+Using alternating supply of the gas beams Si$_2$H$_6$ and C$_2$H$_2$ in GSMBE, 3C-SiC epilayers were obtained on 6H-SiC substrates at temperatures between \unit[850]{$^{\circ}$C} and \unit[1000]{$^{\circ}$C} \cite{yoshinobu92}.
+On \hkl(000-1) substrates twinned \hkl(-1-1-1) oriented 3C-SiC domains are observed, which suggest a nucleation driven rather than step-flow growth mechanism.
+On \hkl(0-11-4) substrates, however, single crystalline \hkl(001) oriented 3C-SiC grows with the c axes of substrate and film being equal.
+The beneficial epitaxial relation of substrate and film limits the structural difference between the two polytypes in two out of six layers with respect to the stacking sequence along the c axis.
+Homoepitaxial growth of 3C-SiC by GSMBE was realized for the first time by atomic layer epitaxy (ALE) utilizing the periodical change in the surface superstructure by the alternating supply of the source gases, which determines the growth rate giving atomic level control in the growth process \cite{fuyuki89}.
+The cleaned substrate surface shows a C terminated $(2\times 2)$ pattern at \unit[1000]{$^{\circ}$C}, which turns into a $(3\times 2)$ pattern when Si$_2$H$_6$ is introduced and it is maintained after the supply is stopped.
+A more detailed investigation showed the formation of a preceeding $(2\times 1)$ and $(5\times 2)$ pattern within the exposure to the Si containing gas \cite{yoshinobu90,fuyuki93}.
+The $(3\times 2)$ superstructure contains approximately 1.7 monolayers of Si atoms, crystallizing into 3C-SiC with a smooth and mirror-like surface after C$_2$H$_6$ is inserted accompanied by a reconstruction of the surface into the initial C terminated $(2\times 2)$ pattern.
+A minimal growth rate of 2.3 monolayers per cycle exceeding the value of 1.7 is due to physically adsorbed Si atoms not contributing to the superstructure.
+To realize single monolayer growth precise control of the gas supply to form the $(2\times 1)$ structure is required.
+However, accurate layer-by-layer growth is achieved under certain conditions, which facilitate the spontaneous desorption of an additional layer of one atom species by supply of the other species \cite{hara93}.
+Homoepitaxial growth of the 6H polytype has been realized on off-oriented substrates utilizing simultaneous supply of the source gases \cite{tanaka94}.
+Depending on the gas flow ratio either 3C island formation or step flow growth of the 6H polytype occurs, which is explained by a model including aspects of enhanced surface mobilities of adatoms on a $(3\times 3)$ reconstructed surface.
+Due to the strong adsorption of atomic hydrogen \cite{allendorf91} decomposited of the gas phase reactants at low temperatures, however, there seems to be no benefit of GSMBE compared to CVD.
+Next to lattice imperfections, incorporated hydrogen effects the surface mobility of the adsorbed species \cite{eaglesham93} setting a minimum limit for the growth temperature, which would preferably be further decreased in order to obtain sharp doping profiles.
+Thus, growth rates must be adjusted to be lower than the desorption rate of hydrogen, which leads to very low deposition rates at low temperatures.
+SSMBE, by supplying the atomic species to be deposited by evaporation of a solid, presumably constitutes the preffered method in order to avoid the problems mentioned above.
+Although, in the first experiments, temperatures still above \unit[1100]{$^{\circ}$C} were necessary to epitaxially grow 3C-SiC films on 6H-SiC substrates \cite{kaneda87}, subsequent attempts succeeded in growing mixtures of twinned 3C-SiC and 6H-SiC films on off-axis \hkl(0001) 6H-SiC wafers at temperatures between \unit[800]{$^{\circ}$C} and \unit[1000]{$^{\circ}$C} \cite{fissel95,fissel95_apl}.
+In the latter approach, as in GSMBE, excess Si atoms, which are controlled by the Si/C flux ratio, result in the formation of a Si adlayer and the formation of a non-stoichiometric, reconstructed surface superstructure, which influences the mobility of adatoms and, thus, has a decisive influence on the growth mode, polytype and crystallinity \cite{fissel95,fissel96,righi03}.
+Therefore, carefully controlling the Si/C ratio could be exploited to obtain definite heterostructures of different SiC polytypes providing the possibility for band gap engineering in SiC materials.
+
+To summarize, much progress has been achieved in SiC thin film growth during the last few years.
+However, the frequent occurence of defects such as dislocations, twins and double positioning boundaries limit the structural and electrical quality of large SiC films.
+Solving this issue remains a challenging problem necessary to drive SiC for potential applications in high-performance electronic device production \cite{wesch96}.
+
+\subsection{Ion beam synthesis of cubic silicon carbide}
+\label{subsection:ibs}
+
+Although tremendous progress has been achieved in the above-mentioned growth methods during the last decades, available wafer dimensions and crystal qualities are not yet statisfactory.
