RHEED investigation of Ge surface segregation during gas source MBE of SiSi1 − xGex heterostructures

Surface Science 284 (1993) 305-314 Nor~“Holland

RHEED investigation of Ge surface segregation during gas source MBE of Si/Si,_, Ge, heterostructures N, Ohtani I, SM. Molder, M.H. Xie, J. Zhang and B.A. Joyce i~er~~c~~a~ Research Centre for Se~‘co~~tor Materials,The B~c~tt Laboratory,hnpetil College of .Yc&ce, Technology and Medicine, ~~i~ersi~ of London, La& SW7 23Z, UK Received 6 October 1992; accepted for publication 6 November 1992

Ge segregation during silicon gas source moi~lar beam epitaxy (Si-GSMB~) has been studied by in situ growth rate measurement using rejection-~~h*ener~-electron di~ction (FEDS intensity o~llations. Growth rate of Si on Si,_,Ge, ~aduaily decreases to the Si hom~pit~a~ growth rate, which is attributed to Ge surface segregation at the growth interface. This segregation has been modelled using a mass balance equation and it has been found that the observed growth rate enh~cement can be used as a direct measure of the Ge segregation. Using this novel in situ technique, concentration dependence of Ge segregation was studied, and it was found that the segregation decay curve is nonlinear, resulting in two segregation regimes dependent upon the Ge concentration, consistent with previous studies. Temperature dependence studies reveal that surface hydrogen, which is produced by the dissociation of S&H, and GeH, on the surface during growth, may act as a growth controlling surfactant, and comparison with solid source growth results suggests that it significantly suppresses the Ge segregation, leading to a more precise controf of the interface. Finally, the thermal stability of the segregated surfaces was examined. Growth interruption and annealing during Si overlayer growth on Sit _XGe, resulted in a small increase in the surface Ge concentration, which may be ascribed to the outdiffusion effect of Ge from the near surface region.

The real~ation of high quail heterost~ctnres, quantum wells, and superlattices compatible with existing Si based technology would open the possibility of a number of new device applications. Si,_,Ge,/Si layered structures are among the most promising candidates, and there has recently been a great deal of activity in this field [l]. IIowever, several problems still remain and must be addressed before ideal heterointerfaces can be achieved. Ge surface segregation is one of the most important issues in this respect. There have been many reports [2-71 on Ge segregation or “floating” at the growth front during Si i +Ge, molecut Permanent address: Semi~ndu~or Basic Technology Research Laboratory, Nippon Steel Corporation, 1638 Ida, Nakahara-ku, Kawasaki 211, Japan. ~39~g/93/$~.~

lar beam epitaxy MI3E;). This can greatly degrade the interface quality and thus is currently a matter of serious concern. Ge segregation has been attributed to a surface mediated effect [S], indicating that dynamical processes at the growth front play an important role in defining the abruptness of the heterostructures. Jesson et al. [9] have provided an atomistic picture of the Ge surface segregation on Si(OOl>(Z X 1) reconstructed surfaces. In their model, an interchange of Si and Ge atoms occurs at step edges of monolayer height islands. The interchange is chemically driven at the rebonded step edge configuration, which results from the (2 x 1) reconstruction. Thus, the (2 X 1) surface reconstruction and in particular its rebonded Sn-type step ~nfig~ation (using the notation of Chadi [lo]) acts as a chemically driven atom “pump” for Ge surface segregation. They have also identified the step-driven segregation as the origin of the

