Si(100) systems

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ELSEVIER

Empty states investigation of epitaxial Sn/Si( 100) systems M. Pedio

*, V. Ghisalberti, Istituto di Struttura

C. Ottaviani, M. Capozi, F. Lama, C. Quaresima, P. Perfetti della Materia-CNR,

V. E. Fermi 38, I-00044 Frascati, Italy

(Received 29 July 1993; accepted for publication 4 October 1993)

Abstract The Sn on the Si(100) system constitutes an epitaxial interface of Group-IV elements. Its phase diagram presents several surface reconstructions, whose structure depends on the Sn adlayer coverage and the annealing temperature. The interfaces of Sn grown at room temperature on Si(lOO)(Z x 1) and the superstructures obtained after annealing have been studied by means of inverse photoemission (IPES), and Auger and LEED spectroscopy, to characterize the electronic empty states of the different overlayer phases and relate them with other electronic properties. Sn/Si(lOO) systems, both at room temperature deposition and after thermal treatment, present the onset of a metallic overlayer for tin coverages above 1 ML. For the interfaces grown at room temperature the IPES spectra with Sn thicknesses below 1 ML show a peak at about 1.0 eV above the Fermi level. At the E, level no emission is present. The IPES spectra of Sn/Si(lOO> after annealing clearly show that the empty states evolution can be related to the kind of reconstruction observed by LEED. These results indicate the formation of interfaces with peculiar properties. In particular in the Sn/Si(lOO) c(8 X 4) system the phase is semiconducting and the gap is reduced; a new prominent state appears at 1.1 eV above E,. The role played at the interface by the Sn-Sn dimers, related to the dimensionality, is discussed.

1. Introduction The study of the first stages in the growing process of epitaxial interfaces represents a very rich tool that enlightens the properties of low-dimensionality systems and the role played by the interface states on macroscopic properties like, e.g., Schottky barriers in metal-semiconductor junctions, whose formation is still not completely understood.

* Corresponding

author.

In the last few decades, much work has been devoted to the study of the filled electronic states of interfaces by means of photoemission spectroscopy, while the characterization of empty electronic states has been made possible only recently, by means of direct techniques like inverse photoemission spectroscopy (IPES). The aim of this paper is the study of the early stages of tin-silicon (100) interface formation. Sn on SXlOO) forms an epitaxial interface [l-4]. In contrast to the majority of metal-silicon interfaces, tin belongs to a small class of metals that do not form compounds with silicon [5]. A new surface reconstruction induced by a thin over-

0039-6028/94/$07.00 0 1994 Elsevier Science B.V. All rights reserved SSDI 0039-6028(93)E0575-F

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layer is an example of a low-dimensionahty system. We will show that Sn/Si interfaces, presenting a large variety of surface reconstructions [1,2], can be considered model systems to study the possible correlation between electronic and structural properties. Tin has two allotropic forms: a-Sn, semiconducting, stable at a temperature below 13.2”C (diamond structure, lattice constant a = 6.485 A, density 5.765 g/cm3), and @Sri, metallic phase, stable for T > 132°C (tetragonal structure, lattice constants a = 5.831 A, c = 3.182 A, density 7.285 g/cm3). The phase transition from /?-Sn to cY-Sn causes a big enhancement of the unit cell’s volume. It is possible to obtain high-quality films of metastable a-Sn at higher T, by means of epitaxial growth on different substrates, with a lattice mismatch < 1% (CdTe(llO), CdTe(lOO), InSb(ll0)) [6,7]. However, the lattice mismatch of Lu-Sn and the low-index Si surfaces is > 10%. Therefore, it is important to understand whether an a-Sn phase can be formed on SXlOO) and what is the Sn critical coverage in order to observe the semiconductor-metal transition. The phase diagram of Sn/Si(lOO), as a function of the tin coverage and annealing temperature, is extremely rich [1,2]. For Sn/Si(lOO) interfaces grown at room temperature CRT), the lowenergy electron diffraction (LEED) pattern at low coverages is (2 X 0, turning to (1 X 1) and progressively more diffuse as the tin thickness increases. In the literature the evolution of the electronic properties has been analyzed [3] by means of Auger and partial yield spectroscopy. Annealing at 550°C of interfaces with different tin coverages led to various superstructures [1,4] differing from the RT case. The electronic properties of such ordered superstructures have been described in the literature by direct photoemission [ll, together with a structural characterization RHEED and by scanning tunneling microscopy @TM) [2]. 2. Experimental

details

The experiments were performed in a UHV system formed by two interconnected chambers. The preparation chamber was equipped with a

