilicide YSi, [35] (Ak’ for YSi,, is +0.2 eV). In these two cases of electron charge transfer between Y and Si atoms, the transfer occurs from Y to Si. Relative to their bulk values the shift of Y 3d is toward higher binding energies and the shift of Si 2s/2p is toward lower binding energies. A study has been carried out under the same experimental conditions by Azizan [36], on four refractory metal (Nb, Ta, W, Mo)/silicon interfaces. In the Nb/Si and Ta/Si cases, the abrupt interface becomes unstable above a critical metal overlayer of 1 ML, and a chemical reaction occurs up to 6 and 4 metal monolayers respectively. The resulting interfacial compound has a near disilicide composition. The following stages are a pure metallic phase for the formation for the Ta/Si system in Nb deposition and a Taloo-, Si, compound which the Si concentration decreases with increasing Ta thickness. On the contrary, the W/Si and Mo/Si interfaces stay atomically abrupt.

A. Pellissier et al. / Mechanism of Y/ Si interface formation

111

A three-step-mechanism has already been found for rare-earth metals (Eu [37], La [24], Sm [38], Yb [39]) deposited on Si substrates, in the same experimental conditions but using synchrotron radiation photoemission. The first stage of growth is the metal chemisorption on the clean Si surface for low coverages. Above a critical metal thickness (Eu: 5 ML, La: 0.1 ML, Sm: - 1 ML, Yb: - l-2 ML), a reactive interdiffusion of RE and Si atoms occurs, allowed by the screening of the covalent bonding in the upper Si substrate atoms. The intermixing ceases when the metal coverage reaches a limiting thickness (15 ML for Eu, 2 ML for La and 20-25 ML for Sm), above which pure rare-earth metal grows. The intermediate Eu-Si and Sm-Si compounds are thus thicker than the Y-Si one. The mechanism that governs the formation of the Ce/Si interface is more complex [29] because in the very first stages of chemisorption Ce clustering dominates. The final interface is mixed. The La/Si interface [24] shows a similar trend to the Y/Si one. For very low coverages ( < 2 ML), a charge transfer occurs from La to Si, as for Y to Si; then around 2 ML the reacted layer is silicide-like, as observed with the presence of peak A. In order to predict the abrupt or diffuse nature of interfaces between transition metals and silicon, Azizan et al. [36,40] have introduced a parameter R, which is characteristic of the abrupt or diffuse nature of interfaces between transition metals and silicon. R is defined as the ratio of the silicide formation -A Hf (kcal/metal atom) to the electronic heat capacity y (mol enthalpy, K2/mJ). The interface is abrupt if R is higher than 2.6 and diffuse if R is lower than 1.7. W/Si, Mo/Si, Nb/Si and Ta/Si whose values of R are 5.6, 5.2, 1.6 and 1.4 respectively, verify this model. In our Y/Si interface, y = 10.2 the Ta/Si and Y/Si [41] and -AH, = 16.1 [42], thus R = 1.6. Although interfaces have the same value of R, the Si diffusion is much greater and more extensive in the first case than in the latter one. The Y/Si interface therefore falls outside the range of validity of this model.

4. Conclusion This paper has presented an experimental study of the yttrium/silicon interface, obtained by yttrium evaporation onto a clean Si(lll)-7 X 7 surface, at room temperature. We propose two possible mechanisms for the yttrium growth description, in order to interpret XPS and UPS analyses that give rather different results from that of the refractory metal/silicon interfaces. Up to 2 ML, the atomically abrupt interface is stable. But as the yttrium overlayer becomes thicker, silicon atoms diffuse. In the first proposed mechanism, they accumulate between the first and the second yttrium monolayers, while the yttrium keeps growing. In the second one, they diffuse widely and form a homogeneous solid solution of

112

A. Pehsier

e; al. / ~e~h~~~rn

of Y,/Si

interface forrn~~~~~

Si in the yttrium overlayer (which electronic structure looks like the silicide one). This mechanism seems to be the most probable. Above 13= 4 ML ( - 12 A), silicon atoms do not longer diffuse and yttrium continues to grow. It seems that hydrogen, always present in a vacuum chamber, plays a role at the interface and eventually during the silicon migration.

