\documentstyle{article}% If there are no figures in a paper %\documentstyle[nanosymp,epsf]{article}% If there are PostScript figures \begin{document}
\textwidth=160mm
\textheight=237mm
\topmargin=-2cm
\leftmargin=1cm
%\rightmargin=-2cm
\oddsidemargin 0 cm
\renewcommand{\arraystretch}{1.5}
\renewcommand{\baselinestretch}{2.0}
\begin{document}

\title{\LARGE \bf HETEROEPITAXIAL GROWTH OF InAs ON Si: THE NEW TYPE OF QUANTUM DOTS}
\author{\empty}
\date{\empty}
\maketitle
\centerline{\large \underline{G.E.CIRLIN$^{1,2}$}, N.K.POLYAKOV$^{1,2}$, 
V.N.PETROV$^1$, V.A.EGOROV$^1$,}
\centerline{\large D.V.DENISOV$^2$, B.V.VOLOVIK$^2$, 
V.M.USTINOV$^2$, Zh.I.ALFEROV$^2$, 
N.N.LEDENTSOV$^{2,3}$,}
\centerline{\large R.HEITZ$^3$, D.BIMBERG$^3$, N.D.ZAKHAROV$^4$, 
P.WERNER$^4$, U.G\"OSELE$^4$
}
\bigskip
\centerline{\sl $^1$ Institute for Analytical Instrumentation RAS, St.Petersburg, Russia}
\centerline{Tel: 007--812--2479350, Fax: 007--812--2517038, e-mail: cirlin@beam.ioffe.rssi.ru } 
\centerline{\sl $^2$ A.F.Ioffe Physico-Technical Institute RAS, St.Petersburg, Russia}
\centerline{\sl $^3$ Technical University, Berlin, Germany }
\centerline{\sl $^4$ Max-Plank-Institute for Microstructure Physics, 
Halle/Saale, Germany}


\thispagestyle{empty}
\eject

\centerline{\bf Abstract}
Under certain growth conditions InAs/Si heteroepitaxial growth proceeds via 
Stranski--Krastanow or Volmer--Weber growth modes depending on the growth 
parameters. Critical thickness at which three dimensional InAs islands start 
to appear at the Si(100) surface is within the range of 0.7~ ---~ 4.0 monolayers 
(substrate temperature range is $350^{\circ}C~ - 430^{\circ}C$). Their size depends critically 
on the growth conditions and is between 5 nm and 80 nm (uncapped islands). 
Critical lateral size of the coherent (Si capped) dislocation--free island is 
equal to 2~ ---~ 5 nm depending on the island height. Islands having larger size 
are dislocated. Optical properties of InAs nanoscale islands capped with Si 
reveal a luminescence band in the $1.3 \mu m$ region. 


\section*{Introduction} 

Self--organisation effects at semiconductor surfaces during molecular beam 
epitaxy (MBE) attract strong interest during the last decade. In this 
direction the most attention was paid to the quantum dots (QDs) in III/V~ 
---~ III/V 
({\it e.g.}, InAs~ ---~ GaAs)\cite{r1,r2}, II/VI~ ---~ II/VI ({\it e.g.}, 
CdSe~ ---~ ZnSe)\cite{r3} or IV~ ---~ IV 
(Si~ ---~ Ge)\cite{r4} materials systems. The last system attracted much efforts 
because silicon remains a key material in modern microelectronics industry. 
For optoelectronic applications silicon is not well suitable because of its 
indirect band gap nature. Attempts to improve the situation using SiGe~ ---~ Si QDs 
did not led to significant progress as these nanostructures provide indirect 
band alignment both in $k$~-- and in real space. In order to increase luminescence 
efficiency of silicon--based structures we propose to insert direct bandgap 
InAs QDs in a Si matrix using MBE growth\cite{r5}. We emphasize that this approach 
is different from the growth utilizing InAs QDs on thick GaAs buffers\cite{r6}. 
Small coherent InAs islands additionally offers a possibility to overcome the 
problems of strong lattice mismatch and formation of antiphase domains inherent 
for the growth of thick III~ ---~ V epilayers ({\it e.g.} GaAs) on Si. 

In this paper we report on the MBE growth and properties of the InAs nanoscale 
islands formed on silicon. We have found that under certain growth conditions 
InAs/Si heteroepitaxial growth proceeds via Stranski--Krastanow or Volmer--Weber 
growth modes depending on the growth parameters leading to the formation of 
nanoscale islands. The lateral size depends critically on the growth conditions 
and is in the range of 2~ ---~ 80 nm. Optical properties of InAs QDs capped with 
Si reveal a luminescence band in the $1.3 \mu m$ region. A pronounced excitation 
density dependence of the photoluminescence (PL) peak position and a decay 
time of order of 400 ns were observed. 