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cylindrical mirror analyzer with a coaxial electron gun for electron energy loss spectroscopy (EELS) and Auger electron spectroscopy (AES) measurements (base pressure 2 X lO_“’ mbar). LEED analysis and angle-integrated IPES measurements in isochromat mode were recorded in the other chamber (base pressure 5 X lo-” mbar), using a home-made Geiger-Miiller-type detector filled with I, and He and closed by a SrF, window to select 9.5 eV photons. The electrons were supplied by a BaO cathode and the IPES spectra were taken with an overall resolution 5 0.35 eV, as measured by the Fermi level width of a tantalum polycrystalline sample. The reference position of the Fermi level (E,) was frequently measured by interchanging the sample with the tantalum foil. Before insertion into UHV, p-type double-domain Si(100) crystals (p = 7-13 R. cm) were cleaned with a wet chemical etching to obtain oxygen-free surfaces. The samples were then outgassed in situ by resistive heating, slowly increasing the temperature up to 850°C for several minutes, and then flashed at about 1100°C. The pressure in the preparation chamber during the cleaning procedure never exceeded 9 X lo-” mbar. The surface cleanliness was checked by AES. No detectable C and 0 Auger transitions were present after the cleaning procedure and clear double-domain LEED (2 X 1) patterns were observed. Tin was evaporated by a resistively heated crucible, keeping the Si samples at room temperature. The evaporation rate was calibrated by a quarz-crystal microbalance and the coverage was controlled by Auger spectroscopy. Typical deposition rates were between 0.2 and 1.0 ML/min, where 1 ML corresponds to 6.8 X 1014 atoms/ cm’. The different Sn/Si(lOO) interfaces have been characterized by means of LEED and Auger spectroscopy. IPES spectra were taken on Sn/Si(lOO) interfaces grown at room temperature and on the same interfaces heated at temperatures between 550 and 650°C for 2-5 min, to obtain different superstructures. The measurements on the reconstructed interfaces were performed at room temperature.

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3. Experimental 3.1. Auger spectroscopy results

Fig. 1 shows a set of SnMNN Auger spectra for interfaces grown at room temperature. The line shapes show the appearance of a new peak at about 407 eV (vertical bars in Fig. 1) for tin coverages above 1 ML. Fig. 2 shows the evolution of the SnNVV line as a function of tin coverage. Taking the minimum as a fingerprint of the Auger NVV transition, it is evident that there is a shift towards higher kinetic energies, with increasing Sn coverage; the Auger-transition energy shifts from 17.4 eV at 0.4 ML to 18.0 eV at 2.8 ML. It is worthwhile to note that a similar behavior of AES low-energy transitions involving a core and two valence band electrons (CVV) has been observed [S] for other elements grown on Si surfaces, corresponding to an increasing metallic character of the interface. For annealed systems the Auger results are similar to the previous ones, with the exception of the structure at about 407 eV that is clearly

Fig. 2. Evolution of SnNW Auger transition as a function of tin coverage, at room temperature.

present only for the Sn/Si(lOOKi x 1) superstructure (see Section 3.2.2) for which the SnNW Auger transition is at 17.8 eV. 3.2. Inverse photoemission spectroscopy results 3.2.1. Sn/Si(lOO)

430

480

J&W)

Fig. 1. SnMNN Auger spectra for different deposited at room temperature.

