Nucleation and growth of gallium arsenide on silicon

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Applied Surface Science 56-58 (19921 589-596 North-Holland

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Nucleation and growth of gallium arsenide on silicon B. Bartenlian,

R. Bisaro,

J, O!ivier, J,-P. Hirtz

Thomson-(Sl.~ LCR, I)tmlahte de Corberilh' 91404 ()~ay Cedex, l:ran¢'e M, Pitaval, J. Meddeb Dt;partentent tit' Physique de.~ Mat~;riatLr, Unit'crsitt t Chtttth' Bernard L~'ott 1, 4.t honlcrard dn I I Not'etnbre ]l) ~, 69622 Vilh'ttrbatme. l"rance

and

A. Rochcr Cetttre d'Elaboration ~h's Matt;rkut.v et d'[-tudt's Strt~t'ittrah's, Lahnr.toirc d'Optique Eh'ctroniq~:e Jt,atltl(, l~[arl'i[,,, 31400 Touh)nse, [:ralu't'

(CEMES-I ~ )E) du CNRS, 29 rue

Received 6 May 1991: accepted fi~r publication 18 June 1991

We have studied the early stages of gr(:wth of GaAs on Si ((){)1), mis~)liented h~, ~to towards [11~)]. t~, migration-enhanced epituxy (MEE) in a molecular beam epila~,'y (MBE) system. Wc present results using in situ and ex situ analyses, reflection high-energy electron diffraction (RHEED), transmission electron microscopy (TEM), X-ray photoelectron spectroscopy (XPS), X-ray photoelectron diffraction (XPD) and X-ray double crystal diffractometry (DCD) performed at each stage of growth. We show that the nucleation by MEE induces a surhJce roughness decreasing as the layer becomes thicker. XPD experiments at the onset of the growth show a stoichiometric GaAs without antiphase domains. The relaxation of stress for the layers deposited at low temperature (3lit)o C) occurs viu partial dislocation migration developing between them. stacking faults and mierotwins. The post annealing of films with thicknesses less than 61) nm drives the formation of three-dimensional islands on the silicon surface. These islands develop (I 14) facets in the [1111] direction where the Ga migration is greater. The height over the base ratio of these islands is uniform and from Bauer's relation, we calculate the interface energy which can be correlated to the strain energy due to the disloeution~ near the interface. There are two reasons lor the restructuration of the surface. First, from a thermodynamical approach including surhlces and interface energies of the G a A s / S i system, we can demonstrate that the growth of GaAs/Si is a three-dimensional Volmer-Weber growth. Second, the bulk energy due to the compressive strain of a continuoux pseudomorphie GuAs film is higher than that of an islanding GaAs where relaxation from the free surfaces of tile islands occurs. The post annealing of films with thicknesses higher than 60 nm has a smoothing effect and can give a perfect two-dimensional tt),~)l) surface. Growth at higher lemperatures (580 ° C) suppresses plane defects, creates a Lomur dislocation network at the interface and induces the 611° dislocations in the bulk. We perform X-ray DCD experiments on 90 nm thick layers. We observe, before annealing a compressive biaxial stress of about - 1.9 × I11~ Pa which is the driving force for the elimination of the plane defects.

1. Introduction G r o w t h o f I i I - V s e m i c o n d u c t o r s o n Si s u b stratus has received considerable attention recently because of the potential for monolithic i n t e g r a t i o n o f I i l - V o p t i c a l w i t h a d v a n c e d Si 0169-4332/92/$05.1](1

electronic circuitry [1-3] on the same wafer. Furt h e r i m p r o v e m e n t s in t h e G a A s q u a l i t y a r e n e c e s sary to realize high-quality minority-carrier devices such as laser diodes. Research progress m a d e d u r i n g t h e p a s t t h r e e y e a r s is i m p o s i n g a s i g n i f i c a n t c h a n g e in o u r o u t l o o k f o r h e t e r o e p i -

a~ 1992 - Elsevier Science Publishers B.V. All rights reserved

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B. Bart('nlian

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aL / Nucleation and growth of GaAs on SHOO1)