\section*{Experimental details}

      The growth experiments were carried out using EP1203 (Russia) or Riber 
Supra (France) MBE machines on exactly oriented Si(100) substrates. The 
Si(100) surface preparation was made in a way similar to that described in \cite{r7}. 
Thermal desorption of silicon native oxide was performed in a growth chamber 
at substrate temperature of $830^{\circ}C - 870^{\circ}C$ during 15 min. After that, well 
resolved ($2\times 1$) or mixed ($1\times 2$) and ($2\times 1$) surface reconstructions typical for 
cleaved Si(100) surface were observed. Then the substrate temperature was 
smoothly decreased to the desired value and the InAs deposition was initiated 
in a conventional MBE mode. The InAs deposition rate was typically 0.1 
monolayer (ML) per second. After the deposition of the desired average 
thickness of InAs has been completed, the sample was immediately quenched to 
the room temperature and removed from the growth chamber in order to provide 
scanning tunnelling microscopy (STM) and scanning electron microscopy (SEM) 
studies. SEM measurements were performed using CamScan setup. Pieces for STM 
studies were covered with silicon vacuum oil immediately after exposure to 
atmosphere, a procedure used before for our studies of InAs QDs on GaAs 
surfaces \cite{r8}. A calibration of the growth rate, III--V flux ratio, and 
monitoring of the surface morphology during growth was performed using 
reflection high energy electron diffraction (RHEED) system composed of a high 
sensitivity video camera, a video tape recorder and a computer, all 
interconnected via a specially--designed interface \cite{r9}. Transmission electron 
microscopy (TEM) images were obtained on the samples when InAs islands were 
covered with 30~ --- 50 nm of silicon using JEOL 4000EX microscope. The same 
structures were used for luminescence experiments. A 514.5 nm line of the $Ar^{+}$ 
laser was used as an excitation source and a Ge photodiode as a detector. The 
excitation density was $5 - 100 {W\over cm^2}$. The structures for TEM and PL studies 
were grown in a following manner. Just after InAs deposition, a 10~ ---~ 20 nm Si 
cap layer was grown at the same substrate temperature $T_s$ as for InAs 
deposition followed by a 10 min annealing procedure at $650^{\circ}C - 700^{\circ}C$. Then 
20~ ---~ 40 nm of Si was grown at the same temperature with optional consequent 
10 min annealing at $700^{\circ}C - 800^{\circ}C$ in order to smooth resulting surface.

\section*{Results and discussion}

       In Fig.1 we summarise RHEED data of the dependencies of the substrate 
temperatures and $\rm As_4/In$ fluxes ratio in the ranges $350^{\circ}C - 450^{\circ}C$ and 2~ ---~ 12, 
respectively, on the critical thickness $d_{crit}$ (mean thickness at which 3D 
nanoscale islands start to appear at the surface). The lowest limit of  $T_s$ 
was chosen for the two following reasons: first, determination of $T_s$ at very 
low temperatures is difficult and second, because of poor quality of Si cap 
layer grown at too low temperatures. Poor quality of the Si cap layer may 
hinder device applications. On the other hand we have found that at temperatures 
higher then $450^{\circ}C$ {\it no nanoislanding is observed on the surface} independently 
on the V/III flux ratio, total arsenic overpressure, {\it etc}. limiting the range 
for the highest $T_s$ as indicated in Fig.1. (At higher $T_s$ only mesoscopic 
dislocated clusters with lateral size of $\sim$ 400 nm appeared on the Si(100) 
surface \cite{r10}). Following the data presented in Fig.1 one can conclude that 
depending on the growth conditions, 2D~ ---~ 3D transition can be tuned in the 
thickness range of 0.7~ ---~ 4.0 ML. On the other hand, we have observed two 
different growth modes for the InAs/Si heteroepitaxial growth. First, 
Volmer--Weber growth mechanism could be realised when InAs nanoscale islands 
are formed on the bare Si substrate ({\it i.e.} for $d_{crit}<1$ ML). Typical RHEED 
pattern for this case is presented in Fig. 2,a. For higher $T_s$ ($400^{\circ}C - 450^{\circ}C$) 
we observed Stranski--Krastanow growth mechanism when $d_{crit}$ exceeded 1.0 ML 
and the formation of InAs nanoscale islands occurs on the top of the wetting 
layer. RHEED pattern for this particular case is presented in Fig. 2,b. These 
results are in agreement with the 2D~ ---~ 3D transition intervals observed by 
other groups \cite{r11,r12}.


