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 \ac{IBS} of SiC.
-Further studies revealed the possibility to form buried layers of SiC by IBS at moderate substrate and anneal temperatures \cite{lindner95}.
+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 orientation of the Si substrate, corresponding to a \unit[50]{at.\%} C concentration at the implantation peak, C atoms get redistributed appropriately resulting in the formation of a stoichiometric buried layer of SiC exhibiting a well-defined interface to the Si host matrix.
-Redistribution is hindered for overstoichiometric doses ... and higher implantation energies resulting in randomly distributed SiC precipitates ...
+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.
+
+To summarize, by understanding some basic processes, \ac{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 \ac{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.
-high t -> direct SiC formation -> no redistribution ...
+\section{Substoichiometric concentrations of carbon in crystalline silicon}
+
+In the following some basic properties of C in crystalline Si are reviewed.
+A lot of work has been done contributing to the understanding of C in Si either as an isovalent impurity as well as at concentrations exceeding the solid solubility limit.
+A comprehensive survey on C-mediated effects in Si has been published by Skorupa and Yankov \cite{skorupa96}.
+
+\subsection{Carbon as an impurity in silicon}
+
+Below the solid solubility, C impurities mainly occupy substitutionally Si lattice sites in Si \cite{newman65}.
+Due to the much smaller covalent radius of C compared to Si every incorporated C atom leads to a decrease in the lattice constant corresponding to a lattice contraction of about one atomic volume \cite{baker68}.
+The induced strain is assumed to be responsible for the low solid solubility of C in Si, which was determined \cite{bean71} to be
+\begin{equation}
+c_{\text{s}}=\unit[4\times10^{24}]{cm^{-3}}
+\cdot\exp(\unit[-2.3]{eV/k_{\text{B}}T})
+\text{ .} \text{{\color{red}k recursive!}}
+\end{equation}
+
+The barrier of diffusion of substitutional C has been determined to be around \unit[3]{eV} \cite{newman61}.
+However, as suspected due to the substitutional position, the diffusion of C requires intrinsic point defects, i.e. Si self-interstitials and vacancies.
+Similar to phosphorous and boron, which exclusively use self-interstitials as a diffusion vehicle, the diffusion of C atoms is expected to obey the same mechanism.
+Indeed, enhanced C diffusion was observed in the presence of self-interstitial supersaturation \cite{kalejs84} indicating an appreciable diffusion component involving self-interstitials and only a negligible contribution by vacancies.
+Substitutional C and interstitial Si react into a C-Si complex forming a dumbbell structure oriented along a crystallographic \hkl<1 0 0> direction on a regular Si lattice site.
+This structure, the so called C-Si \hkl<1 0 0> dumbbell structure, was initially suspected by local vibrational mode absorption \cite{bean70} and finally verified by electron paramegnetic resonance \cite{watkins76} studies on irradiated Si substrates at low temperatures.
+Measuring the annealing rate of the defect as a function of temperature reveals barriers for migration ranging from \unit[0.73]{eV} \cite{song90} to \unit[0.87]{eV} \cite{tipping87}, which is highly mobile compared to substitutional C.
+% diffusion pathway?
+% expansion of the lattice (positive strain)
+
+%\subsection{Agglomeration phenomena}
+% c-si agglomerattion as an alternative to sic precipitation (due to strain)
+% -> maybe this fits better in prec model in next chapter
+
+\subsection{Suppression of transient enhanced diffusion of dopant species}
+
+The predominant diffusion mechanism of most dopants in Si based on native self-interstitials \cite{fahey89} has a large impact on the diffusion behavior of dopants that have been implanted in Si.
+The excess population of Si self-interstitials created by low-energy implantations of dopants for shallow junction formation in submicron technologies may enhance the diffusion of the respective dopant during annealing by more than one order of magnitude compared to normal diffusion.
+This kind of diffusion, labeled transient enhanced diffusion (TED), which is driven by the presence of non-equilibrium concentrations of point defects, was first discovered for implantations of boron in Si \cite{hofker74} and is well understood today \cite{michel87,cowern90,stolk95,stolk97}.
+The TED of B was found to be inhibited in the presence of a sufficient amount of incorporated C \cite{cowern96}.
+This is due to the reduction of the excess Si self-interstitials with substitutional C atoms forming the C-Si interstitial complex \cite{stolk97,zhu98}.
+Therefore, incorporation of C provides a promising method for suppressing TED enabling an improved shallow junction formation in future Si devices.