+Thus, alternative approaches to fabricate SiC have been explored.
+The ion beam synthesis (IBS) technique, i.e. high-dose ion implantation followed by a high-temperature annealing step, turned out to constitute a promising method to directly form compound layers of high purity and accurately controllable depth and stoichiometry.
+A short chronological summary of the IBS of SiC and its origins is presented in the following.
+
+High-dose carbon implantation into crystalline silicon (c-Si) with subsequent or in situ annealing was found to result in SiC microcrystallites in Si \cite{borders71}.
+Rutherford backscattering spectrometry (RBS) and infrared (IR) spectroscopy investigations indicate a \unit[10]{at.\%} C concentration peak and the occurence of disordered C-Si bonds after implantation at room temperature (RT) followed by crystallization into SiC precipitates upon annealing demonstrated by a shift in the IR absorption band and the disappearance of the C profile peak in RBS.
+Implantations at different temperatures revealed a strong influence of the implantation temperature on the compound structure \cite{edelman76}.
+Temperatures below \unit[500]{$^{\circ}$C} result in amorphous layers, which is transformed into polycrystalline 3C-SiC after \unit[850]{$^{\circ}$C} annealing.
+Otherwise single crystalline 3C-SiC is observed for temperatures above \unit[600]{$^{\circ}$C}.
+Annealing temperatures necessary for the onset of the amorphous to crystalline transition have been confirmed by further studies \cite{kimura81,kimura82}.
+Overstoichiometric doses result in the formation of clusters of C, which do not contribute to SiC formation during annealing up to \unit[1200]{$^{\circ}$C} \cite{kimura82}.
+The amount of formed SiC, however, increases with increasing implantation temperature.
+The authors, thus, concluded that implantations at elevated temperatures lead to a reduction in the annealing temperatures required for the synthesis of homogeneous layers of SiC.
+In a comparative study of O, N and C implantation into Si, the absence of the formation of a stoichiometric SiC compound layer involving the transition of a Gaussian into a box-like C profile with respect to the implantation depth for the superstoichiometric C implantation and an annealing temeprature of \unit[1200]{$^{\circ}$C} in contrast to the O and N implantations, which successfully form homogeneous layers, has been observed \cite{reeson86}.
+This was attrubuted to the difference in the enthalpy of formation of the respective compound and the different mobility of the respective impurity in bulk Si.
+Thus, higher annealing temperatures and longer annealing times were considered necessary for the formation of homogeneous SiC layers.
+Indeed, for the first time, buried homogeneous and stoichiometric epitaxial 3C-SiC layers embedded in single crystalline Si were obtained by the same group consequently applying annealing temperatures of \unit[1405]{$^{\circ}$C} for \unit[90]{min} and implantation temperatures of approximately \unit[550]{$^{\circ}$C} \cite{reeson87}.
+The necessity of the applied extreme temperature and time scale is attributed to the stability of substitutional C within the Si matrix being responsible for high activation energies necessary to dissolve such precipitates and, thus, allow for redistribution of the implanted C atoms.
+In order to avoid extreme annealing temperatures close to the melting temperature of Si, triple-energy implantations in the range from \unit[180-190]{keV} with stoichiometric doses at a constant target temperature of \unit[860]{$^{\circ}$C} achieved by external substrate heating were performed \cite{martin90}.
+It was shown that a thick buried layer of SiC is directly formed during implantation, which consists of small, only slightly misorientated but severely twinned 3C-SiC crystallites.
+The authors assumed that due to the auxiliary heating rather than ion beam heating as employed in all the preceding studies, the complexity of the remaining defects in the synthesized structure is fairly reduced.
+Even better qualities by direct synthesis were obtained for implantations at \unit[950]{$^{\circ}$C} \cite{nejim95}.
+Since no amorphous or polycrystalline regions have been identified, twinning is considered to constitute the main limiting factor in the IBS of SiC.
+
+Further studies revealed the possibility to form buried layers of SiC by IBS at moderate substrate and anneal temperatures \cite{lindner95,lindner96}.
+Different doses of C ions with an energy of \unit[180]{keV} were implanted at \unit[330-440]{$^{\circ}$C} and annealed at \unit[1200]{$^{\circ}$C} or \unit[1250]{$^{\circ}$C} for \unit[5-10]{h}.
+For a critical dose, which was found to depend on the Si substrate orientation, the formation of a stoichiometric buried layer of SiC exhibiting a well-defined interface to the Si host matrix was observed.
+In case of overstoichiometric C concentrations the excess C is not redistributed.
+These investigations demonstrate the presence of an upper dose limit, which corresponds to a \unit[53]{at.\%} C concentration at the implantation peak, for the thermally induced redistribution of the C atoms from a Gaussian to a box-shaped depth profile upon annealing.
+This is explained by the formation of strong graphitic C-C bonds for higher C concentrations \cite{calcagno96}.
