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Geometry and interface structure of island nuclei for GaSb buffer layers grown on (001) GaAs by metalorganic vapour phase epitaxy
Journal of Crystal Growth 133 (1993) 168—174 North-Holland
~
CRYSTAL GROWT H
Geometry and interface structure of island nuclei for GaSb buffer layers grown on (001) GaAs by metalorganic vapour phase epitaxy M. Aindow
a T.T. Cheng a N.J. Mason b T.-Y. Seong b and P.J. Walker b School of Metallurgy and Materials, The University of Birmingham, Elms Road, Edgbaston, Birmingham B15 217’, UK b Physics Department, Clarendon Laboratory, The University of Oxford, Parks Road, Oxford OX] 3PU, UK a
Received 22 May 1993; manuscript received in final form 2 July 1993
Atomic force microscopy and transmission electron microscopy have been used to investigate the geometry and interface structure of island nuclei formed in the initial stages of buffer layer growth for MOVPE GaSb on (001) GaAs. There is a bimodal distribution of island sizes with a high density of small, homogeneous nuclei and a lower density of larger, secondary nuclei. The smaller islands have pronounced crystallographic facets which are consistent with those which would be expected for minimization of surface energy and lateral growth anisotropy. The secondary islands are present at junctions between primary nuclei and may have formed due to enhancement of growth rates at emergent threading segments of misfit dislocations. The lattice misfit is accommodated by a regular square arrangement of edge-type misfit dislocations but unusual strain contrast arises in HREM images due to either a “stand-off’ of dislocations from the interface or a corrugated interface.
1. Introduction Epitaxial films of GaSb on GaAs are of considerable interest for optoelectronic applications since the heterostructure junction is rectifying and diodes produced from such structures could be integrated with conventional GaAs device architectures. The deposition of GaSb films onto GaAs by metalorganic vapour phase epitaxy (MOVPE) is usually carried out by depositing a low temperature “buffer layer” of 20—50 nm of GaSb before ramping to the appropriate deposition temperature for the remainder of the film [1]. The role of the buffer is to promote nucleation and to ameliorate the effects of the 7.8% lattice mismatch by “turning over” the dislocations produced at the interface. In the present study we have concentrated on the initial stages of buffer layer formation, i.e. prior to the coalescence of island nuclei. The topography, size and distribution of these features have been investigated using atomic force microscopy (AFM) and the interface structure has been examined using transmission electron 0022-0248/93/$06.0O © 1993
—
microscopy (TEM). It is important to understand the nature of nucleation and growth for these buffer layers, since it may determine the types of crystallographic defects which are produced in response to the misfit. This could have important consequences for the nature and density of defects in any deposit grown onto the buffer. Moreover, the geometry of nucleation could have a greater significance; as with many semiconducting materials, GaSb shows a pronounced quantum confinement effect, thus if regular arrays of island nuclei of an appropriate size could be grown, these might be used directly as quantum dots.
2. Experimental details Gallium antimonide heteroepitaxial deposits were grown by atmospheric pressure MOVPE in a horizontal silica cell reactor with an RF-heated graphite susceptor, which is described in detail elsewhere [2]. The substrates used were Epi-ready GaAs: Si wafers with a nominal dislocation density of