tin coverages

grown at room temperature

Fig. 3 shows angular-integrated IPES spectra for the Si(lOOI(2 x 1) double-domain surface and for Sn/Si(lOO) as a function of tin coverage. In the clean Si(100) IPES spectrum it is possible to distinguish a shoulder above the Fermi level up to about 0.6 eV, and a prominent state at 1.1 eV. The figure also shows the energy positions S, and S, of surface states, as obtained by angle-resolved IPES studies of the clean Si(100) single domain [9]. The two states have a nearly flat dispersion in the I-J direction of the Brillouin zone and are attributed to empty dangling-bond states for the 2 X 1 symmetric and asymmetric dimer models. Our experimental geometry is such that we detect all the electronic empty-state contributions with k values between 0” and 40” off normal. In the clean Si(100) spectrum of Fig. 3 we clearly see the S, contribution, while the shoulder above E, and up to 0.6 eV can certainly be

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156

Fig. 3. Angle-integrated IPES spectra of the Sn/Si(lOOI system at room temperature, for increasing tin coverage. On the left, the corresponding LEED patterns are indicated.

attributed to a convolution of the three contributions S,, S; and S;, reported in Ref. [91. Summarizing, in our angle-integrated TPES data on Si(100) double domain, the shoulder above E, and the peak at 1.1 eV correspond to Si surface empty states, while the peaks at 3.0 and 4.3 eV are assigned to bulk empty states [9] (see Table 1). Tin deposition at room temperature leads to a (2 x 1) LEED pattern for coverages up to 0.5 ML. The LEED for coverages of about 1 ML

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turns to a (1 X 1) pattern with an increasing background and the spots become progressively diffuse for higher coverages. The IPES spectra of Sn/Si(lOO) interfaces, corresponding to tin coverages below 1 ML, do not show any emission at the Fermi level. For low coverages the surface states close to the bottom of the conduction band are no more distinguishable, while peaks at about 1.1 and at 2.9 eV are clearly present. By increasing the tin coverage the shoulder at the bottom of the conduction band vanishes and for tin coverages of about 1.1 ML a structure develops at 0.8 eV. A slight tail of emission is detectable at the Fermi level (E,). The spectra assume a different shape after tin deposition exceeding 1 ML. At E, a clear emission is present together with a broad peak at 2.5 eV. Partial yield measurements [3] confirm the metallic behavior for this coverage range. The emission at E, becomes unambiguously evident at 3 ML.

3.2.2. Sn /Si(lOO) thermal treatment Fig. 4 shows the IPES spectra of four Sn/Si(lOO> reconstructions obtained after annealing. There is a clear evolution of the empty states as a function of tin coverage and for the different superstructures observed by LEED. Table 2 reports the energy positions of the peaks as detected by IPES in the different reconstructions. For the (2 x 1) superstructure that corresponds [l] to tin coverages 2 0.2 ML the IPES spectrum is quite similar to that of the Sn/Si(100)(2 X 1) grown at room temperature, with, at about 1.1 eV, the state more pronounced.

Table 1 Energy position of the structures as measured by IPES for Sn/Si(lOO) interfaces grown at room temperature; as detected by STM measurements (Ref. [2]) are also listed System

LEED

Dimer

Energy position of the structures by IPES feV)

Si(100) 0.4 ML

(2 x 1) ‘(2 x 1)

O-O.6

1.1 ML 2.3 ML 3ML

(1 x 1) Diffuse Diffuse

Si-Si Si-Si Sn-Sn _ _ _

The coverages are estimated by AES calibration.

3.0 2.9

1.1 1.1

2.8

0.8 2.5 2.5

the kinds of dimers

4.3

M. Pedio et al. /Surface

I”’

data. At 1.1 eV a prominent new state appears, together with a broad peak at about 2.5 eV. For superstructures corresponding to coverages below 1 ML there is no emission at the Fermi level. The metallic character of the interface is evident for the Sn/Si(lOOX5 x 11, corresponding to coverages between 1 and 1.5 ML.

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..J-

\.

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Sn/Si ~(8x4)

4. Discussion

0123456

E-EF (eV) Fig. 4. Angle-integrated Sn/Si(lOO) superstructures.