taxy of mismatched systems. There has been considerable progress recently in understanding the stability of vicinal surfaces of silicon (0111) [4-6]. The configuration of steps, their roughness, spacing, have been analysed by scanning tunneling microscopy [4,7], and step and kink energies directly determined [4] for ideal vicinal (001) Si surfaces. The configuration of the steps affects a number of surface p h e n o m e n a including the nature of the growth of heteroepitaxial oveflayers. The surface energy of the vicinal Si (0111) surfaces is dependent on the nature of the first atomic layer adsorbed on it [8]. Very few authors [9] have focussed their attention on the influence of the substrate temperature T,, at which the atomic prelayer is adsorbed, on the configuration of the steps. The aim of this work is firstly to study As adsorption on vicinal Si (001) and its influence on the G a A s overlayer orientation and sect,udly to study the initial stages of nucleation and growth of G a A s in near thermodynamical equilibrium conditions.

2. Experimental To avoid the formation of antiphasc domains, Si (001) substrates misorientcd by 4 ° towards [110] have been used for the MBE growth of GaAs layers. Silicon substratcs have been cleaned using a modified Shiraki method [10]. The protective thin oxide layer on silicon substrates have been desorbed at 1(101)°C for 15 min in the MBE chamber. For the two first series of samples O. and O~, the surfaces have then been exposed to As 4 flux during 5 min at T, = 400 and 7011° C, respectively. Samples of series 1, 3, 4 and 2 have a prelayer of O,, and O~ type, respectively. T h e n we have grown on them G a A s layers at T, = 3 0 0 ° C by MEE. M E E ronsists of N times the following sequence: 2 s of Ga flux exposure, 9 s Ga and As fluxes off, 2 s of As flux exposure, 2 s Ga and As fluxes off. Each cycle appr-ximately corresponds to one G a A s monolayer deposition. N is equal to 14 and 40 for samples of series 1 and 2, 3, 4, respectively. Conventional MBE growth procedure is then pursued for samples of series 3 and 4, at T, = 300 o C and a growth rate

of (1.5 # m / h with Ga and As fluxes on simultaneously. Total thicknesses of samples of series 1, 2, 3, 4 are 3, 8, 15, 90 nm, respectively. MBE growth bcing a growth method far from thermodynamical equilibrium conditions, wc have further annealed samples of series 1, 2, 4 at T, = 630 o C for 20 min under an As 4 exposure, those samples are labeled lr, 2r, 4r respectively. We have then assumed that the state of samples It, 2r is identical to the state resulting from a growth at thermodynamical equilibrium conditions. XPS and XPD measurements are performed ex situ. Samples are transferred from the Riber 23(10 MBE system to the VG E S C A L A B XPS system under nitrogen gas in a special attachment.

3. Results 3.1. A s adsorption on vicinal Si( lO0) stoface

Vicinal surfaces can exhibit different structural phases, since steps of different types may be favored depending on temperature T,, angle of misorientation, and nature of atoms adsorbed on it. The Si (01}1) surface reconstructs to form rows of dimerized atoms. Because of the symmetry of the diamond lattice, dimer rows are parallel to each other on terraces. Just after de-oxidation of the surface at T, = 1000 ° C and before As 4 exposure. the R H E E D pattern is (2 × 2), which m e a n s that both (2 x 1) and (1 x 2) domains are present. However, the stronger intensity of the half-order diffraction line in the [110] azimuth perpendicular to the step ledges, indicates a majority of (2 x I) domains. R H E E D patterns remain (2 × 2) after As 4 expo.~ure either at T~ = 400 or 700 ° C. if G a A s layers are grown on O,, and O r, under As-stabilized growth conditions, R H E E D patterns show (4 × 2) and (2 × 4) G a A s surface reconstructions respectively. In the former case As dimer rows are parallel to the step ledges, in the latter they are perpendicular. No XPS peaks corresponding to S i - O and A s - O bonds could be detected on samples of series O~; in contrast, such peaks are detected on oxide-free silicon and samples of series O,,. Naturally oxygen is coming from sample transfer to the XPS system.