+Increased temperatures exceeding the Si melting point are expected to be necessary for the dissociation of these C clusters.
+Furthermore, higher implantation energies were found to result in layers of variable composition exhibiting randomly distributed SiC precipitates.
+In another study \cite{serre95} high dose C implantations were performed at room temperature and \unit[500]{$^{\circ}$C} respectively.
+Implantations at room temperature lead to the formation of a buried amorphous carbide layer in addition to a thin C-rich film at the surface, which is attributed to the migration of C atoms towards the surface.
+In contrast, implantations at elevated temperatures result in the exclusive formation of a buried layer consisting of 3C-SiC precipitates epitaxially aligned to the Si host, which obviously is more favorable than the C migration towards the surface.
+Annealing at temperatures up to \unit[1150]{$^{\circ}$C} does not alter the C profile.
+Instead defect annihilation is observed and the C-rich surface layer of the room temperature implant turns into a layer consisting of SiC precipitates, which, however, are not aligned with the Si matrix indicating a mechanism different to the one of the direct formation for the high-temperature implantation.
+
+Based on these findings, a recipe was developed to form buried layers of single-crystalline SiC featuring an improved interface and crystallinity \cite{lindner99,lindner01,lindner02}.
+Therefore, the dose must not exceed the stoichiometry dose, i.e. the dose corresponding to \unit[50]{at.\%} C concentration at the implantation peak.
+Otherwise clusters of C are formed, which cannot be dissolved during post-implantation annealing at moderate temperatures below the Si melting point \cite{lindner96,calcagno96}.
+Annealing should be performed for \unit[5-10]{h} at \unit[1250]{$^{\circ}$C} to enable the redistribution from the as-implanted Gaussian into a box-like C depth profile \cite{lindner95}.
+The implantation temperature constitutes the most critical parameter, which is responsible for the structure after implantation and, thus, the starting point for subsequent annealing steps.
+Implantations at \unit[400]{$^{\circ}$C} resulted in buried layers of SiC subdivided into a polycrystalline upper and an epitaxial lower part.
+This corresponds to the region of randomly oriented SiC crystallites and epitaxially aligned precipitates surrounded by thin amorphous layers without crystalline SiC inclusions in the as-implanted state.
+However, an abrupt interface to the Si host is observed after annealing.
+As expected, single-crystalline layers were achieved for an increased temperature of \unit[600]{$^{\circ}$C}.
+However, these layers show an extremely poor interface to the Si top layer governed by a high density of SiC precipitates, which are not affected in the C redistribution during annealing and, thus, responsible for the rough interface.
+Hence, to obtain sharp interfaces and single-crystalline SiC layers temperatures between \unit[400]{$^{\circ}$C} and \unit[600]{$^{\circ}$C} have to be used.
+Indeed, reasonable results were obtained at \unit[500]{$^{\circ}$C} \cite{lindner98} and even better interfaces were observed for \unit[450]{$^{\circ}$C} \cite{lindner99_2}.
+To further improve the interface quality and crystallinity a two-temperature implantation technique was developed \cite{lindner99}.
+To form a narrow, box-like density profile of oriented SiC nanocrystals \unit[93]{\%} of the total dose of \unit[$8.5\cdot 10^{17}$]{cm$^{-2}$} is implanted at \unit[500]{$^{\circ}$C}.
+The remaining dose is implanted at \unit[250]{$^{\circ}$C}, which leads to the formation of amorphous zones above and below the SiC precipitate layer and the desctruction of SiC nanocrystals within these zones.
+After annealing for \unit[10]{h} at \unit[1250]{$^{\circ}$C} a homogeneous, stoichiometric SiC layer with sharp interfaces is formed.
+Fig. \ref{fig:sic:hrem_sharp} shows the respective high resolution transmission electron microscopy micrographs.
+\begin{figure}[t]
+\begin{center}
+\includegraphics[width=0.6\columnwidth]{ibs_3c-sic.eps}
+\end{center}
+\caption[Bright field (a) and \hkl(1 1 1) SiC dark field (b) cross-sectional TEM micrographs of the buried SiC layer in Si created by the two-temperature implantation technique and subsequent annealing.]{Bright field (a) and \hkl(1 1 1) SiC dark field (b) cross-sectional TEM micrographs of the buried SiC layer in Si created by the two-temperature implantation technique and subsequent annealing as explained in the text \cite{lindner99_2}. The inset shows a selected area diffraction pattern of the buried layer.}
+\label{fig:sic:hrem_sharp}
+\end{figure}
+
+To summarize, by understanding some basic processes, IBS nowadays has become a promising method to form thin SiC layers of high quality exclusively of the 3C polytype embedded in and epitaxially aligned to the Si host featuring a sharp interface.
+Due to the high areal homogeneity achieved in IBS, the size of the layers is only limited by the width of the beam-scanning equipment used in the implantation system as opposed to deposition techniques, which have to deal with severe wafer bending.
+This enables the synthesis of large area SiC films.