8 1993 - Elsevier Science Publishers B.V. Ah rights reserved

306

N. Ohtani et a!. / Ge segregation during S~-eS~3E

long-range ordered phase observed in Si,_,Ge, alloys El11 and succeeded in explaining the details of the obtained HRTEM images [12]. Recently, it has been shown that the Ge surface segregation can be suppressed by using a monolayer of As 131,Ga [5], or Sb ISI as a growth controlling “su~a~tant,” which can tune surface energetics chemically. These surfactants have a lower surface energy than Si and Ge, thus preventing the chemicaliy driven interchange of Si and Ge. In addition, chemical vapor deposition (CVD) results in an atmospheric H, ambient [13] and recent medium energy ion scattering (ME-IS) [ 141 and time-of-flight scattering and recoiling spectroscopy (TOF-SARS) [15] measurements reveal that surface hydrogen also acts as a surfactant. Thus, a surface terminated by hydrogen may lead to the suppression of the Si-Ge interchange at the growth front and result in a more precise control of the interface. In this paper we report the results of an investigation of Ge surface segregation monitored during Si gas source molecular beam epitaxy (SiGSMBE) by in situ growth rate measurements using the reflection-high-energy-electron diffraction (RHEED) intensity oscilIation technique. Its strong ability to probe the dynamical processes of Ge surface segregation has recently been demonstrated by the present authors [161 and Werner et al. 1171.Here, concentration and temperature dependent results of Ge surface segregation obtained from growth rate data are presented. Suppression of Ge segregation due to surface hydrogen, which is produced by the dissociation of disilane and germane on the surface during growth [I$], is also discussed.

2. Experimental procedure All experiments were carried out in a VG Semicon pu~ose-built double”ended MBE system which has been modified to acco~odate the hydrides of both Si and Ge (SiI-I,, Si,H,, GeH,). Base pressures below 5 X lo-i1 mbar are routinely achieved and the background pressure of the chamber never exceeds 3 x 10e5 mbar during growth. The gas sources are highly colli-

of Si/Si,

$e,

mated, forward peaked beams of Si,H, (Air Products, 99.99% purity) and GeH, gas (Union Carbide, 99.995% purity) with the source slits located symmetrically 50 mm from the substrate surface. RHEED studies were carried out using a VG LEG 110 electron gun (energy = 15 keV) mounted on a manipulator equipped with a stepper-motor. The latter feature has enabled us to have a precise control over the polar angle of the incident electron beam which is important in studying RHEED intensity oscillations and diffraction patterns [19,20]. Two inch Si(OO1) singular (+0.5”) substrates were prepared by a modified [213 Shiraki 1221 etch prior to loading into the system. The resulting thin surface oxide film was thermally removed by heating the sample at 825°C in a low flux of Si,H, which has been shown previously to aid in the removal of the surface oxide [23,24]. To eliminate the initial surface roughness and any residual cont~inat~n problems, a buffer layer of approximately 1000 A was grown at 725°C. After the surface cleaning and buffer layer deposition, a RHEED pattern consisting of well-developed diffraction features indicating a two-domain (2 x 1) + (1 x 2) surface structure was observed. Intensity oscillations from the diffraction features along a [OlO] azimuth were measured directly from the RHEED screen using a fiber optic and photomuItiplier tube (PMTj. The output of the PMT was recorded directly onto an X-Y plotter. Details of the Si substrate preparation methods and the RHEED arrangement have been reported elsewhere [21]. The growth experiments proceeded as foliows: After the buffer layer deposition, 500 A of Si,_,Ge, was grown at 560°C. Alloy compositions were determined ex situ by TEM/EDX (transmission electron microscopy/ energy dispersive X-ray analysis>. The substrate was then flashed to 800°C for 5 min immediately prior to Si overfayer growth. This was done to promote surface relaxation as well as remove surface hydrogen [ZS], leaving a well-ordered Si,_,Ge, alloy surface. The annealing made the RHEED intensity oscillations during the subsequent Si overlayer growth more distinct, but induced no observable change in the oscillation periods. Following annealing,

N. Ohtani et al. / Ge segregation during Si-GSMBE

TIME (set) Fig. 1. RHEED intensity oscillations observed in the specularly reflected beam along the [OlO]azimuth during Si growth on Si,_,Ge, at a substrate temperature of 560°C; a constant Si,H, flux of 3.4X1O’5 molecules cm-* s-l was used. The gradual change in GR is attributed to Ge segregation to the surface layer. The inset shows the decrease of the growth rate change (0) and the fit from eq. (10) (solid line).

the substrate temperature was lowered to the overlayer growth temperature, and the Si layers were grown on the Si,_,Ge, alloy while the RHEED intensity oscillations were monitored.