     Typical surface morphology extracted from the STM images for these cases 
revealed high density arrays of nanoscale islands (with surface density of 
the order of $(1\div 6)\times 10^{11} cm^{-2}$) and the lateral sizes of 10~ ---~ 20 nm \cite{r13,r14}. 
These islands are basically pyramidal in shape with rectangular or triangular 
base. 

     We have found that the growth conditions (the amount of InAs deposited, 
fluxes ratio {\it etc}.) influenced strongly on the arrangement, sizes and surface 
density of InAs/Si islands \cite{r15}. In order to clarify how growth parameters 
influence on the surface morphology we grew three samples, labelled 1--3. 
Their MBE growth conditions and geometrical characteristics measured with SEM 
are presented in the Table 1.

    The main trends for the uncapped InAs/Si island formation are the 
following. For higher As pressure and the same InAs mean thickness the island 
size becomes smaller as compared to the case of the decreased As/In fluxes 
ratio (samples 1 and 3). We believe that this is due to the suppressing of 
the surface migration length. The density of the islands for the sample 3 is 
four times higher than that of the sample 1. The islands size distribution 
became narrower, too. With increasing the average InAs thickness (sample 2) 
the islands start to form large conglomerates. These clusters are elongated 
and exhibit crystallographic shape with the base oriented along $[011]$ and 
$[0\bar 11]$ directions. Similar situation was observed in \cite{r11} where such clusters 
were formed during the InAs growth on Si(100) passivated (hydrogen--terminated) 
surfaces with approximately same size and surface density. We note that for 
the sample 2 the critical thickness is larger by the value of 1.8 ML, {\it i.e.} 2.5 
times higher. In InAs/GaAs heteroepitaxial system such exceeding and large 
sizes should lead to the formation of the dislocations. In order to check the 
crystalline quality we studied the same uncapped samples with cross--sectional 
TEM. 

      Most of the islands exhibit considerably good crystalline quality 
although misfit dislocations appeared at the InAs/Si interface. In Fig. 3 
typical cross--section TEM image for single InAs nanoisland is presented. From 
high resolution TEM data taken from the Si capped samples we have found that 
the critical lateral size of the coherent dislocation--free islands is equal 
to 2-5 nm depending on the island height. Islands having larger size are 
defected due to misfit dislocations or twins. We also suggested that most of 
the InAs reevaporated during the deposition (or annealing) of the cap layer.
Fig.3. Typical high resolution cross-section TEM image of the single InAs 
nanoscale island at Si(100) surface. 

     The samples with such InAs quantum dots capped with 30~ ---~ 50 nm Si layer 
show a luminescence band in the $1.3 \mu m$ region for temperatures up to 
$\sim 170^{\circ}K$. 
The PL exhibited a pronounced blue shift with increasing the excitation density 
and decays with a time constant of $\sim$ 400 ns \cite{r16}. 
Fig.4  compares  low  temperature 
PL spectra of the 4.8 ML and the 1.6 ML InAs samples deposited in different 
growth modes (Stranski--Krastanow for 4.8 ML and Volmer--Weber for 1.6 ML \cite{r17}). 
These spectra reveal TO-- phonon assisted exciton emission of the Si substrate 
and barrier at $\sim$ 1.11 eV. The broad emission band in the 0.9 to 1.05 eV 
region was observed for both InAs deposition thickness suggesting its origin 
to be associated with the InAs nanoobjects embedded in Si matrix. This broad 
line was not observed on the samples with InAs insertions smaller than the 
critical thickness or on a Si substrate. However, the size of these insertions 
evaluated from the TEM measurements do not exceed 5 nm \cite{r16} what is definitely 
not enough for charge localisation and quantum confinement.  

     Our very recent detailed TEM investigation of these samples shows that 
besides a random distribution of InAs molecules, two kinds of ordering were 
detected. Such regions ($\phi < 6 nm$) are characterised by larger InAs concentration. 
The first kind shows an ordering of InAs molecules in $(101)$,$(10\bar 1)$ planes 
inclined to the $[001]$ growth direction, while in the second one the ordered 
$(110)$, $(1\bar 10)$ planes are parallel to the $[001]$ growth direction. The formation 
of such a structure results in energy lowering due to the decrease of the 
number of mixed Si--As and Si--In bonds. We believe that the appearance of the 
broad PL peak is the result of solid solution formation and/or ordering, 
because the incorporation of InAs molecules into Si matrix changes the band 
gap of material and possibly shifts the relative positions of conduction and 
valence bands. Detailed results on TEM study will be presented in a separate 
paper \cite{r18}.