Elsevier Science Publishers B.V. All rights reserved
M. Aindow eta!.
/
Geometry and interface structure of island nucleifor GaSb on (001) GaAs
tation of (001), i.e. no intentional offcut. The substrate preparation and pre-growth treatments were identical to those used previously [3]. The reactants used were trimethyl gallium (TMGa) (Epichem, 9°C)and trimethyl antimony (TMSb) (Alfa, 0°C),respectively with palladium diffused hydrogen as the carrier gas. Deposits were grown at 530°Cfor 5 mm: at this temperature, deposition rates of <1 ~sm/h are typical. Topographic measurements were obtained from these specimens with a Digital Instruments Nanoscope II contact-mode AFM using a conventional Si 3N4 with at a force constant 1 cantilever and scanning 4 Hz with of 0.06 N m The data files are measurements constant force. of specimen “height” as a function of probe position over a grid of 400 X 400 points within the scan area. Cross-sectional TEM specimens were produced by glueing two portions of the wafer together at 90° to one another and then cutting slices such that the surface normal, n, is parallel to the two orthogonal directions [110] and [110] which lie in the substrate surface. Discs cut from these slices were dimpled and ~ ion beam milled to perforation at liquid nitrogen temperatures. Plan-view specimens were produced in a similar manner and ion beam milled from the substrate side only. These specimens were exammed using a Philips CM2O TEM operating at 200 kV. —
169
±2%. There are two distinct types of island present; a large number of thin islands with an aspect ratio 5 (e.g. at A) and a smaller number of thicker islands with an aspect ratio 1 (e.g. at B). Islands of the first type are 80—110 nm thick, 180—300 nm wide and 1100—1300 nm in length. The density of these islands is difficult to estimate, but ignoring the possible effects of coalescence gives a lower bound of io~cm2. Islands of the second type are 140—250 nm thick, 350— 1000 nm wide and 900—1300 nm in length; the density of these islands is 8 x 106 cm2. We note thattwo these latterofislands always appear to straddle or more the thinner ones. Scans of more restricted areas give better resolution and reveal more clearly the effect of crystallography on the island topography. Fig. 2 is a perspective view of a surfaceplot with no exaggeration of vertical relief, constructed from a data file for a 3.25 X 3.25 jxm scan of the area marked X in fig. 1. In this case the “image” shows the shadows which would be expected if a surface with this geometry were illuminated from the left side. The thinner islands are bounded by a small upper facet which lies parallel to the substrate surface and by inclined facets on each side. Measurements from the data show that in each case, the inclined facets lie at an angle of 55°±3°to the substrate surface (we note that whilst
3. Results Topographic data were obtained from several areas on the sample using AFM and they revealed a homogeneous distribution of island nuclei with pronounced crystallographic facets. Fig. 1 is a topview “image” with shading corresponding to the height, constructed from a typical data file for a 16.5 x 16.5 ~m scan. The figure shows a large number of light rectangular protrusions on a dark background; comparing the scan axes to the {110} cleavage facets on the edges of the specimen, it is clear that these islands have edges parallel to [110] and [110]. The proportion of the surface which is covered by these islands is 83%
Fig. 1. Top view AFM image showing the distribution of GaSb islands — scan size 165 x 165 ~sm.
170
M. Aindow et a!.
/ Geometry and
interface structure of island nuclei for GaSb on (001) GaAs
Fig 2. Surface plot AFM image showing a detail of the region marked X in fig. 1 (top illuminated). Scan size 3.25 x 3.25 tim, vertical scale 0.7 urn — no exaggeration of vertical relief.
contact-mode AFM images can give considerable distortion to such angular measurements, the angles measured here correspond very closely to those measured from TEM images of cross-sectional specimens). Since the crystallographic axes of the substrate and deposit are parallel to one another, the upper facet is an (001) surface and the side facets are (111) surfaces. One example of
_________
the thicker islands is also present at Z in fig. 2. The topography is less faceted and exhibits some structure which is not well defined. We should emphasise that these are not isolated observations, data collected from various locations on the wafer revealed island nuclei with similar sizes, shapes and densities to those shown in figs. 1 and 2.
IL,
Fig. 3. Determination of polarity using CBED patterns obtained with B [270]: (a) pattern obtained with 002 at the Bragg condition; (b) pattern obtained with 002 at the Bragg condition.
M. Aindow eta!.
/ Geometry and
A
interface structure of island nuclei for GaSb on (001) GaAs
Ad
lOOnm Fig. 4. Many beam bright field TEM micrograph obtained from a plan view specimen with B
Fig. 5. Weak beam dark field TEM micrograph obtained from a plan view specimen with g
[001].
400 and B close to [001].
171
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M. Aindow et al.
/
Geometry and interface structure of island nuclei for GaSh on (001) GaAs
Fig. 6. HREM image obtained with B = [110] showing a regular array of 90°type misfit dislocations (arrowed).