591

K Bartenlian el al. / Mwleation and growth of GaAs on Si(O01) 3.2. Study o f thbz layers o f G a A s on Si

Samples of series 1, 2, 3, lr, 2r are used in this study. TEM cross-section micrographs show that growth of a 3 nm thick GaAs layer (sample of series 1) is just sufficient to cover continuously the silicon surface. The thicker the layers, the smoother the morphology of the surface. For the sample of series 1, the surface is rough Jnd gives a spotty RF~EED pattern with spots elongated in the [001] direction. XPD experiments performed in both [170] and [110] azimuth~ show no evidence of antiphase domains but a better crystalline quality after annealing. Stacking faults and microtwins are present in the as-grown layers (fig. 1). Partial dislocations of the ~(112) type are always aligned in the [l'i0] ~irection. Figs. 2a and 2b show TEM plan view micrographs of a sample of series 2. It must be noticed that the [170] direction for GaAs is perpendicular and parallel to the step ledges for O~, and O c procedure, respectively. After annealing at T, = 630°C for

20 min (samples of series Ir and 2r), three-dimensional islands of GaAs are formed on the silicon surface as shown in TEM cross-section micrographs (figs. 3a and 3b). The height-over-base ratio h / l of the pyramid-shaped islands is uniform and equal lo 0,2, Facets are present in the [110] azimuth. These facets make a q, angle of about 19.5 ° with respect to the (001) Si planes (fig. 3a), these are {114} facets. The R H E E D pattern consists of chevrons. Neither st~,:king faults nor microtwins are present in the [110] azimuth, but extra {111} planes in silicon indicates the presence at the interface of 60 °-dislocation lines parallel to the step ledges (fig. 3a). In contrast, the GaAs islands have a hemispherical shape in the [li0] azimuth. Very few stacking faults and microtwins are detected in the TEM cross-section micrograph (fig. 3b). The R H E E D pattern shows small round spots. The TEM planview micrograph (fig. 4) of a sample of series 2r shows a Lomer dislocation network at the interface of individual islands.

S~ckin 9 fault

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Z11Ol(~

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Fig. I. TEM cress-sectionr.-,..:.,~raphsof GaAs on Si grownat T, = 300 ° C (sampleof series 3).

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B. Bartcnlian el al. / Nuch'atioa and growth o]"GaA~"on SiffXH )

reduction of defects by a factor of 3. Coherence length is measured to be 13 and 21 nm for samples of series 4 and 4r, respectively. Subsequent annealing is beneficial to the crystallinity of the material and has a smoothing effect on the surface morphology. A desorientation of about 0.1 ° is measured between (001) planes of GaAs and Si. Samples are strained, X-ray DCD measurements give e~ = 5.4 × 10 -~ (lattice in compression) and ~ = - 5 . 4 × 10 -~ (lattice in tension), i.e., o-0 = - 7 . 3 × 10 7 Pa and o-, = 7.3 × 107 Pa for

3.3. Study of thick layer~ , GaAs on Si Samples of series 4 and 4r are used in this study. The surface morphology of the as-grown sample is smoother than that of as-grown samples of series 1, 2 and 3. Spots on the R H E E D pattern are more elongated in the [001] direction. After annealing (sample of series 4r), the surface is fiat and reconstructed (4 × 2). The full width at half maximum (FWHM) of X-ray DCD rocking curves decreases by a factor of 1.7 corresponding to a

[110]

/[1101

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,~

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Fig. 2. T E M p l a n - ~ e w micrographs o f a 8 nm thick G a A s / S i layer grown at T, = 3 0 0 ° C after one A s monolayer deposition at

T, = 700 oC. Sampleof series 2: (a) (220) reflection; (b) (220) reflection.

B. Bartenlian et al. /Nucleatton and growth o f GaAs on Si(O01)

4° Discussion and conclusion

samples of series 4 and 4r, respectively. During cooling of samples of series 4 and 4r to room temperature, the thermal expansion coefficient difference gives a tensile strain plastically relaxed until 350 ° C. A calculated value for compressive stress during growth at T, = 3 0 0 ° C is o-ii= - 1.9 x 10 a Pa. The corresponding G a A s lattice deformations are e ~ _ = l . 4 × 1 0 -'~ and eli = - l . 6 x 10 -'~ near the limit of the bond breaking.