\section*{Conclusions}

     We have demonstrated that InAs/Si(100) heteroepitaxial growth could 
proceed via Stranski--Krastanow or Volmer--Weber growth modes depending on 
the MBE growth parameters. Critical thickness at which three dimensional 
InAs islands start to appear at the Si(100) surface varies within the range 
of 0.7~ ---~ 4.0 monolayers (substrate temperature is $350^{\circ}C - 430^{\circ}C$, 
As/In fluxes 
ratio is 2~ --- 12). Their size depends critically on the growth conditions and 
is between 5 nm and 80 nm (uncapped islands). Critical lateral size for Si 
capped InAs coherent dislocation--free islands is equal to 2~ --- 5 nm depending 
on the island height. Islands having larger size are dislocated. Optical 
properties of the InAs nanoscale islands capped with Si reveal a luminescence 
band in the $1.3 \mu m$ region independently on the growth mode used for InAs 
deposition. The appearance of this peak can be the result of the solid 
solution formation and/or ordering due to the incorporation of InAs molecules 
into Si matrix during post growth annealing changes the band gap of material 
and possibly shifts the relative positions of conduction and valence bands. 
InAs/Si heteroepitaxial system seems to be promising for Si--based 
optoelectronic applications, in particular for short distance fiber optics 
and integration of microelectronics and optical parts at the same silicon 
wafer.

\section*{Acknowledgements}

       The authors thanks to N.P.Korneeva, V.N.Demidov, S.A.Masalov and 
V.M.Busov for assistance during experiments. This work was supported by 
INTAS Grant No. 96--0242, Russian Foundation for Basic Research Grant 
No 98--02--18317 and Grant No 99--02--16799, Russian Scientific Programmes 
\lq\lq Physics of Solid State nanostructures" Grant No 98--2029 and \lq\lq Methods 
and Devices of Micro-- and Nanoelectronics" Grant No 02.04.5.1.40.E.46.  

\eject

\begin{table}[h]
\caption{MBE growth conditions and geometrical characteristics of the InAs/Si(100)
 quantum dot samples}
\begin{center}
\begin{tabular}{|c|c|c|c|c|}
\hline
Sample&InAs thickness,&As/In fluxes&Island lateral&Surface density of\\
No.&ML&ratio&size, nm&islands, $\rm 10^{10} cm^{-2}$\\
\hline\hline
1&1.2&3&15---50&4\\
\hline
2&2.5&3&20---80&1.6\\
\hline
3&1.2&10&3--15&16\\
\hline
\end{tabular}
\end{center}
\end{table}
\eject