The polarity of the samples was established by using convergent beam electron diffraction (CBED) from cross sectional specimens in the TEM. The patterns were indexed with respect to a unit cell with basis Ga at 0, 0, 0 and As/Sb at 1/4, 1/4, 1/4. Following the method of Taftø and Spence [4], patterns were obtained from the substrate under dynamic two beam conditions with the beam direction B [270] as shown in fig. 3. The bright cross contrast inside the 002 reflection (fig. 3a) and the dark cross contrast inside the 002 reflection (fig. 3b) indicate that this portion of the specimen has n parallel to [110]. In equivalent patterns obtained from the portions of the specimen with n parallel to [110], the cross contrast is inverted. By examining images obtamed from these specimens, it is clear that the major and minor axes of the elongated islands are [110] and [110] respectively. Thus the larger indined facets are (111) and (111) which are the Sb-terminated (111)B surfaces and the smaller inclined facets are (111) and (111) which are the Ga-terminated (111)A surfaces. Diffraction contrast TEM images obtained from plan-view specimens confirm the distribution of island nuclei indicated in AFM images such as fig. 1. Fig. 4 is a typical bright field image obtained with the specimen oriented with B = [001]. The positions of the islands are revealed by the presence of two sets of perpendicular moire fringes with a spacing of 2.9 nm. This spacing corresponds to that which would be expected for interference between the 220-type reflections in GaSb and GaAs and indicates that the deposit is
fully relaxed. In weak-beam dark field images the amplitude of moire fringe contrast is attenuated strongly and dislocation contrast dominates the images. A representative micrograph is shown in fig. 5 which was obtained with the specimen onented for g/3g with g = 400, where g is the reciprocal lattice vector of the diffracting planes. A regular orthogonal arrangement of dislocations is observed with line directions [1101 and [110], each array having a spacing of 5.8 nm. This is consistent with the arrangement that would be expected for full accommodation of misfit by arrays of edge type misfit dislocations with Burgers vectors b ~[110] and ~[110], respectively. This has been confirmed by obtaining high resolution electron microscopy (HREM) images of the crystal lattices in the interface region from cross sectional specimens. One such image, obtained with B = [110] is presented in fig. 6 with the edges of terminating half planes indicated by arrows. It is interesting to note that such images show unusual strain contrast at the interface with “cusps” at the position of the dislocations; this contrast was present for all values of defocus which gave lattice resolution. =
4. Discussion The AFM data presented here show that the number density of the thinner island deposits is five orders of magnitude higher than the stated dislocation density for the substrates. These islands must, therefore, arise from homogeneous
M. Aindow eta!.
/
Geometry and interface structure of island nucleifor GaSb on (001) GaAs
nucleation on steps or on the terraces between them rather than on heterogeneous sites such as emergent substrate dislocations, The initial supracritical clusters will tend to form hemispherical nuclei since this will minimise the surface area of the deposit. For crystalline deposits, however, the surface energy is anisotropic and the shape with minimum surface energy will be given by a modification of the Wulif construction [5]. This predicts that a flucleus in equilibrium with its surroundings will be bounded by facets with the lowest surface energy and that the area of a facet will be inversely proportional to the surface energy. The implication is that growth occurs more rapidly on onientations with high surface energy and that large facets of low energy orientations are left in their wake. Whilst such models would account for the presence of faceted islands, it seems implausible that the surface energy of the {111)B surfaces would be a factor of 5 lower than that of the { 1l1}~surfaces and thus other factors must be involved. One possibility is the mechanism proposed by Asai [6] for anisotropy in the lateral growth of homoepitaxial GaAs by MOVPE. By considering the bond configurations at surfaces and the adsorption/ desorption behaviour of As, it was shown that, at high AsH3 overpressures where the surface is covered by a monolayer of As, lateral growth was faster along [110] than along [110]. For a faceted island this would give faster growth on the Ga-terminated (111}A surfaces than on the As-terminated {111)B surfaces. It was further shown that, at lower AsH3 overpressures where As atoms may desorb, the opposite is true. One must, of course, be cautious in drawing parallels between these observations and the present system which is heteroepitaxial and involves a larger group V carrier molecule. It should also be noted that the present growth was performed in an atmosphere with no intentional excess of either TMSb or TMGa. If the thinner islands represent the optimum shape for an island formed by homogeneous nucleation under these conditions then it is not clear how the thicker islands arise. Their density is nearly three orders of magnitude higher than that of the dislocations in the substrate and they
173