3. Results and discussion 3.1. Obseruat~onof Ge sur$ace segregation using RHEED Fig. 1 shows RHEED intensity oscillations of the specular beam obtained in the [OlO] azimuth during Si overlayer growth on Si,_,Ge, at 560°C using a SiBH, beam flux of 4.6 x 10” molecules cm2 s-l, from which it is apparent that the period of the oscillations increased gradually during the Si overlayer growth. An accompanying increase of the absolute diffraction intensity was also observed. The frequency and intensity changes appear to saturate after w 20 monolayers (MLs) of Si growth. The gradual frequency change indicates a decreasing growth rate (GR) as Si is deposited and following N 20 MLs of Si overlayer growth the GR approaches the typical Si on Si homoepitaxial value observed at 560°C. This transient growth rate behaviour was not

ofSi /Si,

_ xGe,

301

observed at substrate temperatures above 600°C, where the reaction mechanism changes from hydrogen desorption to hydride adsorption/ dissociation limited kinetics as observed previously during Si-GSMBE growth from Si,H, [26]. In fact, the change in the GR in the transitional region, between 52O”C-6OO*C,shows a very strong temperature dependence [27]. The change in growth rate (AGR) during Si overlayer growth beiow 600°C can be attributed to a gradual decrease in the number of centres from which hydrogen can desorb easily, which have been equated with Ge surface atoms resulting from Ge segregation effects which are known to occur in Si/Si,_,Ge, heterost~ctures f2-71. During Si homoepitaxy from hydride sources at growth temperatures below 6OO”C,it is generally believed that surface hydrogen inhibits pyrolysis by blocking adsorption sites, thus limiting the Si growth rate. A number of studies [28-301 have shown that the activation energy, EA, for low temperature Si growth from hydrides closely resembles that found by Sinnah et al. [31] for the desorption of hydrogen from Si(OO1) surfaces, 47 kcal/mol. On the other hand, a lower EA of 35 t 5 kcal/mol for hydrogen desorption from Ge(OO1) and (111) surfaces was measured by Sumev and Tikhov 1321, and, consequently, Ge atoms on the growing surface should act as efficient hydrogen desorption centres. We have also reported a lower I?, (35 kcal/mol) during Si,_,Ge, heteroepitaxial growth than during Si homoepitaxial growth (40.7 kcal/mol). Thus, Ge relaxes the hydrogen blocking effect of the surface adsorption sites, resulting in a growth rate enhancement in the Si overlayer growth on Si,_,Ge,. The gradual recovery to the Si homoepitaxial growth rate after commencing the overlayer growth corresponds to the Ge profile, which is smeared out over about 20 MLs by Ge segregation. Our observation demonstrates that RHEED intensity oscillations are a very powerful tool for studying the Ge segregation effect in Si/Si, _xGe, heterointerfaces. By analyzing the gradual recovery in detail, we can achieve, with monolayer resolution, an in situ investigation of dynamical aspects of the Ge segregation.

N.

308

Ohtani et al. / Ge segregation durhg Si-GSMBE

3.2. Mode&g

fi, _ *Ge,<

requires that we explicitly include the following conditions: (1) The constant k must be expressed in terms of the growth rate, i,e., k(GR). (2) The experimental conditions used here (GR < 0.2 ML/s and Tg = 560°C) result in segregation which occurs in the equilibrium or steady-state regime, as described previously by Harris et al. 1331. Fig. 2a represents schematically the two-siteexchange model introduced by Ilarris et al. [33]. We can distinguish the two regions; 6) equilibrium (or steady state) where segregation is fast relative to the growth rate so that an equilibrium (steady state) dist~butio~ of se~egatiug material is established between surface and subsurface layer during growth; (ii> ~netically-limited, where segregation is slow relative to the growth rate so that a steady state distribution is not achieved. (i> For the steady state case, assuming m = 1 and ak/GR * 1, the incorporation constant is given

In general, the dynamic surface segregation process observed for Si, _XGeX can be represented by a mass balance equation of the form t33f

-d&k =J-k(N,,)m-k&jN&, dt

where h& is the number of Ge atoms per unit area of surface, and J is the nominal incident flux of Ge from GeH,. k(N,,)” represents the rate of i~co~ration of Ge into the growing film, of order m with respect to the Ge concentration, and kdNGe represents Ge devotion. Under the experimental conditions used here, J = 0 and assuming there is no Ge desorption from the surface at these temperatures [34], k, = 0. Eq. (1) therefore reduced to:

In situ measurements of the RHEED intensity oscillations reveal that the growth rate, GR, changes doing growth. This unusual behaviour

