Self-organized SiC nanostructures on silicon

Superlattices and Microstructures 36 (2004) 345–351 www.elsevier.com/locate/superlattices

Self-organized SiC nanostructures on silicon V. Cimallaa,∗, A.A. Schmidtb, Ch. Foerstera, K. Zekentesc, O. Ambachera, J. Pezoldta a Center for Micro- and Nanotechnologies, Technical University Ilmenau, 98693 Ilmenau, Germany b A.F. Ioffe Physico-Technical Institute of RAS, 26 Polytechnicheskaya Street, 19402 St. Petersburg, Russia c Foundation of Research and Technology-Hellas, P.O. Box 1527, Heraklion, Crete, 71110, Greece

Available online 18 September 2004

Abstract Self-organization in a system with chemical interactions between the substrate (silicon) and deposit (carbon) is demonstrated by growing SiC dots on silicon substrates. The large lattice mismatch between silicon and SiC of 20% stimulates a three-dimensional nucleation on the substrate. This spontaneous formation of islands is a powerful tool for the formation of dots. However, the chemical interaction leads to an instability of the Si surface during the nucleation and the growth: the need for Si for the SiC formation as well as the Si evaporation results in a depletion of the area surrounding the SiC islands. As a result, well-defined pyramids with a fourfold symmetry are formed on Si(001) substrates with SiC nuclei on the top. The nucleation sites were controlled by the formation of equally spaced monatomic and biatomic steps on Si(001) and Si(111), respectively. The resulting terraces promote an alignment of the SiC dots along the step edges. By applying atomic force microscopy we demonstrate a lateral ordering of SiC dots in linear chains and in dense dot arrays. Depending on the process conditions, the SiC dot separation was adjusted between 20 and 500 nm. © 2004 Elsevier Ltd. All rights reserved.

∗ Corresponding address: Zentrum für Mikro- und Nanotechnologien, Fachgebiet für Nanotechnologie, TU Ilmenau, Kirchhoffstrasse 7, 98693 Ilmenau, Germany. Tel.: +49 3677 69 3408; fax: +49 3677 69 3355. E-mail address: [email protected] (V. Cimalla).

0749-6036/$ - see front matter © 2004 Elsevier Ltd. All rights reserved. doi:10.1016/j.spmi.2004.08.001

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1. Introduction Silicon carbide is a versatile material with a wide range of applications. The outstanding physical and chemical properties of SiC, especially its wide band gap, make this material suitable for different applications in the field of high frequency, high power devices for operation at high temperatures and in harsh environments. Up to now, less attention has been drawn to the fabrication and application of this material in nanoscale device structures. SiC in and on silicon, if scaled down to quantum dot dimensions, can be applied as a tunneling barrier in various nanoheterostructure and tunneling devices [1–3]. Other possible application fields are those of antidot structures for lateral conductivity modulations of two-dimensional electron gases and artificial anisotropies with applications in narrow band photodetectors and magnetic frequency dependent switches. In the field of optics and optoelectronics, nanoscale SiC can be used as an emitter in the UV [4] and far infrared regions [5]. Regardless of the bright prospects for applications, the formation of nanometric sized SiC nuclei and their regularities have been studied only to a minor extent [6–13]. First attempts to model the processes occurring during the nucleation and growth of nanoscale SiC clusters are published in [6,14–16]. To our knowledge, only in [17] have attempts been made to fabricate and characterize the growth of ordered SiC dots on silicon surfaces. In a previous paper we demonstrated the first results on self-ordering of SiC nanoscale crystallites on silicon surfaces [17]. In this paper the focus will be on the physical mechanisms which lie behind the observed ordering phenomena. 2. Experimental The experiments were accomplished in two different SSMBE configurations with electron beam evaporators for Si and C. The growth on Si(111) was investigated in a UMS 500 Balzers system and an RCA cleaning with a finishing HF dip was used before a subsequent annealing was carried out in the MBE chamber. For the growth on Si(100) a VG80S system was used, and the substrates were cleaned chemically with a subsequent in situ removal of the protective oxide. The reproducibility of the processes was checked by performing experiments with representative conditions in both growth chambers. The substrate temperature and carbon flux in both chambers were varied between 600 and 1050 ◦ C and between 1013 and 1015 atoms cm−2 s−1 , respectively. The growth processes were monitored in real time by means of RHEED in both SSMBE chambers, and additionally by means of ellipsometry in the case of Si(111). The layers grown were further analyzed ex situ by atomic force microscopy (AFM) and scanning electron microscopy (SEM). 3. Results and discussion The first stage of the silicon carbon interactions consists in a carbon induced change of the silicon surface reconstruction [11]. Within the period of time where the phase