are distributed evenly over the specimen, thus local fluctuations in the dislocation density cannot be responsible. Since they are usually located at the junction of two or more islands, it is probable that they are related to the misfit dislocations. As island nuclei develop, the strain energy which would be associated with a coherent deposit is released in a stepwise fashion by the introduction of 90° “Lomer” dislocations at the island edge [7]. The periodic repetition of this process leads to the misfit dislocation arrays observed in figs. 5 and 6. If some of the islands are slightly misaligned with one another, as indicated by the small variations in moire fringe spacing in fig. 4, then when two islands meet, some of these 90°dislocations may be forced out of the interface producing threading segments. The intersections of these segments with the deposit surface could act as preferential sites for secondary nucleation and growth. Whilst such configurations could give enhanced growth, leading to the bimodal distribution of island sizes observed experimentally, there is no direct evidence to support this hypothesis. The misfit dislocation arrangement revealed in weak beam images such as fig. 5 corresponds to that which would be expected for high misfit heteroepitaxy on an (001) substrate. The presence of unusual strain contrast in HREM images such as fig. 6 does, however, indicate a more complex structure. Two previous studies of the GaSb on (001) GaAs interface indicate what form this might take: (a) In a TEM study of early growth specimens deposited at 600°Cby MOVPE, Mallard et al. [8] observed a coherent crystalline layer 1.0—1.5 nm in thickness covering the substrate between the islands. They inferred that growth must occur in Stranski—Krastanov rather than Volmer—Weber mode despite the fact that the Dodson and Taybr model [9] indicates that there can be no coherent GaSb deposit of any thickness in stable equilibrium for this system. If such a coherent GaSb layer were present, then during the growth of islands, edge-type misfit dislocations would be introduced at the junction of the island with this layer rather than at the interface. In the absence of a significant driving force for climb processes,
174
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Geometry and interface structure of island nucleifor GaSb on (00]) GaAs
this would result in a misfit dislocation network at a distance from the interface equal to the thickness of the prior coherent layer. (b) In a more recent grazing incidence X-ray scattering study of continuous layers deposited at 470°Cby MBE [10], Bourret and Fuoss showed that their measurements were consistent with a corrugated interface with a period equal to the misfit dislocation spacing and an amplitude of 0.5—1.0 nm. It was suggested that this could arise by a complex diffusional mechanism or by the prior formation of a coherent layer containing “tiles” of GaSb and GaAs. A tiled coherent layer of this form would have both a lower interfacial energy and a lower elastic strain energy than a continuous coherent layer of GaSb alone, Both of the above structural models would certainly give rise to strain contrast at the interface in HREM images from cross sectional TEM specimens but detailed matching with computed image simulations would be required to establish which is a more appropriate model for the interface structure under the present deposition conuitiofls.
5. Summary Microscopical studies of the initial stages of buffer layer growth for MOVPE GaSb on (001) GaAs using AFM and TEM have enabled the following points to be established. (1) There is a high density of GaSb islands on the substrate surface; these have resulted from homogeneous nucleation. The geometry of these nuclei is consistent with the formation of crystallographic facets both to minimize surface energy and in response to anisotropic lateral growth rates. (2) There is a lower density of secondary islands present at junctions between primary nuclei.
These may have formed due to enhanced growth rates on threading segments of misfit dislocations. (3) The lattice misfit is accommodated by a very regular square arrangement of edge-type misfit dislocations. Unusual strain contrast in HREM images of the interface could be ascribed to the formation of a coherent layer prior to island growth leading to either a “stand-off’ of dislocations from the interface or a corrugated interface.
Acknowledgements The authors would like to thank Dr. I.P. Jones for helpful discussions and SERC for financial support. The AFM experiments were performed in the University of Birmingham STM/AFM Facility.
References [1] T. Nishimura, K. Kadoiwa, N. Hayafuji, K. Mitsui, H. Kumabe and T. Murotani, J. Crystal Growth 107 (1991) 468. [2] 5K. Haywood, N.J. Mason and P.J. Walker, J. Crystal Growth 93 (1988) 56. [3] E.T.R. Chidley, S.K. Haywood, RE. Mallard, N.J. Mason, R.J. Nicholas, R.J. Warburton and P.J. Walker, AppI. Phys. Letters 54 (1989) 1241. [4] J. Taftø and J.C.H. Spence, J. AppI. Cryst. 15 (1982) 60. [5]1. Markov and S. Stoyanov, Contemp. Phys. 28 (1987) 267. [6] H. Asai, J. Crystal Growth 80 (1987) 425. [7] R. Vincent, Phil. Mag. A 19 (1969) 1127. [8] R.E. Mallard P.R. Wilshaw, N.J. Mason, P.J. Walker and G.R. Booker, in: Microscopy of Semiconducting Materials 1989, Inst. Phys. Conf. 5cr. 100, Eds. AG. Cullis and J.L. Hutchison (Inst. Phys., London—Bristol, 1989) p. 331. [9] B.W. Dodson and PA. Taylor, AppI. Phys. Letters 49 (1986) 642. [10] A. Bourret and PH. Fuoss, AppI. Phys. Letters 61(1992) 1034.