(3)

b

a

Subsurface

of Si /

Surface

Fig. 2. (a) Schematic energy diagram of the ~~site”~cha~~e model. E, is the kinetic barrier height and AC;,,, is the free energy difference between the surface and subsurface. The desorption energy, Ed, is much larger than both Et and AGs,,. Each minimum corresponds to a point on the Si lattice, with a spacing a = a,/4 in the [Ool] direction. tb) Temperature dependence of the f/e decay length. The Ge segregation undergoes a transition from an equilibrium to a kineticai~-limited regime as the temperature is decreased. In the equilibrium regime, segregation is enhanced as the temperature is lowered, while, in the kinetically-limited regime, segregation decreases as the temperature is lowered.

N. Ohtaniet al. / Ge segregationdwikg Si-GSMBE of Si / Si, _ xGe,

where Q is the ML thickness and (b is the steady state ratio of Ge between surface and subsurface sites and can be expressed in terms of the Gibbs energy difference AG,, which provides the driving force:

309

where CYand y are constants, eq. (‘7) can then be rewritten as d AGR -=; dx

-y;AGR,

(9

so that += ( z).,y=kl

exp( -%),

(4) AGR(x)

=B exp (

where A is a constant. Profiles will be independent of growth rate. (ii) For the case where the process is kinetically limited: k(GR)

where T is the ML growth time and P the transition pr~babili~ between surface and subsurface sites. It can be expressed in terms of a barrier height E, (see fig. Za): P=P,exp

(-2 1,

(6)

where PO is a constant similar to A in eq. (4). In this regime se~egation profiles will be dependent on the growth rate. The temperature dependence of segregation in the two regimes is shown ~hematic~ly in fig. 2b. Here the segregation can be characterised by the l/e decay length (0 In the equilibrium regime, enhanced segregation is obtained as the temperature is decreased as a result of the ~l~rnan~ factor in eq. (4). At lower temperatures, in the kinetically-limited regime where eq. (5) describes the segregation beha~our, segregation decreases as the temperature is decreased due to the decrease of P. There is a transition from the kinetic~ly-l~ted to the eq~ibrium regime, with a m~mum segregation decay length. Substituting eq. (31 to eq. (2), we obtain (assuming m = 1) dNGe -=--

(b

N Ge’ a dx If the observed change in growth rate is a function of the surface Ge ~n~ntration, i.e. AGR=LY(N&)~,

(8)

y$x1 ,

where B is a constant.

W

We have found that this simple exponential form fits well the observed AGR decay seen in fig. 1, as shown in the inset in fig. 1. Furthermore, all data reported in subsequent sections can also be fit by this exponential equation, imply~g the assumptions made (m = 1 and AGR N (NoJy> are valid for the range of conditions used here. 3.3. Ge concentration dependence To examine the influence of the Ge concentration on the segregation, we monitored the transient growth rate changes during Si overlayer growth at constant temperature on Si,_,Ge, alloys of different Ge film com~sitions (x = 0.01, 0.02, 0.035, 0.045, 0.07); it should be noted that the actual surface Ge concentrations, N& are expected to be higher than the film ~m~sitions due to segregation. Growth rate changes (AGR) of the Si overlayer on each starting surface calculated from the period of the RHEED intensity oscillations are shown in fig. 3. In the figure, it can be seen that the exponential decay (eq. (101) fits well the observed data on each Si,_,Ge, surface. The decay length 0) is constant throughout whole the concentration range and calculated to be h = 14 A. Thus, the observed curves are simply shifted parallel to each other. In fig. 4 we plot the initial growth rate change of Si on Si,_,Ge, uersm the alloy composition of the starting films. As seen in the figure AGR is propo~ional to x. In a previous study 1271 we found that Ge segregation does not result in a surface Ge saturation at the Ge concentrations used here, and it is therefore assumed that x is pro~~ional to A&.. Hence, AGR is also assumed to be proportional to No@ implying that y = 1 in eq. (8).

310

N. Ohtani et al. / Ge segregation during Si-GSMBE of Si / Sit _ $Ik,

0

. a

* ‘

_..--. “^--.--

CR x=0.01 CR x=002

.“-- “-.