IPES

spectra

for

the

various

In the Sn/Si(lOOX6 x 2) IPES spectrum strong differences are detectable: at about 1.0 eV the state disappears, while at 1.8 eV a peak is present. The features in the spectrum appear slightly washed. The ~$8 >(:4) reconstruction measurements present some similarity with the (2 x 1) IPES

Tin and silicon belong both to the Group-IV elements, have similar electronegativity, are thermodynamically insoluble and do not form alloys [lo]. These factors lead to a very rich phase diagram for the Sn/Si(lOO> systems, with a large surface reconstruction variety. For the case of Sn/Si(lOO> interfaces grown at room temperature and with low tin coverages a (2 X 1) LEED pattern is still present. The 0.4 ML IPES spectrum of Fig. 3 is similar to that of clean Si(lOO), showing a broad emission at the bottom of the conduction band and a feature at 1.1 eV above the Fermi level (Table 1). At this coverage, STM measurements [2] find a locally ordered anisotropic growth of Sn dimers and still reveal the presence of Si-dimer dangling bonds. The existence of a Si-dimer dangling-bond region is confirmed by the persistence of the S, feature at 1.1 eV. The more blurred emission at the bottom of the conduction band could be explained by a more disordered configuration due likely to the presence of both Si-Si and Sn-Sn dimers.

Table 2 Energy position of the structures as measured by IPES for Sn/Si(lOO) interfaces for annealed interfaces; the kinds of dimers as detected by STM measurements (Ref. [2]) are also listed System

LEED

Dimer

Energy position of the structures by IPES (eV1

SK1001 0.2 ML

(2X 1) (2x 1)

O-O.6

0.3-0.5 ML

(6 X 2)

0.5-1.0 ML 1.0-I/5 ML

~(8 x 4) (5 x 1)

Si-Si Si-Si Sn-Sn Mixed Sn-Si Sn-Sn

1.1 1.0

3.0 2.5

1.8

2.6

1.1

2.5 2.5

The coverages are estimated by AES calibration and correspond to the sharper respective LEED patterns.

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spectroscopy measurement findings [3]. The energy position of the SnNVV AES transition is 17.9 eV and the MNN transition shows a new feature at 407 eV, typical of a metal-like behavior; the IPES spectra for 13~”> 2 ML are in good agreement with the theoretical prediction on p-Sn

1121.

Fig. 5. Atomic buckled Sn-Sn dimer model of the Sn/Si(lOO) ~(8x4) interfaces as proposed by STM (Ref. [2]). Each Sn dimer consists of a pair of gray and solid circles. The solid circle represents the Sn atom buckled out of the surface plane.

For the 1.1 ML coverage the LEED pattern changes from (2 x 1) to (1 X 1). This is a surface reconstruction that is never observed during our experiment after the annealing procedures (see Figs. 4 and 5). This observation shows that the adlayer is strongly disordered, with the (1 X 1) LEED pattern arising from the substrate underneath. In fact an ideal Sn (1 X 1) LEED pattern for coverages above 1 ML is unfavorable, due to large surface strains induced by the exceedingly high lattice mismatch between tin and silicon (> 10%). The contraction of the unit cell volume in the Sn adlayer corresponds to the formation of an adlayer with metallic character. On the other hand, the SnNVV Auger line of Fig. 2 shifts toward higher kinetic energies and the SnMNN transition shows the appearance of a new structure at 407 eV. These findings suggest that in this coverage range the interface is developing from a semiconducting to a metal-like phase. For tin coverages up to 1 ML the states found theoretically for the (w-SN [ll] do not fit with our IPES results and growth of a-Sn on Si(100) can be excluded. The presence of a tail at E,, for tin coverages at 2.3 ML and above, suggests that Sn-Sn interaction is becoming strong and induces a metallic overlayer, in agreement with the partial yield