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transitions of the surface reconstructions occur, carbon will be incorporated into the near surface layers [18]. Thus a Si1−x Cx alloy is formed. Subsequent to alloying, the first 3CSiC nuclei are formed [18,19]. The SiC formation and growth during the carbon interaction with the silicon surface (carbonization) could be divided roughly into two temperature regions depending on the carbon flux [20]. At substrate temperatures below 800 ◦ C the interaction process leading to SiC formation is similar to the “complete condensation” process and the Si substrate surface morphology is preserved during the initial stages of the silicon carbide formation [20]. Under these nucleation conditions a pseudo-twodimensional growth of SiC occurs. This growth mode results in an early coverage of the silicon surface. The dimensions of the SiC nuclei are below 30 nm. At substrate temperatures above 800 ◦ C re-evaporation of the deposited carbon species, sublimation of the silicon surface and silicon surface diffusion starts to degrade the “complete condensation” conditions. This leads to a SiC growth behavior similar to “incomplete condensation”, i.e. pronounced 3D nucleation [20]. The SiC islands appear very pronounced on the AFM pictures and are well separated (Fig. 1). The AFM pictures obtained were analyzed by using power density spectral analysis. From the obtained spectra, the density of the islands and their periodicity were extracted. The temperature and time dependences of the SiC island density in the cases of Si(100) and Si(111) surfaces are shown in Figs. 2 and 3, respectively. In agreement with the classical nucleation theory [21] the island density increases with increasing temperature and obeys an Arrhenius law. The time dependence for the SiC islands shows three distinct regions (Fig. 3). The first region is the transient nucleation region, i.e. nonsteady state nucleation density. In this region the density of the SiC nuclei increases with increasing process time. This behavior was observed for both silicon surfaces and extends approximately up to an effective carbon coverage of 10% of the silicon surface. At higher coverages, i.e. growth times, the observed densities remain constant, i.e. the equilibrium density is reached. At very large growth times, i.e. at a carbon coverage above 3 ML with respect to the silicon surface, a decrease of the nuclei density due to coalescence effects was observed (see also Fig. 1). It is noteworthy that the nucleation density on Si(111) surfaces is higher than on Si(100) ones. The critical nuclei size is given from the classical nucleation theory by the following expression [21]: rcrit =

−2γ G v

(1)

where G v is the free energy change per unit volume when the cluster is formed by condensation from the vapour phase and γ is the surface free energy of the SiC cluster formed. If we take into account that SiC(100) surfaces exhibit a higher surface energy compared to SiC(111), the density of the islands on Si(100) has to be lower compared to that of Si(111) surfaces, if the other growth parameters are held constant during processing. A similar observation has been reported in [22], where a higher nucleation density during the initial stages of epitaxial growth of 3C-SiC on Si(100) and Si(111) substrates was obtained by using hexamethylsilane in a hydrogen–argon carrier during chemical vapour deposition. For both surfaces, at certain process conditions a two-dimensional ordering of the SiC dots grown was observed (Fig. 4). To clarify the reason for the ordering effect, a detailed

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Fig. 1. AFM images of the 3C-SiC nucleation on Si(100) ((a), (c), (e)) and Si(111) ((b), (d), (f)) at different growth times: ((a), (b)) for the initial nucleation at 3 s, ((c), (d)) after reaching the equilibrium nucleation density and ((e), (f)) after the coalescence of the SiC islands.

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Fig. 2. SiC island density versus substrate temperature on Si(100) and Si(111) at carbon fluxes of 6 × 1013 cm−2 s−1 and 3 × 1013 cm−2 s−1 , respectively, for a growth time of 180 s.

analysis of the surface morphology was carried out. The analysis revealed that the SiC dots were located at the step bands formed on the silicon substrate [17]. Similarly, the initial introduction of carbon into the silicon surface at step edges was proved by STM in [23]. The density of the SiC nuclei in the case of ordering was in agreement with the nucleation kinetics shown in Figs. 2 and 3. A correlation of the mean distance between the SiC dots L D and the measured terrace width on the silicon surface L T revealed that the ordering occurs only in the case where L T is in the region of L D —therefore, when only one SiC nucleus could be formed on the Si terrace. In the case of L D  L T SiC nuclei did not appear on every terrace, whereas for L D  L T a random nucleation was initiated. So, for controlling the ordering of the SiC dots on Si surfaces the values of L D have to be correlated with L T as regards order. This can be done by using two approaches. The first approach is based on the dependence of the nucleation density on the substrate temperature, supersaturation, and surface and interface energies of the heteroepitaxial system. These parameters allow one to adjust the mean distance between the SiC dots to that of the step bands formed on the Si surface by substrate preparation. In this approach it is important to note that the steady state dot density, i.e. the equilibrium density of the nuclei, has to be reached to obtain the best ordering effect. This is because during the nonstationary nucleation phase at the beginning of the SiC growth smaller nuclei are evident, whereas in the coalescence stage larger features appear on the silicon surface disturbing the narrow dot distribution. The second approach consists in the adjustment of the terrace width to the

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Fig. 3. The time dependent evolution of the SiC island density at 925 ◦ C on Si(100) and Si(111) and carbon fluxes of 6 × 1013 cm−2 s−1 and 3 × 1013 cm−2 s−1 , respectively.

Fig. 4. Examples of the formation of two-dimensional ordered arrays of SiC dots on Si(100) for complete condensation and Si(111) surfaces for incomplete condensation.

nucleation density. For this reason the off-cut angle of the silicon substrates used and the sample preparation conditions leading to different step bunching configurations have to be tuned to a given nucleation condition.

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For both approaches another restriction has to be taken into account. This requirement comes from the need for preservation of the morphology of the step band. With increasing growth time, they can lose their geometry due to the Si consumption necessary to form SiC, because only carbon atoms are supplied to the silicon surface. The source for this Si supply are the step edges where the Si atoms have a lower binding energy as compared to the atoms of the terraces. An example of an appropriate adjustment of the nucleation density to a step band configuration is shown in Fig. 4. 4. Conclusion On the basis of the analysis of the nucleation of SiC dots on Si(100) and Si(111) surfaces, rules for the formation of ordered dot arrays for chemically interacting systems were worked out and demonstrated for the case of SiC nucleation on silicon. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23]

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