+
+% in general: high c diffusion in areas of high damage, low diffusion for substitutional or even sic prec
+
+\subsection{Strained silicon and silicon heterostructures}
+
+% lattice location of implanted carbon
+Radiation damage introduced during implantation and a high concentration of the implanted species, which results in the reduction of the topological constraint of the host lattice imposed on the implanted species, can affect the manner of impurity incorporation.
+The probability of finding C, which will be most stable at sites for which the number of neighbors equals the natural valence, i.e. substitutionally on a regular Si site of a perfect lattice, is, thus, reduced at substitutional lattice sites and likewise increased at interstitial sites.
+Indeed, x-ray rocking curves reveal a positive lattice strain, which is decreased but still remains with increasing annealing temeprature, indicating the location of the majority of implanted C atoms at interstitial sites \cite{isomae93}.
+Due to the absence of dislocations in the implanted region interstitial C is assumed to prevent clustering of implantation-induced Si self-interstitials by agglomeration of C-Si interstitials or the formation of SiC precipitates accompanied by a relaxation of the lattice strain.
+
+% link to strain engineering
+However, there is great interest to incorporate C onto substitutional lattice sites, which results in a contraction of the Si lattice due to the smaller covalent radius of C compared to Si \cite{baker68}, causing tensile strain, which is applied to the Si lattice.
+Thus, substitutional C enables strain engineering of Si and Si/Si$_{1-x}$Ge$_x$ heterostructures \cite{yagi02,chang05,kissinger94,osten97}, which is used to increase charge carrier mobilities in Si as well as to adjust its band structure \cite{soref91,kasper91}.
+% increase of C at substitutional sites
+Epitaxial layers with \unit[1.4]{at.\%} of substitutional C have been successfully synthesized in preamorphized Si$_{0.86}$Ge$_{0.14}$ layers, which were grown by CVD on Si substrates, using multiple-energy C implantation followed by solid-physe epitaxial regrowth at \unit[700]{$^{\circ}$C} \cite{strane93}.
+The tensile strain induced by the C atoms is found to compensates the compressive strain present due to the Ge atoms.
+Studies on the thermal stability of Si$_{1-y}$C$_y$/Si heterostructures formed in the same way and equal C concentrations showed a loss of substitutional C accompanied by strain relaxation for temperatures ranging from \unit[810-925]{$^{\circ}$C} and the formation of spherical 3C-SiC precipitates with diameters of \unit[2-4]{nm}, which are incoherent but aligned to the Si host \cite{strane94}.
+During the initial stages of precipitation C-rich clusters are assumed, which maintain coherency with the Si matrix and the associated biaxial strain.
+Using this technique a metastable solubility limit was achieved, which corresponds to a C concentration exceeding the solid solubility limit at the Si melting point by nearly three orders of magnitude and, furthermore, a reduction of the defect denisty near the metastable solubility limit is assumed if the regrowth temperature is increased by rapid thermal annealing \cite{strane96}.
+Since high temperatures used in the solid-phase epitaxial regrowth method promotes SiC precipitation, other groups realized substitutional C incorporation for strained Si$_{1-y}$C$_y$/Si heterostructures \cite{iyer92,fischer95,powell93,osten96,osten99,laveant2002} or partially to fully strain-compensated (even inversely distorted \cite{osten94_2}) Si$_{1-x-y}$Ge$_x$C${_y}$ layers on Si \cite{eberl92,powell93_2,osten94,dietrich94} by \ac{MBE}.
+Investigations reveal a strong dependence of the growth temperature on the amount of substitutionally incorporated C, which is increased for decreasing temperature accompanied by deterioration of the crystal quality \cite{osten96,osten99}.
+While not being compatible to very-large-scale integration technology, C concentrations of \unit[2]{\%} and more have been realized \cite{laveant2002}.
+
+\section{Assumed silicon carbide conversion mechanisms}
+\label{section:assumed_prec}
-.. lindner limit in dose -> 1250
-... two temp implantation ... sharp interface
-By understanding some basic processes (32-36), \ac{IBS} nowadays has become a promising method to form thin SiC layers of high quality exclusively of the 3C polytype embedded in and epitactically aligned to the Si host featuring a sharp interface \cite{lindner99,lindner01,lindner02}.
+Although high-quality films of single-crystalline 3C-SiC can be produced by means of \ac{IBS} the precipitation mechanism in bulk Si is not yet fully understood.
+Indeed, closely investigating the large amount of literature reveals controversial ideas of SiC formation, which are reviewed in more detail in the following.