In case of annealed growth of Sn on Si(lOO), the IPES spectra show a well-defined dependence on the interface geometry and tin coverage. At low coverages (0.2 ML), the LEED still shows a (2 x 1) pattern and the IPES spectrum is similar to that of the Sn/Si(100)(2 x 1) system grown at room temperature (Fig. 3, Tables 1 and 2). In both cases the S, state at about 1.1 eV is clearly detectable and we can reach the same conclusions, i.e., the Si-Si-dimer dangling-bond empty states are not perturbed very much by the presence of Sn atoms, at this stage. For higher coverages the empty states’ features as detected by IPES are clearly different from the Sn/Si(100>(2 X 1) case (Fig. 4, Table 2). The Sn/Si(100>(6 X 2) IPES spectrum clearly shows a unique feature at 1.8 eV, not present in all other reconstructions. This structure is reproducible and can be considered a fingerprint of the (6 x 2) surface and attributed to Sn-Si dimers’ presence, as has been observed by STM analysis

Dl. In the Sn/Si(lOO>c(8 x 4) IPES spectrum a clear and pronounced state is present at about 1.1 eV above the Fermi level. A model proposed in the STM study of Ref. [2] shows that in this coverage range (corresponding to about 1 ML) the configuration of the Sn layer consists of Sn-Sn asymmetric dimers, forming chains. Fig. 5 shows the model based on the STM results, for clarity. Similar results have been obtained by an STM study [13] on the Sn/GaAs(llO) system, with a similar chain structure of the overlayer and a clear empty state at an energy below 1.2 eV. Theoretical calculations assign this state to the Sn-Sn empty interacting dangling bonds. The same behavior is also present in the Sb/Si(lOO) [14] interface, where Sb-Sb dimers are present in the (2 x 1) reconstruction and the angular-resolved IPES spectra show a prominent

M. Pedio etal./Surface

peak at about 1 eV. This feature is absent in the Sb/Si(lOO)(l X 1) reconstruction, where dimers are not present. In conclusion, we attribute the 1.1 eV peak of the Sn/Si(lOO>c(8 X 4) of Fig. 4 to chains of Sn-Sn adlayer dimers. The (5 x 1) reconstruction IPES spectrum in Fig. 4 shows metallic character and appears to be similar to the 2.3 ML Sn coverage spectrum (Fig. 3). The (5 x 1) reconstruction represents a significant contraction of the unit cell size, i.e., a configuration that makes p-Sn more favorable, as already noted by different authors [l-3]. The above analysis clearly indicates that the semiconductor-metal transition of the Sn/Si(lOO) system takes place above 1 ML of tin coverage. We want to stress that the LEED patterns taken under corresponding conditions as the IPES spectra reported in Fig. 4 were very clear and sharp. Moreover, we obtained, in different experimental runs with intermediate tin coverages, mixed LEED patterns. The corresponding IPES data can help to follow the electronic properties evolution already discussed. We report the IPES data corresponding to the mixed (6 x 2)-~$8 X 4) and c(8 x 4)-(5 x 1) phases in Fig. 6. The SnSn-dimers-induced state of the c(8 X 4) configuration at 1.1 eV is already present in the mixed (6 X 2)-c(8 X 4) IPES spectrum. This last spectrum was obtained with 0.5 ML tin coverage. The l.l-eV feature is disappearing in the c(8 X 4)-(5 x 1) phase and it is not present for the (5 x 1) phase. From the data of Figs. 4 and 6 we have also the indication that in the ~$8 X 4) configuration the bottom of the conduction band is reaching the Fermi level. This conclusion is confirmed by a recent direct photoemission analysis [ 11, where the separation between the Fermi level and the valence band maximum goes to zero for the (5 x 1) reconstruction. In the direct photoemission results, reported in Ref. [l], the separation between the valence band maximum and the Fermi level decreases for coverages up to 0.3 ML, then it presents a plateau in the range of 0.3-0.5 ML, corresponding to the (6 x 2) reconstruction, and subsequently decreases again in correspondence of the c(8 X 4).

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E-EF (eV) Fig. 6. Angle-integrated mixed superstructures.

