GaAs heteroepitaxy characterized as a stress-free system

434

Applied Surface Science 50 (1991) 434-439 North-Holland

G a S b / G a A s heteroepitaxy characterized as a stress-free system C l a u d e Raisin

a,

A n d r 6 R o c h e r b, G e o r g e s L a n d a c, R o b e r t Carles c a n d L o u i s L a s s a b a t e r e a

a Laboratoire d'Etudes des Surfaces, Interfaces et Composants (LES1C), Universit@ des" Sciences et Techniques du Languedoc, F-34095 Montpellier, France b Centre d'Elaboration des Matdriaux et d'Etudes Structurales ( C E M E S / L O E ) , CNRS, 29 Rue Jeanne Maroig, F-31400 Toulouse, France c Laboratoire de Physique des Solides, Universitd Paul Sabatier, F-31062 Toulouse, France Received 27 November 1990; accepted for publication 10 December 1990

Thin layers of GaSb grown on G a A s by molecular beam epitaxy have been studied by R a m a n spectroscopy and transmission electron microscopy. In spite of the high value of the lattice mismatch (8%), these systems reveal a good crystalline quality: as a matter of fact, the lattice dynamics analysed through the resonant R a m a n scattering lead to a large correlation length for the optical modes. Moreover, they evidence the unstrained nature of the epilayers. The absence of the residual strain is connected to the highly regular network of Lomer dislocations observed by TEM. Obtaining such a perfect array of dislocations is a consequence of both the characteristic of the system (high lattice mismatch, same cationic sites) and the peculiar growth conditions (ideal surface preparation due to the buffer layer, adequate temperature allowing the direct relaxation of GaSb by island growth).

1. Introduction

The crystal growth behaviour of the G a S b / GaAs system exhibits interesting features, associated with large lattice mismatch heterostructure. The density of threading defects has been measured at the level of the interface, and is at least two orders of magnitude smaller than in the 4% mismatched G a A s / S i system [1]. In this article, we shall explain this result in terms of the relationship between the crystalline quality of the GaSb epilayer and the peculiar nature of the G a S b / G a A s interface. GaAs and GaSb both have the cubic zincblende structure, with the space group F43m. The lattice parameter of GaAs is 5.641 A, whereas that for GaSb is 6.094 A, giving a lattice mismatch of 8% between the two crystals. In this case, the useful deposited layer is always much thicker than the critical thickness and the system relaxes through dislocation mechanisms. Dislocations appearing in such heterostructures can be classified as follows: (i) Misfit dislocations, accommodating the lattice mismatch. Their interfacial density is of

the order of 8 x 1 0 6 c m - 2 to obtain an unstrained structure. (ii) Threading defects propagating parallel to the growth direction into the epilayer. They have been acknowledged to impair both the electrical and the optical properties of semiconductor epilayers. Their surface density is still as high as 1 0 6 cm -2 even with a 3 /~m thick GaAs layer grown on Si [2,3]. The reduction of this threading defect density in the active layer constitutes one of the main objectives for epitaxial growth techniques such as MBE or OMVPE. A detailed understanding of the nature of the misfit dislocations and of the structure of interfaces is needed to control and reduce the density of these unwanted defects. The crystalline quality of GaSb films is evaluated here by Raman spectroscopy and the interfacial defects are characterized by transmission electron microscopy. The growth of GaSb is processed in two steps: a 1 /~m thick buffer layer of GaAs is first grown by MBE at 580°C. Its role is to give the best possible surface quality of the GaAs substrate. A thin layer of GaSb, 3 or 25 nm thick, is then deposited by MBE at the homoepitaxy tempera-

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C. Raisin et al. / GaSb / Ga,4s heteroepitaxy characterized as a stress-free system

ture, 4 7 0 ° C [4]. These thicknesses are considerably larger than the critical thickness ( < 1 nm) and are then suitable to give significant results concerning the misfit dislocation network already observed for layers of 0.6 nm nominal thickness

[1].

/I

Raman spectra were recorded at room temperature in the backscattering geometry using the 488 n m ( E i = 2.54 eV) line of an Ar + laser. Due to the high value of the GaSb absorption coefficient, the scattering depth is limited to 8.5 nm; moreover, as the incoming photon energy lies in the vicinity of the E 1 + A1 transition (2.49 eV), the Raman efficiency is resonantly enhanced. Both phenomena then allow efficient local probing of the epilayer. For calibrating the Raman data (peak position, linewidth), we used a bulk GaSb sample as a reference. The spectra displayed in fig. 1 only concern the

= 488 nm

GaSb(3nm) I I

r--

I,

I1'

tll,o

r0

t

"

---.;_ 200

, 250

'07

.... 2 300

WAVENUMBER (cm -1) Fig. 1. R a m a n spectra of two heterostructures (3 and 25 n m thick) of G a S b / G a A s ( 1 0 0 ) in comparison with a bulk GaSb spectrum. The inset shows the GaSb acoustical second-order range for the 25 n m thick epilayer.