Si on SiGe at 560 T

n=oox

CR GR x=O.OJS CR n=om

/

--------.-

D

Fig. 3. Growth rate decay (AGR) measured during Si overlayer growth on Si,_,Ge, alloys of varying Ge composition (x = 0.01, 0.02, 0.035, 0.045, and 0.07). The simple exponential form (solid line), eq. (lo), fits well the observed decay of AGR over the whole range of the Ge film compositio?, providing the common characteristic decay length of h = 14 A.

Recent XPS and SIMS studies [34,35] reveal that Ge segregation in Si/Si, _,Ge, does not fit a simple exponential form over a large range of Ge concentration and higher order terms Cm > 1) are required to fit the observed decay. Fukatsu et al. measured, by SIMS, Ge segregation profiles over a wide range of Ge concentration (~.l%-lo%)

of/, o

’ 0.01

r . i . ’ r ’

0.02

0.03

Film

0.04

0.05

I

0.06

’

’ . ’

0.07

0.08

Com~sition x

Fig. 4. The initial growth rate enhancement of Si on SiI_XGe, versus the alloy composition (x) of the starting film. Assuming that x is proportional to Noe, AGR is well described by eq. @with y=l.

IO

10

,111

Number of Monolayer Fig. 5. Normalized Ge surface con~ntration change as a function of the Si overlayer growth on Si,,,,Ge,,O,. Note that the segregation decay curve shows a change in slope following - 30 ML deposition of Si.

and found that the decay profiles change slope near 1% Ge composition, indicative of the existence of two concentration regimes which have different ~hara~te~stic decay lengths. Accounting for bond geometry, which results in two equivalent segregation paths in the [OOll direction, they attributed this non-exponential behavionr to a self-l~itation effect of Ge segregation [361, which arises from the presence of Ge along the segregation path, thus hindering further segregation. In our case, we are actually measuring a GR change which is attributed to an enhanced hydrogen desorption rate due to the presence of surface Ge. Due to the large number of parameters in our experiment, we hesitate to discuss the origin of the observed nonlinear behaviour beyond this at the present stage of this research. In order to shed light on this issue, we obtained prolonged RIIEED intensity oscillations by optimizing the growth surface and the diffraction conditions. The result is shown in fig. 5. In this experiment we again used a Si,.93Ge,,, alloy layer as a starting surface for Si overlayer growth and the growth procedure was the same as in fig. 4. Despite the scattered data in the lower concentration range, it can be clearly seen that the segregation decay curve shows a change in slope following N 30 ML deposition of Si. The decay

N. Ohtani et al. / Ge segregation during Si-GSMBE of Si /Si, _ xGex

lengths before and af:er the break are calculated to be A = 14 and 70 A, respectively. These observations are consistent with the SIMS results reported by Fukatsu et al. [35] and the decay lengths obtained here are quite similar to those observed by them during solid source MBE at Tg = 380°C. They attained x = 0.01 as the Ge concentration for which the transition from fast to the slow decay region occurred. If we assume that in our experiments the change in the decay length also occurs at x = 0.01, then for the Si,,,Ge,,, alloy layer used as a starting surface in fig. 5 approximately 10% is obtained by extrapolation as the actual Ge surface concentration. This is about 50% more than the film composition, and is consistent with our previous studies which show that segregation does not result in a saturated Ge surface concentration. Although this value appears reasonable, a full evaluation of Ge surface concentration is still required. 3.4. Temperature dependence Fig. 6 shows the temperature dependence of the Ge segregation profiles. Here, we fixed the Ge film composition at x = 0.07 and the sub-

0

5

IO

IS

20

25

Number of Monolayer

Fig. 6. Temperature dependence of the Ge segregation during GSMBE. The segregation is drastically suppressed as the substrate temperature is lowered. At 54o”C, the decay length in GSMBE (A = 11 A) is comparable to that at 150°C in solid source MBE. The data for solid source MBE (solid lines) follow Nakagawa and Miyao (ref. [34]).

311

strate temperature during Si overlayer growth was changed; Tg = 560, 540, 520°C. In fig. 6, the Ge surface concentrations are calculated from the growth rates of the Si overlayer at each substrate temperature using eq. (8) and y = 1. In the equilibrium regime, as shown in fig. 2b, segregation is expected to be enhanced as the temperature is lowered due to the decrease of the Boltzmann factor in eq. (4). However, as seen in fig. 6, the Ge segregation is actually suppressed at lower temperatures. This apparent inconsistency may be understood in terms of the recently reported “surfactant” effect of surface hydrogen, which is produced by the dissociation of S&H, and GeH, and passivates the surface during growth when Tg < 600°C.