IPES

spectra

for

the

Sn/Si(lOO)

Inverse photoemission and direct photoemission results suggest that the Sn/Si(lOO)c(8 x 4) interface would correspond to a structural and electronic phase with Egap < 0.5 eV. This phase is still semiconducting but is different from the a-Sn. This last conclusion was inferred in Ref. [l] by UPS data; the ~$8 x 4) valence band has indeed a very different shape than that observed for the cu-Sn [ll]. A recent theoretical study has shown that the energy-gap evolution for ordered Sn/Si(lll> systems is a function of the number of tin atoms in the unit surface cell [153. Our study gives an indication about the feasibility of Sn-Si systems to form interfaces with small and adjustable energy gaps. Further studies are necessary to demonstrate experimentally the formation of layered Si-Sn superlattices that maintain the elec-

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tronic face.

properties

found,

e.g., in the c(8 x 4) inter-

5. Conclusions We have discussed the evolution of the empty electronic states of the epitaxial Sn/Si(lOO) interfaces both grown at room temperature and after annealing. Angle-integrated IPES spectroscopy results prove extremely sensitive to the different reconstructions of the system including the cases where a mixed configuration has been found. The IPES study of Sn/Si(lOO) interfaces at room temperature has shown that the tin overlayer is metallic for coverages > 1 ML. For the annealed interfaces it is possible to detect a clear evolution of the empty states, characteristic of the various formed superstructures. In particular for the ~(8 X 4) the phase is semiconducting with a small Egap < 0.5 eV. In analogy with other similar systems the origin of the pronounced states at about 1 eV above the Fermi level has been assigned, with the help of STM studies [3], to Sn-Sn asymmetric dimers. This phase is semiconducting but is different from the a-Sn phase. It is typical of 1 ML Sn grown on substrates having a strong lattice mismatch (> 10%) with the a-Sn. We suggest that the peak at about 1 eV above E, in the IPES spectra is a sort of fingerprint related to the existence of dimers at the interface. The Sn/Si(100)(5 X 1) superstructure, corresponding to tin coverages above 1 ML, has a

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metallic character, toemission results

in agreement [2].

with direct pho-

6. References [ll D.H. Rich, T. Miller, A. Samsavar,

H.F. Lin and T.C. Chiang, Phys. Rev. B 37 (1988) 10221. I21 A.A. Baski, CF. Quate and J. Nogami, Phys. Rev. B 44 (1991) 11167. J.P. Lacharme and C.A. Sebenne, 131I. Adriamantenasoa, Surf. Sci. 189/190 (1987) 563. [41 K. Ueda and K. Kinoshita, Surf. Sci. 145 (1984) 261. Spectroscopy and Photoelectron l51 S. Kono, Photoelectron Diffraction Studies of Submonolayer Metal/Sic1 11) Interfaces, in: Core-Level Spectroscopy in Condensed Systems, Eds. J. Kotani and A. Kotani (Springer, Berlin, 1987) p. 253, and references therein. Appl. Phys. Lett. 54 l61 J.L. Reno and L.L. Stephenson. (1989) 2207; M.T. Asom, A.R. Kortan, L.C. Kimerling and R.C. Farrow, Appl. Phys. Lett. 55 (1989) 1439. Surf. Sci. 126 171 H. Htichst and I. Hernandez-Calderon, (1983) 2s; H.U. Middelmann. L. Sorba, V. Hinkel and K. Horn, Phys. Rev. B 35 (1987) 718. U. de1 Pennino, S. Sassaroli and S. b31 0. Bisi, C. Calandra, Valeri, Phys. Rev. B 30 (1984) 5696. and B. Reihl, Surf. Sci. 269/270 (1992) [91 L.S.O. Johansson 810. of Binary Alloys (McGraw-Hill, [lOI M. Hansen, Constitution New York, 1958). [ill S. Groves and W. Paul, Phys. Rev. Lett. 11 (1963) 194. [I21 G. Weisz, Phys. Rev. 149 (1966) 504. and K.C. Pandey, [131 C.K. Shih, E. Kaxiras, R.M. Feenstra Phys. Rev. B 40 (1989) 10044. [141 A. Cricenti, private communication. and N.W. Ashcroft, J. Phys. Condensed [I51 A.P. Horsfield Matter 4 (1992) 7333.