\begin{thebibliography}{17}

\bibitem{r1} L.Goldstein, F.Glas, J.Y.Marzin, M.N.Charasse, and G.Le Roux~//
Appl.Phys.Lett. {\bf 47}, p.1099 (1985).
\bibitem{r2} D.Leonard, M.Krishnamurthy, C.M.Reaves, S.P.Denbaars and 
P.M.Petroff~// Appl.Phys.Lett. {\bf 63}, p.3203 (1993).
\bibitem{r3} M.Strassburg, V.Kutzer, U.W.Pohl, A.Hoffmann, I.Broser, 
N.N.Ledentsov, D.Bimberg, A.Rosenauer, U.Fischer, D.Gerthsen, I.L.Krestnikov, 
M.V.Maximov, P.S.Kop'ev and Zh.I.Alferov~// Appl. Phys. Lett. 72, p. 942 (1998). 
\bibitem{r4} Feng Liu and M.G.Lagally~// Surf.Sci. 386, p.169 (1997).
\bibitem{r5} N.N.Ledentsov in {\em Proceedings of the 23rd International 
Conference on the Physics of Semiconductors, Berlin, Germany, July 21-26, 1996}, 
edited by M. Scheffler and R. Zimmermann, World Scientific, Singapoure, 1996, 
{\bf 1}, p. 19-24.
\bibitem{r6} J.M.Gerard, O.Cabrol and B.Sermage~// Appl.Phys.Lett. {\bf 68}, p.3123 (1996).
\bibitem{r7} A.Ishisaka and Y.Shiraki~// J. Electrochem. Soc. {\bf 133}, p.666 (1986).
\bibitem{r8} G.E.Cirlin, G.M.Guryanov, A.O.Golubok, S.Ya.Tipissev, 
N.N.Ledentsov, P.S.Kop'ev, M.Grundmann and D.Bimberg~// Appl. Phys. Lett. {\bf 67}, p.97 (1995). 
\bibitem{r9} G.M.Gur'yanov, V.N.Demidov, N.P.Korneeva, V.N.Petrov, 
Yu.B.Samsonenko and G.E.Tsyrlin~// Tech.Phys. {\bf 42}, p.956 (1997).
\bibitem{r10} G.E.Cirlin, V.N.Petrov, V.G.Dubrovskii, S.A.Masalov, A.O.Golubok, 
N.I.Komyak, N.N.Ledentsov, Zh.I.Alferov and D.Bimberg~// Tech.Phys.Lett. {\bf 24}, p.290 (1998).
\bibitem{r11} T.Mano, H.Fujioka, K.Ono, Y.Watanabe and M.Oshima~// Appl.Surf.Sci. {\bf 130--132}, p.760 (1998). 
\bibitem{r12} P.C.Sharma, K.W.Alt, D.Y.Yeh and K.L.Wang~// Appl. Phys. Lett. {\bf 75}, p.1273 (1999). 
\bibitem{r13} G.E.Cirlin, V.G.Dubrovskii, V.N.Petrov, N.K.Polyakov, 
N.P.Korneeva, V.N.Demidov, A.O.Golubok, S.A.Masalov, D.V.Kurochkin, 
O.M.Gorbenko, N.I.Komyak, V.M.Ustinov, A.Yu.Egorov, A.R.Kovsh, M.V.Maximov, 
A.F.Tsatsul'nikov, B.V.Volovik, A.E.Zhukov, P.S.Kop'ev, Zh.I.Alferov, 
N.N.Ledentsov, M.Grundmann and D.Bimberg~// Semicond.Sci.Technol. {\bf 13}, p.1262 (1998).
\bibitem{r14} G.E.Cirlin, V.N.Petrov, V.G.Dubrovskii, Yu.B.Samsonenko, 
N.K.Polyakov, A.O.Golubok, S.A.Masalov, N.I.Komyak, V.M.Ustinov, A.Yu.Egorov, 
A.R.Kovsh, M.V.Maximov, A.F.Tsatsul'nikov, B.V.Volovik, A.E.Zhukov, 
P.S.Kop'ev, N.N.Ledentsov, Zh.I.Alferov and D.Bimberg~// Semiconductors {\bf 33}, p.972 (1999).
\bibitem{r15} G.E.Cirlin, N.K.Polyakov, V.N.Petrov, V.A.Egorov, Yu.B.Samsonenko, 
B.V.Denisov, V.M.Busov, B.V.Volovik, V.M.Ustinov, Zh.I.Alferov, N.N.Ledentsov, 
D.Bimberg, N.D.Zakharov and P.Werner~// Czeck.J.Phys. {\bf 49}, p.1547 (1999).
\bibitem{r17} R.Heitz, N.N.Ledensov, D.Bimberg, A.Yu.Egorov, M.V.Maximov, 
V.M.Ustinov, A.E.Zhukov, Zh.I.Alferov, G.E.Cirlin, I.P.Soshnikov, N.D.Zakharov, 
P.Werner and U.G\"osele~// Appl. Phys. Lett. {\bf 74}, p.1701 (1999).
\bibitem{r17} InAs critical thickness for both samples was exceeded by the value of 0.8 ML.
\bibitem{r18} N.D.Zakharov et.al., In {\em MRS'99 Fall meeting proceedings}, to be published.



\end{thebibliography}

\eject
\section*{\centerline{Figure captions}}
\noindent
Fig.1. Critical thickness dependence on substrate temperature and $J_{As_4}/J_{In}$ 
fluxes ratio during InAs/Si growth.
\vspace{1in}

\noindent
Fig.2. Typical RHEED patterns in [011] direction taken after the formation of 
InAs 3D islands in Si(100) surface via (a) Volmer--Weber and (b) Stranski--Krastanow 
growth modes.
\vspace{1in}

\noindent
Fig.3. Typical high resolution cross-section TEM image of the single InAs nanoscale 
island at Si(100) surface.
\vspace{1in}

\noindent
Fig.4. PL data for the InAs nanostructures embedded in Si matrix: (a) InAs grown 
via Stranski--Krastanow, (b) InAs grown via Volmer--Weber growth modes.

\eject
\end{document}