GaSb/GaAs (100)

]1 N

(25 nm)

h-488 nm J

P-~2OOmW

II

g

2. Raman spectroscopy

435

T =300K

,o//co

,o

l

7 U.I I'--

z_ z< <

e~

z (xY)2 200

250

300

WAVENUMBER (cm -1) Fig. 2. Raman spectra recorded from GaSb/GaAs(]O0) in three different polarization configurations. X, Y and Z refer to (110), (110) and (001), respectively.

optical frequency range. They evidence signatures of both the GaSb epilayer (220-240 cm -1) and the GaAs substrate (250-300 cm-1). First, the observation of the two well-defined peaks assigned to transverse and longitudinal optical (TO and LO) modes of GaSb, the dominance of the LO structure expected in a (001) backscattering geometry (see GaAs substrate signal), testifies to the high crystalline quality and good orientation of such ultrathin epilayers. Moreover the correlation lengths of vibrational modes are sufficient to allow clear observation of secondorder scattering as shown in the inset in fig. 1. These points have been corroborated by wellobeyed selection rules observed in polarized Raman spectra analysis. Fig. 2 presents, as an example, the Raman data recorded from the 25 nm thick sample in three different configurations: the scattering rate of the LO modes follows the expected extinction rules in the crossed configuration. Second, no frequency shift of the GaSb modes is observed by comparison with the reference signal, thus revealing the absence of any strain in the

436

C. Raisin et al. / GaSb / GaAs heteroepitaxy characterized as a stress-free system

epilayers at room temperature unlike in other heterostructures (as G a A s / S i ) [5]. As was previously shown [5], the residual strain ell at room temperature Ta is given by: =

+

f ZX (T')

dT',

where the first term accounts for the lack of complete relaxation of the lattice mismatch (Aa/a) at the growth temperature Tc and the second term for the thermoelastic contribution, Aa being the difference between the thermal expansion coefficients. Using C%asb= 6.8 X 10 - 6 K -1 and a~aAs = 5.7 X 10 - 6 K -1 [6], the latter contribution takes the value 0.5 x 10-3. The value of ell can be deduced from the LO singlet relative frequency shift by [7]: A~o ~s

1 6(sll + s12 ) [ (sll + 2s12)(/~11 + 2/('12 ) -- (Sll -- S12)(/~11

-- /~12)] '11,

where the sij are the elasticity constants and /£u the reduced phonon deformation potentials. Using the corresponding GaSb values [8] we obtain: Ao~Lo(Cm -1) = --3.7 X 102~11. Since

the

experimental

uncertainty

in

the

frequency position is lower than 0.2 c m - 1 , we can estimate that I~tl(Ta)l is lower than 0.6 x 10 -3 and IclI(TD I lower than 1.1 X 10 -3. The relaxation of the lattice mismatch strain is then about 99% during growth. However some unexpected results have been encountered in the R a m a n investigation. In particular, the observation of the substrate signal through the 25 nm epilayer (i.e. thicker than Raman depth), the activation of the T O modes (specially for GaAs) and the non-equivalent scattering efficiency in the ( Y Y ) and ( X X ) configurations (see fig. 2) required further T E M investigations.

3. TEM observations

The G a S b layers thinner than 80 nm are suitable for easy preparation of plan-view samples and for weak-beam T E M observations. This preparation needs only the chemical etching of the rear side of the substrate because the G a S b film is already thin enough to give G a S b / G a A s samples transparent to the electron b e a m accelerated at 200 kV. Such observations, with the electron b e a m parallel to the growth direction [001], give a twodimensional image of the interface. They furnish

t

Fig. 3. Bright-field image of the 3 nm GaSb film deposited on GaAs: note the amsotropic shape of the island characterized as a stress-free system by the perfect moir6 patterns.

C. Raisin et al. / GaSb / GaAs heteroepitaxy characterized as a stress-free system

more complete information on the defects than those given by (110) cross-sections where the (001) interface is seen as a line and only the set of (110) dislocations parallel to the electron beam are seen as dot contrast. With the nominal thickness deposition smaller than 30 nm the G a S b film is not continuous. As revealed in fig. 3, the 3 nm thick layer of G a S b consists of isolated islands with a wide range of sizes. The island shapes are well-defined by crystalline elements, particularly the (111) planes as facets. The elongated shape of the islands can be related to an anisotropy of the growth rates of the island facets. The coverage varies with the nominal deposited thickness. It appears that the island height rapidly reaches a m a x i m u m value dictated by the growth conditions. Increasing the deposited G a S b thickness .results in an increase in the lateral island size, together with some coalescence of neighbouring islands. The latter results in irregularly shaped islands. Similar results have been