~
CRYSTAL GROWT H
Geometry and interface structure of island nuclei for GaSb buffer layers grown on (001) GaAs by metalorganic vapour phase epitaxy M. Aindow
a T.T. Cheng a N.J. Mason b T.-Y. Seong b and P.J. Walker b School of Metallurgy and Materials, The University of Birmingham, Elms Road, Edgbaston, Birmingham B15 217’, UK b Physics Department, Clarendon Laboratory, The University of Oxford, Parks Road, Oxford OX] 3PU, UK a
Received 22 May 1993; manuscript received in final form 2 July 1993
Atomic force microscopy and transmission electron microscopy have been used to investigate the geometry and interface structure of island nuclei formed in the initial stages of buffer layer growth for MOVPE GaSb on (001) GaAs. There is a bimodal distribution of island sizes with a high density of small, homogeneous nuclei and a lower density of larger, secondary nuclei. The smaller islands have pronounced crystallographic facets which are consistent with those which would be expected for minimization of surface energy and lateral growth anisotropy. The secondary islands are present at junctions between primary nuclei and may have formed due to enhancement of growth rates at emergent threading segments of misfit dislocations. The lattice misfit is accommodated by a regular square arrangement of edge-type misfit dislocations but unusual strain contrast arises in HREM images due to either a “stand-off’ of dislocations from the interface or a corrugated interface.
1. Introduction Epitaxial films of GaSb on GaAs are of considerable interest for optoelectronic applications since the heterostructure junction is rectifying and diodes produced from such structures could be integrated with conventional GaAs device architectures. The deposition of GaSb films onto GaAs by metalorganic vapour phase epitaxy (MOVPE) is usually carried out by depositing a low temperature “buffer layer” of 20—50 nm of GaSb before ramping to the appropriate deposition temperature for the remainder of the film [1]. The role of the buffer is to promote nucleation and to ameliorate the effects of the 7.8% lattice mismatch by “turning over” the dislocations produced at the interface. In the present study we have concentrated on the initial stages of buffer layer formation, i.e. prior to the coalescence of island nuclei. The topography, size and distribution of these features have been investigated using atomic force microscopy (AFM) and the interface structure has been examined using transmission electron 0022-0248/93/$06.0O © 1993
—
microscopy (TEM). It is important to understand the nature of nucleation and growth for these buffer layers, since it may determine the types of crystallographic defects which are produced in response to the misfit. This could have important consequences for the nature and density of defects in any deposit grown onto the buffer. Moreover, the geometry of nucleation could have a greater significance; as with many semiconducting materials, GaSb shows a pronounced quantum confinement effect, thus if regular arrays of island nuclei of an appropriate size could be grown, these might be used directly as quantum dots.
2. Experimental details Gallium antimonide heteroepitaxial deposits were grown by atmospheric pressure MOVPE in a horizontal silica cell reactor with an RF-heated graphite susceptor, which is described in detail elsewhere [2]. The substrates used were Epi-ready GaAs: Si wafers with a nominal dislocation density of
