A more detailed investigation showed the formation of a preceding $(2\times 1)$ 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.
+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.
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 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 destruction 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.
+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}
\subfigure[]{\label{fig:sic:db_agglom:seq03}\includegraphics[width=0.30\columnwidth]{sic_prec_seq_03.eps}}
%SiC formation and release of excess Si atoms
\end{center}
-\caption[Two dimensional schematic of the assumed SiC precipitation mechanism based on an initial C-Si dumbbell agglomeration.]{Two dimensional schematic of the assumed SiC precipitation mechanism based on an initial C-Si dumbbell agglomeration. C atoms (red dots) incorporated into the Si (black dots) host form C-Si dumbbells (a), which agglomerate into clusters (b) followed by the precipitation of SiC and the emission of a few excess Si atoms (black circles), which are located in the interstitial Si lattice (c). The dotted lines mark the atomic spacing of c-Si in \hkl[1 0 0] direction indicating the $4/5$ ratio of the lattice constants of c-Si and 3C-SiC.}
+\caption[Two dimensional schematic of the assumed SiC precipitation mechanism based on an initial C-Si dumbbell agglomeration.]{Two dimensional schematic of the assumed SiC precipitation mechanism based on an initial C-Si dumbbell agglomeration. C atoms (red dots) incorporated into the Si (black dots) host form C-Si dumbbells (a), which agglomerate into clusters (b) followed by precipitation of SiC and the emission of a few excess Si atoms (black circles), which are located in the interstitial Si lattice (c). The dotted lines mark the atomic spacing of c-Si in \hkl[1 0 0] direction indicating the $4/5$ ratio of the lattice constants of c-Si and 3C-SiC.}
\label{fig:sic:db_agglom}
\end{figure}
The incorporated C atoms form C-Si dumbbells on regular Si lattice sites.
-With increasing dose and proceeding time the highly mobile dumbbells agglomerate into large clusters.
+With increasing dose and proceeding time, the highly mobile dumbbells agglomerate into large clusters.
Finally, when the cluster size reaches a critical radius, the high interfacial energy due to the 3C-SiC/c-Si lattice misfit is overcome and precipitation occurs.
Due to the slightly lower silicon density of 3C-SiC, excessive silicon atoms exist, which will most probably end up as self-interstitials in the c-Si matrix since there is more space than in 3C-SiC.
However, the exact atomic rearrangement involved within this topotactic transformation is not identified.
Furthermore, IBS studies of Reeson~et~al.~\cite{reeson87}, in which implantation temperatures of \unit[500]{$^{\circ}$C} were employed, revealed the necessity of extreme annealing temperatures for C redistribution, which is assumed to result from the stability of substitutional C and consequently high activation energies required for precipitate dissolution.
The results support a mechanism of an initial coherent precipitation based on substitutional C that is likewise valid for the IBS of 3C-SiC by C implantation into Si at elevated temperatures.
-The fact that the metastable cubic phase instead of the thermodynamically more favorable hexagonal $\alpha$-SiC structure is formed and the alignment of these cubic precipitates within the Si matrix, which can be explained by considering a topotactic transformation by C atoms occupying substitutionally Si lattice sites of one of the two fcc lattices that make up the Si crystal, reinforce the proposed mechanism.
+The fact that the metastable cubic phase instead of the thermodynamically more favorable hexagonal $\alpha$-SiC structure is formed and the alignment of these cubic precipitates within the Si matrix, which can be explained straightforward by considering a topotactic transformation by C atoms occupying substitutionally Si lattice sites of one of the two fcc lattices that make up the Si crystal, reinforce the proposed mechanism.
-To conclude, a controversy with respect to the precipitation of SiC in Si exists in literature.
+To conclude, a controversy with respect to the precipitation model of SiC in Si exists in literature.
Next to the pure scientific interest, solving this controversy and gaining new insight in the SiC conversion mechanism might enable significant progress in the heteroepitaxial growth of thin films featuring non-coherent interfaces in the C/Si system.
On the other hand, processes relying upon prevention of precipitation in order to produce strained heterostructures will likewise benefit.
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