References [I] J.J. Scheer and J. van Laar, Surface Sci. 18 (1969) 130. [2] S.M. Sze, Physics of Semiconductor Devices (Wiley-Interscience, New York, 1981). [3] J.J. Andrews and J.C. Phillips, Phys. Rev. Letters 35 (1975) 56; CRC Crit. Rev. Solid State Sci. 5 (1975) 405. [4] J.G. Clabes, G.W. Rubloff, B. Reihl, R.J. Purtell, P.S. Ho, A. Zartner, J.F. Himpsel and D.E. Eastman, J. Vacuum Sci. Technol. 20 (1982) 684. [S] J.E. Rowe, J. Vacuum Sci. Technol. 13 (1976) 798. [6] G. Margaritondo, J.E. Rowe and S.B. Christman, Phys. Rev. B 14 (1976) 5396. [7] E.H. Rbodevick, Mets-~~conductor Contacts (Clarendon, Oxford, 1978). [8] J.M. Poate, K.N. Tu and J.W. Mayer, Eds., Thin films - Interdiff~ion and Reactions (Wiley, New York, 1978). [9] J.L. Freeouf, Solid State Commun. 33 (1980) 1059. [lo] G. Ottaviani, K.N. Tu and J.W. Mayer, Phys. Rev. Letters 44 (1980) 4; G. Ottaviani, J. Vacuum Sci. Technol. 18 (1981) 924. IllI M. Wittmer, P. Oelhafen and K.N. Tu, Phys. Rev. B 35 (1987) 9073. WI R.J. Purtell, P.S. Ho, G.W. Rubloff and P.E. S&mid, Physica B 117/118 (1983) 834. [I31 B.Y. Tswar and C.H. Anderson, Jr., Appl. Phys. Letters 47 (1985) 527. 1141 T.N. Nguyen Tan, M. Azizan, R. Cinti, G. Chauvet and R. Baptist, Surface Sci. 162 (1985) 651. WI M. Azizan, T.A. Nguyen Tan, R. Cinti, G. Chauvet and R. Baptist, Solid State Commun. 154 (1985) 895. M. Azizan, T.A. Nguyen Tan, R. Cinti, R Baptist and G. Chauvet, Surface Sci. 178 (1986) 17. [I71 5. Lajzerowicz, These Docteur ingenieur de I’ENST, novembre (1986). 1181J.E. Baglin, F.M. d’Heurle and C.S. Petersson, Appl. Phys. Letters 36 (1980) 594. 1191 R.D. Thomson, B.Y. Tsaur and K.N. Tu, Appl. Phys. Letters 38 (1981) 535. WI J.A. Knapp and ST. Picraux, Appl. Phys. Letters 48 (1986) 466. WI V.M. Koleshko, V.F. Belitsky and A.A. Khodin, Thin Solid Films 141 (1986) 277. Appl. Phys. Letters 51 (1987) 311. WI M. Gurvitch, A.F.J. Levi, R.T. Tung and S. Nakahara, 1231 G. Rossi, Surface Sci. Rept. 7 (1987) 1. 1241 E. Puppin, H. Guyot, 2.X. Shen, J.H. Wang and I. Lindau, Solid State Commun. 67 (1988) 23. f251 R. Baptist, A. Pellissier and G. Chauvet, Z. Phys. B 73 (1988) 107. 1261J.H. Scofield, J. Electron Spectrosc. 8 (1976) 129. ~271C.R. Brundle, Surface Sci. 48 (1975) 99. et des Composants Electroniques (Masson, 12s1II. Mathieu, Physique des Semi-conducteurs Paris, 1987) p. 105. S. Petersson and P.A. Tove, Appl. [x91 H. Norde, J. de Sousa Pires, F. d’Heurle, F. Pesavento, Phys. Letters 38 (1981) 865. 1301 A. Hiraki, Surface Sci. 168 (1986) 74. I311 D.R. Penn, J. Electron Spectrosc. 9 (1976) 29.

A. Pellissieret al. / ~ee~anj~~ of Y/Si

interfaceformation

[32] M. dei Giudice, J.J. Joyce and J.H. Weaver, Phys. Rev. B 36 (1987) 4761. [33] P. Streubel, R. Fellenberg and A. Reif, J. Electron Spectrosc. Related Phenomena 261. [34] L. Ley, in: Semiconductors and Semimetals, Vol. 21, Ed. J. Pankove (Academic York, 1984) p, 385. [35] R. Baptist, A. Pellissier and G. Chauvet, Solid State Commun. 68 (1988) 555. [36] M. Azizan, These de Doctorat es Sciences Physiques de I’USTMG (Juillet 1987) [37] G. Rossi, J. Nogami, I. Lindau and J.J. Yeh, Surface Sci. 152/153 (1985) 1247. [38] A. Franciosi, J.H. Weaver, P. Perfetti, A.D. Katnani and G. Margaritondo, Corm-nun. 47 (1983) 427. [39] L. Braicovich, I. Abbati, C. Carbone, 3. Nogami and I. Lindau, Surface Sci. 168 [40] M. Azizan, T.A. Nguyen Tan and J. Derrien, Le Vide, Les Couches Minces 42 [41] Ch. Kittel, Introduction to Solid State Physics (Wiley, New York, 1976). 1421 K.N. Tu and J.W. Mayer, in: Thin Films Interdiffusion and Reactions, Eds. J.M. Tu and J.W. Mayer (Wiley, New York, 1978) p. 359.

113

34 (1984) Press, New

p. 262. Solid

State

(1986) 193. (1987) 219. Poate, K.N.