XPS measurements performed on O~ type samples, indicate that at low temperatures, As atoms are either physisorbed or chemisorbed on the surface. The As chemisorption is not sufficiently completed and T, too low to drive the system to a new minimal energy configuration. G a A s growth on such substrates necessarily gives

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Fig. 3. TEM cross-section micrographs of a 3 nm thick ~ a A s / S i layer grown at ~ = 300°C and annealed at T~ = 630°C for 20 min. Sample of series I r: (a) [ 110] azimuth; (h) [ 1/0] azimuth.

594

B. Baltenlian et at / Ntwh'alion ,~oulgrowth o[ {haA,~ on SitO01)

a (4 ,'< 2) s u r f a c e as if the initial A s / S i s u b s t r a t c w a s c o m p l e t e l y (i × 2) single d o m a i n , in fact furt h e r o b s e r v a t i o n s by s c a a n i n g e l e c t r o n mic r o s c o p y have s h o w n a n a c c u m u l a t i o n o f a r s e n i c a t o m s on p a r a l l e l s t r a i g h t lines s e p a r a t e d by 50

nm. If t h e a c c u m u l a t i o n t a k e s p l a c e at s t e p ledges, s u p e r - s t e p s a r e also f o r m e d as a l r e a d y p r o p o s e d by P u k i t e et al. [t)] a n d K o c h [11]. In c o n t r a s t , a vicinal (I){ll) Si e x p o s e d ~9 A s 4 flux at T, = 7 0 0 ° C b e c o m e s -~trongly stable. X P S m e a s u r e m e n t s h a v e

Fig. 4. TEM plan-view micrograph of an 8 nm thi~k GaAs/Si layer grown at 7~.= .300°C after one As monolayer deposition at T~= 700 '~C and annealed at 630 ° C for 20 rain.

B. Bartenlian et aL / Nuch'ation and growth of GaAs on Si(001)

shown no other chemical bonds than S i - A s bonds. As atoms are chemisorbed on the vicinal C0011 Si surface. Thermal energy drives the system to a new stabilized (2 × l) A s / S i surface. Minimum energy configuration is given as in the case of Si(100) by thermodynamical considerations [12]. As dimers are parallel to tile step ledges and then subsequent G a A s surface layer is (2 x 4) reconstructed. The Ga diffusion coefficient is greater along the [110] direction perpendicular to the dimer rows than along [ll0] parallel to the dimer rows [13]. Hence for O,, procedure, the Ga diffusion is easier normal to the step ledges and ( l l 4 ) facets can be developed. In contrast, in the [110] direction, the growth is both controlled by Ga surface diffusion and mass transportation coming from the other azimuth. Islands are hemispherical in this direction. The energy stored in the bulk by a defect-free layer of thickness h is given by:

E = (Ctt + C,2 - ( 2 C ~ f f C , , ) ) E ~ h , where C ~ t = l l . 8 8 x 1 0 m Pa and C ~ 2 = 5 . 3 8 × 10 m Pa. For h = 10 rim, E = 1.91 × 10 -4 J / c m 2 is similar to a surface energy. At high temperature we can use the thermodynamical model of Bauer [14] which gives for the height-over-base ratio of islands: h / / I = ( O d e p + Ori -- O'sup) /2Ordcp,

where o'aep, tri, O~,up are the specific free surface energies of the deposit, interface and substrate, respectively. O'dcp(001)= 1.11 X 10 -4 J / c m -~ for the G a A s and o-~p(001)= 10 -4 J / c m 2 for the A s / S i surface [8]. We find tri = 3 × I0 - s J / c m 2, hence the fundamental relation for a three-dimensional growth is satisfied 4.1

X

10 -'~ = O'd~p + O"i -- oh, p >. O.