Elsevier Science Publishers B.V. All rights reserved
M. Aindow eta!.
/
Geometry and interface structure of island nucleifor GaSb on (001) GaAs
tation of (001), i.e. no intentional offcut. The substrate preparation and pre-growth treatments were identical to those used previously [3]. The reactants used were trimethyl gallium (TMGa) (Epichem, 9°C)and trimethyl antimony (TMSb) (Alfa, 0°C),respectively with palladium diffused hydrogen as the carrier gas. Deposits were grown at 530°Cfor 5 mm: at this temperature, deposition rates of <1 ~sm/h are typical. Topographic measurements were obtained from these specimens with a Digital Instruments Nanoscope II contact-mode AFM using a conventional Si 3N4 with at a force constant 1 cantilever and scanning 4 Hz with of 0.06 N m The data files are measurements constant force. of specimen “height” as a function of probe position over a grid of 400 X 400 points within the scan area. Cross-sectional TEM specimens were produced by glueing two portions of the wafer together at 90° to one another and then cutting slices such that the surface normal, n, is parallel to the two orthogonal directions [110] and [110] which lie in the substrate surface. Discs cut from these slices were dimpled and ~ ion beam milled to perforation at liquid nitrogen temperatures. Plan-view specimens were produced in a similar manner and ion beam milled from the substrate side only. These specimens were exammed using a Philips CM2O TEM operating at 200 kV. —
169
±2%. There are two distinct types of island present; a large number of thin islands with an aspect ratio 5 (e.g. at A) and a smaller number of thicker islands with an aspect ratio 1 (e.g. at B). Islands of the first type are 80—110 nm thick, 180—300 nm wide and 1100—1300 nm in length. The density of these islands is difficult to estimate, but ignoring the possible effects of coalescence gives a lower bound of io~cm2. Islands of the second type are 140—250 nm thick, 350— 1000 nm wide and 900—1300 nm in length; the density of these islands is 8 x 106 cm2. We note thattwo these latterofislands always appear to straddle or more the thinner ones. Scans of more restricted areas give better resolution and reveal more clearly the effect of crystallography on the island topography. Fig. 2 is a perspective view of a surfaceplot with no exaggeration of vertical relief, constructed from a data file for a 3.25 X 3.25 jxm scan of the area marked X in fig. 1. In this case the “image” shows the shadows which would be expected if a surface with this geometry were illuminated from the left side. The thinner islands are bounded by a small upper facet which lies parallel to the substrate surface and by inclined facets on each side. Measurements from the data show that in each case, the inclined facets lie at an angle of 55°±3°to the substrate surface (we note that whilst
3. Results Topographic data were obtained from several areas on the sample using AFM and they revealed a homogeneous distribution of island nuclei with pronounced crystallographic facets. Fig. 1 is a topview “image” with shading corresponding to the height, constructed from a typical data file for a 16.5 x 16.5 ~m scan. The figure shows a large number of light rectangular protrusions on a dark background; comparing the scan axes to the {110} cleavage facets on the edges of the specimen, it is clear that these islands have edges parallel to [110] and [110]. The proportion of the surface which is covered by these islands is 83%
Fig. 1. Top view AFM image showing the distribution of GaSb islands — scan size 165 x 165 ~sm.
170
M. Aindow et a!.
/ Geometry and
interface structure of island nuclei for GaSb on (001) GaAs
Fig 2. Surface plot AFM image showing a detail of the region marked X in fig. 1 (top illuminated). Scan size 3.25 x 3.25 tim, vertical scale 0.7 urn — no exaggeration of vertical relief.
contact-mode AFM images can give considerable distortion to such angular measurements, the angles measured here correspond very closely to those measured from TEM images of cross-sectional specimens). Since the crystallographic axes of the substrate and deposit are parallel to one another, the upper facet is an (001) surface and the side facets are (111) surfaces. One example of
_________
the thicker islands is also present at Z in fig. 2. The topography is less faceted and exhibits some structure which is not well defined. We should emphasise that these are not isolated observations, data collected from various locations on the wafer revealed island nuclei with similar sizes, shapes and densities to those shown in figs. 1 and 2.
IL,
Fig. 3. Determination of polarity using CBED patterns obtained with B [270]: (a) pattern obtained with 002 at the Bragg condition; (b) pattern obtained with 002 at the Bragg condition.
M. Aindow eta!.
/ Geometry and
A
interface structure of island nuclei for GaSb on (001) GaAs
Ad
lOOnm Fig. 4. Many beam bright field TEM micrograph obtained from a plan view specimen with B
Fig. 5. Weak beam dark field TEM micrograph obtained from a plan view specimen with g
[001].
400 and B close to [001].
171
172
M. Aindow et al.
/
Geometry and interface structure of island nuclei for GaSh on (001) GaAs
Fig. 6. HREM image obtained with B = [110] showing a regular array of 90°type misfit dislocations (arrowed).