437

obtained on I n S b / G a A s by Zhang et al. [9]. This observation confirms the 3D growth process of GaSb on G a A s observed by in-situ Auger measurements in the first stage of the epitaxy. The misfit dislocation networks have been observed using an unconventional weak-beam technique discussed elsewhere [1]. Weak-beam images are better when the specimen is very thin. In fig. 4, the dislocations are seen as straight lines along [110] and [110] with a Burgers vector perpendicular to the dislocation line and contained in the (001) interface, which are often called Lomer dislocations. Their average spacing is measured to be 5.4 + 0.5 nm, in good agreement with the calculated value: 5.365 nm. They are organized as a periodic array very close to the ideal case. The two families of Lomer dislocations cross each other mainly without interaction creating threading defects. It appears that the threading defects are only related to the imperfections of the misfit dislocation network as observed in fig. 4 for the

Fig. 4. 2D weak-beam image of G a S b / G a A s interface for a 25 nm film of GaSb. Note the perfect array of square misfit dislocations. Threading defect marked D appears at the imperfections of the misfit dislocation network.

438

C. Raisin et al. / GaSb / GaAs heteroepitaxy characterized as a stress-free system

defect marked D. Finally, the high quality of this network explains well the low density of threading defects in the G a S b epilayers.

4. Discussion The T E M observations and R a m a n spectroscopy data show that the GaSb films are stress-free with a very low density of threading defects. This result is directly related to the high quality of the interfacial dislocation network. In the G a S b / G a A s system, the creation of L o m e r dislocations appears to be a mechanism directly related to the 3D growth which relaxes the misfit strain by the dislocation network. The (111) facets of the islands play a major role in the creation mechanism of Lomer dislocations. These facets and the interface make an edge parallel to a (110) direction. The line defect is then formed at the periphery of the island [10]. This process is very likely since, as discussed by Kiely et al. [11], it can occur without the coordinated motion of a large number of atoms and it is not subject to an energy formation barrier. Subsequent dislocations form at approximately even intervals as the strain in the island periodically accumulates. This perfect relaxation is also directly related to the surface quality furnished by the GaAs buffer layer. The second important parameter in the growth process is the temperature, which is not taken into account in the models of critical thickness. The G a S b homoepitaxial temperature is large enough to enable a direct plastic relaxation of the deposited film. The main difference between G a S b / G a A s and G a A s / S i is related to the quality of their misfit dislocation networks. All misfit dislocations, studied in these heterostructures, have been characterized as Lomer type. They are obtained either directly during the growth as observed in G a S b / G a A s or after annealing for G a A s / S i ; the creation mechanism should be different in the two systems. In the case of G a A s / S i , a two-step growth process has been developed to improve the crystalline quality of the G a A s layer. The first step was performed at low temperature to ensure enough

nucleation to cover the entire Si surface with a uniform G a A s layer with poor crystalline quality. The second step is performed at higher temperature which is more conventional for G a A s epitaxy. Under these conditions, G a A s becomes plastic and the misfit strains are relaxed mainly by Lomer dislocations. The dislocation network is then reorganized into small 2D sub-arrays of Lomer dislocations with a mean size related to the growth temperature. The junctions between the dislocation sub-arrays are responsible for a large density of threading defects. Nevertheless, the creation mechanism of Lomer dislocations is not well understood. The driving force involved in the G a A s / S i system is probably related to an effect of recrystallization. 5. Conclusions The R a m a n scattering experiments discussed in this paper show that the G a S b is relaxed more than 99% even for very thin deposited layers. Similar results have been obtained by Bourret et al. using X-ray analysis [12]. We have seen that the threading defects are directly related to the imperfections of the interfacial misfit dislocation network. Their density is much lower in G a S b / G a A s than in G a A s / S i due to the difference in quality of their misfit dislocation network. Contrary to c o m m o n notions, there is then no direct correlation between the value of the mismatch and the density of threading defects. The Lomer dislocations, especially when they are developed with a well-defined mechanism as observed in G a S b / G a A s , are useful for complete relaxation of the stress due to the lattice mismatch and there is no other fundamental reason for creating threading defects in the epilayer. These results furnish some indications concerning the origin of both the misfit dislocations and the threading defects. In some special cases it would be possible to define experimental conditions that considerably reduced the number of threading defects in the epilayer. When the misfit dislocation network is perfectly organized at the interface, the stress due to the lattice mismatch is well relaxed and the density of threading defects can be very low.

C. Raisin et al. / GaSb / GaAs heteroepitaxy characterized as a stress-free system

Acknowledgements T h e a u t h o r s a r e v e r y g r a t e f u l t o J. C r e s t o u f o r specimen preparation for TEM observations. T h i s p a p e r is d e d i c a t e d t o t h e m e m o r y o f o u r friend Claude Raisin.

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