We notice that o-~ = 0.6 e V / d i m e r has the same order of magnitude as the bond-breaking energy. The relaxation mechanism in these layers proceeds firstly by the nucleation at the surface of a V-shaped dislocation (energetically more favorable than half a loop) caused by the strain due to the lattice mismatch between GaAs and Si. In GaAs, perfect 60 °-dislocations of the ½( 1101 type

595

(glide or shuffle) are called either a (As dislocation core) o r / 3 (Ga dislocation core), a dislocations have an As understoichiometric core, more kinks and hence have higher velocities: V ( a ) > V(screw)> V(/3). These perfect dislocations are dissociated a(/3)60 ° --, a(/3)9{) ° + a(/3)30 o. At low temperature, the slowest /330 ° (t,t~,~o> v0s,) is quenched at the surface, the fastest ~390 °, after their nucleation at the surface, migrate to the interface in the glide planes {1111, and generate stacking faults and microtwins between them. The stacking faults and microtwins, regarded as grown-in faults are terminated at both the surface and the interlace with partial dislocations of the ~(1121 type. They are hence visible by T E M plan-view mainly in the [170] azimuth. At high temperature, under the compressive stress due both to the lattice mismatch and to the difference in thermal expansion coefficient, 1330 ° move to the interface, cancel the stacking faults and microtwins and react with the /390° at the interface: /330 ° + B90 o ~ / 3 6 0 °, Then two f160 ° with antiparallel Burgers vector combine at the interface to form the motionless pure-edge Lomer dislocation network clearly visible in the T E M plan-view micrograph (fig. 4). The energy of such a dislocation array is given by E = ( I n ( R / r ) p . b 2 ) / ( 4 ~ r ( 1 - v ) ) w h e r e v is Poisson's ratio = 0.3, b the modulus of the Burgers vector r = b = 0.4 nm,/~ = 3.25 Pa, R = 5 nm. We find E = 3 x 10 -5 J / c m 2 because there are 20 km of dislocation lines at the ;~aerface per cm 2. The energy due to (;';,locations being comparable to o-i, the interface er,ergy is coming mainly from Lomer dislocation network. Both bulk and surface effects drive the G a A s / S i system to a threedimensional growth, probably of the V o l m e r Weber type because of the strong stability of an As monolayer on silicon at high temperature.

Acknowledgements We would like to thank A. Wochenmayer and F. Wyczisk for their help in the experimental

596

B. Bartenlian et hi. / Nnclearion imd growth o f GaAs on Si(O01)

w o r k . T h i s w o r k h a s b e e n p a r t l y s u p p o r t e d by the F r e n c h D i r e c t i o n d e s R e c h e r e h e s et E t u d e s Techniques of the Direction G6n~rale de ['Armement.

References [I] J.C.C. Fan and J.M, Poatc, Eds., Heteroepitaxy on Si, Vol. 67 (Materials Research Society, Pittsburgh, PA, 1986). [21 J.C.C. Fan, J.M. Phillips and B.Y, Tsaur, Eds., Heteroepitaxy on Si I!, Vol. 91 (Materials Research Society, Pittsburgh, PA, 1987), [31 H.K. Choi, R. Hull, H. Ishiwara and RJ. Nemanich, Eds,, Heteroepitax'v on Si: Fundamentals, Structures and Devices, Vol. 116 IMaterials Research Society, Pittsburgh, PA, 1988).

[4] B.S. Swarzentruber, Y.W. Mo, R. Kariotis, M.G. Lagally and M,B. Webb, Phys. Rev. Lett. 65 (1990) 1913. [5] T.W. Poon, S, Yip, P.S. Ho and F.F, Abraham, Phys. Rev. Lelt. 65 (19901 2161. [~,] O.L. Alerhand, A. Nihat Berker, J.D. Joannopoulos, D. Vanderbilt, R.J. Hamers and J.E. Demuth, Phys. Rev. Lett. 64 (1990) 2406. [7] MG. Lagally, R. Kariotis, B.S. Swartzentruber and Y.W. Mo, Ultramieroscopy 31 (1989) 87, [8] J.E. Northrnp, Phys. Rev. Len. 62 (1989) 2487. [q] P.R, Pukite and P.I. Cohen, Appl. Phys. Len. 50 0987) 1739, [10] A. Ishizaka and Y. Shiraki, J. Electrochem. Soc, 133 (1986) 666. Ill] S.M. Koch. PhD Thesis, Stanford University (1986). [12] R.S. Becket, T. Klitsner and J.S. Viekers, J. Microse. 152 (1988) 157. [13] K. Ohta, T. Kojima and T. Nakagawa, J. Cryst. Growth 95 (1989) 71. [14] E. Bauer, Z. Krist. ll0 (19581 372.