The polarity of the samples was established by using convergent beam electron diffraction (CBED) from cross sectional specimens in the TEM. The patterns were indexed with respect to a unit cell with basis Ga at 0, 0, 0 and As/Sb at 1/4, 1/4, 1/4. Following the method of Taftø and Spence [4], patterns were obtained from the substrate under dynamic two beam conditions with the beam direction B [270] as shown in fig. 3. The bright cross contrast inside the 002 reflection (fig. 3a) and the dark cross contrast inside the 002 reflection (fig. 3b) indicate that this portion of the specimen has n parallel to [110]. In equivalent patterns obtained from the portions of the specimen with n parallel to [110], the cross contrast is inverted. By examining images obtamed from these specimens, it is clear that the major and minor axes of the elongated islands are [110] and [110] respectively. Thus the larger indined facets are (111) and (111) which are the Sb-terminated (111)B surfaces and the smaller inclined facets are (111) and (111) which are the Ga-terminated (111)A surfaces. Diffraction contrast TEM images obtained from plan-view specimens confirm the distribution of island nuclei indicated in AFM images such as fig. 1. Fig. 4 is a typical bright field image obtained with the specimen oriented with B = [001]. The positions of the islands are revealed by the presence of two sets of perpendicular moire fringes with a spacing of 2.9 nm. This spacing corresponds to that which would be expected for interference between the 220-type reflections in GaSb and GaAs and indicates that the deposit is
fully relaxed. In weak-beam dark field images the amplitude of moire fringe contrast is attenuated strongly and dislocation contrast dominates the images. A representative micrograph is shown in fig. 5 which was obtained with the specimen onented for g/3g with g = 400, where g is the reciprocal lattice vector of the diffracting planes. A regular orthogonal arrangement of dislocations is observed with line directions [1101 and [110], each array having a spacing of 5.8 nm. This is consistent with the arrangement that would be expected for full accommodation of misfit by arrays of edge type misfit dislocations with Burgers vectors b ~[110] and ~[110], respectively. This has been confirmed by obtaining high resolution electron microscopy (HREM) images of the crystal lattices in the interface region from cross sectional specimens. One such image, obtained with B = [110] is presented in fig. 6 with the edges of terminating half planes indicated by arrows. It is interesting to note that such images show unusual strain contrast at the interface with “cusps” at the position of the dislocations; this contrast was present for all values of defocus which gave lattice resolution. =
4. Discussion The AFM data presented here show that the number density of the thinner island deposits is five orders of magnitude higher than the stated dislocation density for the substrates. These islands must, therefore, arise from homogeneous
M. Aindow eta!.
/
Geometry and interface structure of island nucleifor GaSb on (001) GaAs
nucleation on steps or on the terraces between them rather than on heterogeneous sites such as emergent substrate dislocations, The initial supracritical clusters will tend to form hemispherical nuclei since this will minimise the surface area of the deposit. For crystalline deposits, however, the surface energy is anisotropic and the shape with minimum surface energy will be given by a modification of the Wulif construction [5]. This predicts that a flucleus in equilibrium with its surroundings will be bounded by facets with the lowest surface energy and that the area of a facet will be inversely proportional to the surface energy. The implication is that growth occurs more rapidly on onientations with high surface energy and that large facets of low energy orientations are left in their wake. Whilst such models would account for the presence of faceted islands, it seems implausible that the surface energy of the {111)B surfaces would be a factor of 5 lower than that of the { 1l1}~surfaces and thus other factors must be involved. One possibility is the mechanism proposed by Asai [6] for anisotropy in the lateral growth of homoepitaxial GaAs by MOVPE. By considering the bond configurations at surfaces and the adsorption/ desorption behaviour of As, it was shown that, at high AsH3 overpressures where the surface is covered by a monolayer of As, lateral growth was faster along [110] than along [110]. For a faceted island this would give faster growth on the Ga-terminated (111}A surfaces than on the As-terminated {111)B surfaces. It was further shown that, at lower AsH3 overpressures where As atoms may desorb, the opposite is true. One must, of course, be cautious in drawing parallels between these observations and the present system which is heteroepitaxial and involves a larger group V carrier molecule. It should also be noted that the present growth was performed in an atmosphere with no intentional excess of either TMSb or TMGa. If the thinner islands represent the optimum shape for an island formed by homogeneous nucleation under these conditions then it is not clear how the thicker islands arise. Their density is nearly three orders of magnitude higher than that of the dislocations in the substrate and they
173
are distributed evenly over the specimen, thus local fluctuations in the dislocation density cannot be responsible. Since they are usually located at the junction of two or more islands, it is probable that they are related to the misfit dislocations. As island nuclei develop, the strain energy which would be associated with a coherent deposit is released in a stepwise fashion by the introduction of 90° “Lomer” dislocations at the island edge [7]. The periodic repetition of this process leads to the misfit dislocation arrays observed in figs. 5 and 6. If some of the islands are slightly misaligned with one another, as indicated by the small variations in moire fringe spacing in fig. 4, then when two islands meet, some of these 90°dislocations may be forced out of the interface producing threading segments. The intersections of these segments with the deposit surface could act as preferential sites for secondary nucleation and growth. Whilst such configurations could give enhanced growth, leading to the bimodal distribution of island sizes observed experimentally, there is no direct evidence to support this hypothesis. The misfit dislocation arrangement revealed in weak beam images such as fig. 5 corresponds to that which would be expected for high misfit heteroepitaxy on an (001) substrate. The presence of unusual strain contrast in HREM images such as fig. 6 does, however, indicate a more complex structure. Two previous studies of the GaSb on (001) GaAs interface indicate what form this might take: (a) In a TEM study of early growth specimens deposited at 600°Cby MOVPE, Mallard et al. [8] observed a coherent crystalline layer 1.0—1.5 nm in thickness covering the substrate between the islands. They inferred that growth must occur in Stranski—Krastanov rather than Volmer—Weber mode despite the fact that the Dodson and Taybr model [9] indicates that there can be no coherent GaSb deposit of any thickness in stable equilibrium for this system. If such a coherent GaSb layer were present, then during the growth of islands, edge-type misfit dislocations would be introduced at the junction of the island with this layer rather than at the interface. In the absence of a significant driving force for climb processes,