-\section{Substoichiometric concentrations of carbon in crystalline silicon}
+\begin{figure}[ht]
+\begin{center}
+\subfigure[]{\label{fig:sic:hrem:c-si}\includegraphics[width=0.48\columnwidth]{tem_c-si-db.eps}}
+\subfigure[]{\label{fig:sic:hrem:sic}\includegraphics[width=0.48\columnwidth]{tem_3c-sic.eps}}
+\end{center}
+\caption{High resolution transmission electron microscopy (HREM) micrographs\cite{lindner99_2} of agglomerates of C-Si dimers showing dark contrasts and otherwise undisturbed Si lattice fringes (a) and equally sized Moir\'e patterns indicating 3C-SiC precipitates (b).}
+\label{fig:sic:hrem}
+\end{figure}
-diffusion mechanism, lattice distortion, hmm ... extra section needed?
+\ac{HREM} investigations of C-implanted Si at room temperature followed by \ac{RTA} show the formation of C-Si dumbbell agglomerates, which are stable up to annealing temperatures of about \unit[700-800]{$\circ$C}, and a transformation into 3C-SiC precipitates at higher temperatures \cite{werner96,werner97}.
+The precipitates with diamateres between \unit[2]{nm} and \unit[5]{nm} are incorporated in the Si matrix without any remarkable strain fields, which is explained by the nearly equal atomic density of C-Si agglomerates and the SiC unit cell.
+Implantations at \unit[500]{$\circ$C} likewise suggest an initial formation of C-Si dumbbells on regular Si lattice sites, which agglomerate into large clusters \cite{lindner99_2}.
+The agglomerates of such dimers, which do not generate lattice strain but lead to a local increase of the lattice potential \cite{werner96}, are indicated by dark contrasts and otherwise undisturbed Si lattice fringes in \ac{HREM}, as can be seen in Fig.~\ref{fig:sic:hrem:c-si}.
+A topotactic transformation into a 3C-SiC precipitate occurs once a critical radius of \unit[2]{nm} to \unit[4]{nm} is reached.
+The precipitation is manifested by the disappearance of the dark contrasts in favor of Moir\'e patterns (Fig.~\ref{fig:sic:hrem:sic}) due to the lattice mismatch of \unit[20]{\%} of the 3C-SiC precipitate and the Si host.
+The insignificantly lower Si density of SiC of approximately \unit[3]{\%} compared to c-Si results in the emission of only a few excess Si atoms.
+
+
+The formation of SiC by a preceeding agglomeration of C-Si dumbbells is supported by studies ... \cite{koegler03,eichhorn99}
+
+
+
+In contrast, investigations of strained Si$_{1-y}$C$_y$/Si heterostructures formed by MBE\cite{strane94,guedj98}, which incidentally involve the formation of SiC nanocrystallites, suggest an initial coherent precipitation by agglomeration of substitutional instead of interstitial C.
+
+Coherency is lost once the increasing strain energy of the stretched SiC structure surpasses the interfacial energy of the incoherent 3C-SiC precipitate and the Si substrate.
+
+These two different mechanisms of precipitation might be attributed to the respective method of fabrication.
+
+While in CVD and MBE surface effects need to be taken into account, SiC formation during IBS takes place in the bulk of the Si crystal.
+
+However, in another IBS study Nejim et~al.\cite{nejim95} propose a topotactic transformation that is likewise based on the formation of substitutional C.
+
+The formation of substitutional C, however, is accompanied by Si self-interstitial atoms that previously occupied the lattice sites and a concurrent reduction of volume due to the lower lattice constant of SiC compared to Si.
+
+Both processes are believed to compensate one another.
+
+
+%cites:
+
+% continue with strane94 and werner96
+
+%ibs, c-si agglom: werner96,werner97,eichhorn99,lindner99_2,koegler03
+%hetero, coherent sic by sub c: strane94,guedj98
+%ibs, c sub: nejim95
+%ibs, indicated c sub: martin90 + conclusions reeson8x, eichhorn02
+%more: taylor93, kitabatake contraction along 110, koegler03
+%taylor93: sic prec only/more_easy if self interstitials are present
+
+% -> skorupa 3.2: c sub vs sic prec
+
+% remember!
+% werner96/7: rt implants followed by rta < 800: C-Si db aggloms | > 800: 3C-SiC
+% taylor93: si_i reduces interfacial energy (explains metastability) of sic/si
+% eichhorn02: high imp temp more efficient than postimp treatment
+
+% todo
+% add sharp iface image!
-\section{Assumed cubic silicon carbide conversion mechanisms}
-\label{section:assumed_prec}
on surface ... md contraction along 110 ... kitabatake ... and ref in lindner ... rheed from si to sic ...
also indictaed by other direct synthesis experiments like martin90 and conclusions of reeson8X ...
eichhornXX, koegler, lindner ...
+
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