Cope1 and Tromp [14] have found that surface hydrogen serves a beneficial role in maintaining the abruptness of the Si/Si,_,Ge, interfaces. They measured, by MEIS, a probability of Ge “floating” (Pa,,,,) towards surface during MBE, CVD, and surfactant assisted growth (SAG) using Sb. During each growth experiment the surface was terminated either with dangling bonds, hydrogen, or antimony, respectively. For structures grown by MBE, abrupt interfaces could only be obtained when Tg =g250°C. However, low temperatures resulted in poor crystallinity of the material. On the other hand, for CVD grown structures at Tg = 52o”C, they could observe little evidence of segregation in the MEIS spectrum. At a slightly higher temperature, an asymmetric broadening of the Ge peak in the spectrum could be seen, indicative of segregation into the Si capping layer, and finally, at Tg = 590°C not only does the Ge peak become noticeably broader, but significant quantities of surface Ge, arising from strong segregation, were observed. They concluded that the surface hydrogen acts as a surfactant at low temperatures (T, I 52O”C), inhibiting the interchange of Si and Ge at the growth front, and the dominant factor in determining the interface abruptness during CVD growth was the quantity of the surface hydrogen during growth. To prepare abrupt structures at temperatures above 520°C they found that the surface must be terminated with a “surfactant” such as Sb, which does not desorb as readily as hydrogen.

312

N. Ohtani et al. / Ge se~egat~on during Si-GSMBE

l/T Fig. 7. Schematic representation of the temperature dependence of the l/e decay length (A) of the Ge segregation during GSMBE. In the equilibrium regime, as the temperature is decreased, A is expected to increase along the equilibrium curve. At the same time, however, the reduction of AG,, due to the increase of surface hydrogen decreases the slope of the equilibrium curve, resulting in the temperature dependence indicated by the dashed line.

Surface hydrogen may reduce the free energy difference AG,,, between the surface and subsurface (see fig. 2a). This reduction of AG,,, may compensate the temperature effect, thereby resulting in the observed suppression of the Ge segregation as the temperature is decreased. This effect is shown schematically in fig. 7. As the temperature is decreased, the segregation decay length, A, increases along the equilibrium curve shown in fig. 2b. At the same time, however, the quantity of surface hydrogen, which passivates the growing surface, increases at lower temperatures, causing the reduction of Ah(c;,,,which leads to the decrease of the slope of the equilibrium curve. Although these two effects act in opposite directions, the observed suppression of the segregation at lower temperatures suggests that the latter effect is dominant and results in the temperature dependence shown by the dashed line in fig. 7. In fig. 6, for comparison, the results of Si/Si i _xGex growth from conventional sohd source MBE (SSMBE) by Nakagawa and Miyao [34] are also shown. In their experiments, Si overlayers were deposited on three monolayers of Ge and the segregation profiles were examined ex

of Si / Si, _ xGe,

situ by XPS. For comparison with our results, their data are normalized to the same concentration range. It appears in the figure that the Ge segregation is drastically suppressed in GSMBE compared to SSMBE, indicative of a crucial role of surface hydrogen as a surfactant. The decay length at 540°C in GSMBE (A = 11 A) is comparable to that at 150°C in SSMBE. This result is consistent with the MEIS results obtained by Cope1 and Tromp [141, suggesting that alloy growth from hydride sources may have an inherent advantage over evaporation sources in providing sharp heterointerfaces. 3.5. Growth interruption and annealing Generally the diffusion of Ge in bulk Si is believed to be negligibly small at typical MBE temperatures [37], but the strong affinity for Ge to segregate to the surface during growth could result in enhanced outdiffusion from layers in the near surface region. To examine the thermal stability of Ge segregated surfaces, growth interruption and annealing experiments were performed. The experimental procedure is essentially the same as that in figs. 4 and 5, however, the Si overlayer growth was interrupted following approximately 10 MLs of growth. The substrate was then annealed at 800°C for 2 min, cooled back to the appropriate growth temperature and growth resumed. The results are summarized in fig. 8 for growth at 520 and 560°C. As seen in the figure, growth interruption and annealing does not affect the growth rate and, hence, the Ge surface concentration decay profile at Tg = 560°C. Growth interruption and anneaiing at 52O”C, however, appears to increase the Ge surface concentration which results in a break in the decay curve following the annealing process. Continued growth at 520°C shows that although the Ge composition has increased, the decay profile has the same slope as before the annealing, indicating that the Ge incorporation rate during growth is immediately re-established regardless of the initial Ge composition. This is consistent with the data in fig. 3. This increase of surface Ge seen following annealing reflects the instability of the surface composition which is established during growth at