174
M. Aindow eta!.
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Geometry and interface structure of island nucleifor GaSb on (00]) GaAs
this would result in a misfit dislocation network at a distance from the interface equal to the thickness of the prior coherent layer. (b) In a more recent grazing incidence X-ray scattering study of continuous layers deposited at 470°Cby MBE [10], Bourret and Fuoss showed that their measurements were consistent with a corrugated interface with a period equal to the misfit dislocation spacing and an amplitude of 0.5—1.0 nm. It was suggested that this could arise by a complex diffusional mechanism or by the prior formation of a coherent layer containing “tiles” of GaSb and GaAs. A tiled coherent layer of this form would have both a lower interfacial energy and a lower elastic strain energy than a continuous coherent layer of GaSb alone, Both of the above structural models would certainly give rise to strain contrast at the interface in HREM images from cross sectional TEM specimens but detailed matching with computed image simulations would be required to establish which is a more appropriate model for the interface structure under the present deposition conuitiofls.
5. Summary Microscopical studies of the initial stages of buffer layer growth for MOVPE GaSb on (001) GaAs using AFM and TEM have enabled the following points to be established. (1) There is a high density of GaSb islands on the substrate surface; these have resulted from homogeneous nucleation. The geometry of these nuclei is consistent with the formation of crystallographic facets both to minimize surface energy and in response to anisotropic lateral growth rates. (2) There is a lower density of secondary islands present at junctions between primary nuclei.
These may have formed due to enhanced growth rates on threading segments of misfit dislocations. (3) The lattice misfit is accommodated by a very regular square arrangement of edge-type misfit dislocations. Unusual strain contrast in HREM images of the interface could be ascribed to the formation of a coherent layer prior to island growth leading to either a “stand-off’ of dislocations from the interface or a corrugated interface.
Acknowledgements The authors would like to thank Dr. I.P. Jones for helpful discussions and SERC for financial support. The AFM experiments were performed in the University of Birmingham STM/AFM Facility.
References [1] T. Nishimura, K. Kadoiwa, N. Hayafuji, K. Mitsui, H. Kumabe and T. Murotani, J. Crystal Growth 107 (1991) 468. [2] 5K. Haywood, N.J. Mason and P.J. Walker, J. Crystal Growth 93 (1988) 56. [3] E.T.R. Chidley, S.K. Haywood, RE. Mallard, N.J. Mason, R.J. Nicholas, R.J. Warburton and P.J. Walker, AppI. Phys. Letters 54 (1989) 1241. [4] J. Taftø and J.C.H. Spence, J. AppI. Cryst. 15 (1982) 60. [5]1. Markov and S. Stoyanov, Contemp. Phys. 28 (1987) 267. [6] H. Asai, J. Crystal Growth 80 (1987) 425. [7] R. Vincent, Phil. Mag. A 19 (1969) 1127. [8] R.E. Mallard P.R. Wilshaw, N.J. Mason, P.J. Walker and G.R. Booker, in: Microscopy of Semiconducting Materials 1989, Inst. Phys. Conf. 5cr. 100, Eds. AG. Cullis and J.L. Hutchison (Inst. Phys., London—Bristol, 1989) p. 331. [9] B.W. Dodson and PA. Taylor, AppI. Phys. Letters 49 (1986) 642. [10] A. Bourret and PH. Fuoss, AppI. Phys. Letters 61(1992) 1034.