h? Uhtaniet al. / Ge segregationduring Si-GSMBE of Si/Si,_

xGe,

313

from the surface. Whether it is the hydrogen or germanium surface concentration which controls the outdiffusion effect however remains unclear, and further investigation is required,

4. Conclusions

0

10

20

30

Number of Monolayer Fig. 8. The Ge segregation curve allowing growth interruption and annealing. The arrow indicates the point where growth was interrupted and the sampIe annealed (800°C 2 mm). At 560°C no change is observed in the Ge surface concentration (Noe) following interruption and annealing. At 52O”C, however, a marked increase in NGe is measured following growth interruption and annealing. The inset shows that growth interruption and annealing in the lower NGe range, during 560°C growth, results in a small increase in NGer indicative of the outdi~sio~ of Ge.

this temperature. By annealing, this instability is thermally overcome and the surface Ge concentration increases. The observation of this effect at 520°C but not at 560°C indicates that the Ge surface concentration established during growth at 560°C is already high, as seen by the longer decay length, and any further outdiffusion is suppressed. However, annealing the sample grown at 560°C when the surface Ge composition is low, i.e. after 25 MLs of growth, a small but reproducible increase of surface Ge can be observed, as seen in the inset of fig. 8. Removal of surface hydrogen, which suppresses Ge segregation, during the annealing step presumably influences this process and hence, the increase observed when the sample grown at 520°C (a large hydrogen coverage) is annealed and not observed when annealing the sample grown at 560°C (less hydrogen coverage) may simply be due to the removal of the surfactant

We have investigated Ge segregation during Si-GSMBE by in situ growth rate measurements using RHEED intensity oscillations. The presence of Ge at the alloy surface increases the heteroepit~al growth rate, due to the increased desorption rate of hydrogen from Ge rather than Si at the alloy surface. As Ge segregates to the growth interface RHEED intensity oscillations show a marked change in growth rate. We have modelled the Ge segregation by monitoring the change in growth rate and have shown that the observed growth rate enhancement can be used as a direct measure of the segregation. The surface sensitivity of RHEED makes this technique an excellent probe for high resolution studies of dynamical processes at the growth front. Here, using this novel technique, concentration and temperature dependent results of Ge surface segregation were presented. In the concentration dependence studies the segregation decay curve was found to have a point of infection, where the segregation behaviour shows a transition from a fast to a slow decay mode. This result is in a good a~eeme~t with the nonlinear behaviour of SIMS and XPS profiles recently reported and indicates that the segregation proceeds in a different fashion between these two concentration ranges. Temperature studies reveal that the Ge surface segregation can be minim~ed below 600°C in gas source MBE. This is indicative of a crucial role of surface hydrogen, which is produced by the dissociation of disilane and germane on the surface during Si-GSMBE. This “surfactant” effect of surface hydrogen becomes prominent below 56O*C, and the characteristic segregation decay length at 540°C of h = 11 w is almost equivalent to that at 150°C in solid source MBE. These results are very consistent with the MEIS results recently obtained by others and hold a great

promise for a more precise control of the interface. FinaIly, the effects of growth inter~ptio~ during the surface segregation and the thermal stability of Ge segregated surfaces were examined. Growth interruption and annealing during Si overlayer growth on Si,_,Ge, resulted in a small increase in the surface Ge concentration. This is attributed to the outdiffusioR of Ge from the very near surface region, which is another consequence of surface energetics in this system.

Acknowledgements The authors would like to thank C. Roberts for preparing the substrates, X. Zhang for TEM measurements, and J.H. Neave for his useful insight.

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