A1As heterostructures

Journal of Crystal Growth 102 (1990) 891—898 North-Holland

891

ATOMIC LAYER MOLECULAR BEAM EPITAXY OF InAs/AIAs HETEROSTRUCTURES M. VAZQUEZ, J.P. SILVEIRA, L. GONZALEZ, M. PEREZ, G. ARMELLES, J.L. DE MIGUEL and F. BRIONES Centro Nacional de Microelectronica,

CSIC, Serrano 144,

28006 Madri4 Spain

Received 25 September 1989; manuscript received in final form 20 January 1990

Experimental results on 0.3 ~sm thick (AlAs),~(InAs),,superlattices (m, n = 1, 2,3, 5, 10 monolayers) grown by atomic layer molecular beam epitaxy on (100) GaAs substrates are presented, X-ray and Raman experiments show that they are decoupled from the substrate in the range of composition studied. These results are compared with lattice relaxation data of thick layers of InAs (AlAs) grown on AlAs (InAs) by the same technique.

1. Introduction The possibility of tailoring the electronic and optical properties of semiconductors by growing superlattices has led recently to a research on strained materials combinations. Novel properties of superlattices arise when one or both of the components are under significant strain. Of particular technological interest are those combinations with a band gap in the range 1.3—1.5 ~sm. This range can be achieved by stacking, for exampie, layers of InAs and GaAs or InAs and AlAs, The main problem in obtaining such highly strained superlattices (— 7% lattice mismatch) is how to grow high quality layers with sharp interfaces. It has been demonstrated that strain plays an important role in surface kinetics and consequently in the growth mode [1,2], as well as that the growth mode dominates the kinetics and type of defects appearing in the process of lattice relaxation [3]. Molecular beam epitaxy (MBE), a successful growth technique for coherent epitaxial systems, is not so effective in growing largely mismatched materials, when formation of 3D nuclei at the interface is difficult to avoid. A modification of the conventional MBE technique, atomic layer molecular beam epitaxy (ALMBE) [4,5], while keeping the inherent advantages of MBE (such as high growth rate and monolayer 0022-0248/90/$03.50 © 1990

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control of interfaces) seems to be more appropriate for the growth of such strained layer structures. ALMBE has already been used to grow high quality InAs [6], (GaAs)/(GaP) [7} and GaAs quantum wells confined by GaAs/GaP superlattices on (100) GaAs substrates. The major modification of ALMBE over conventional MBE originates in a cyclic perturbation of the growth front in synchronism with the monolayer by monolayer growth sequence. Practically, this perturbation is induced by interrupting or pulsing the group V beam, or by alternating group V and group III beams. While in conventional MBE, 2D growth (in competition with other mechanisms such as step propagation and 3D island formation) is only reached over an optimized range of temperatures, V/Ill flux ratios and low density of surface steps, the ALMBE process has actually proven to enhance 2D growth in a broad range of growth conditions. In particular, high quality layers of materials that demand different growth conditions by MBE, like InAs and AlAs, can be obtained at a common low substrate temperature (7 350—400°C) by ALMBE. This fact opens the possibility of building up high quality superlattices from such different materials. For this work, two series of (AIA5)m(InAs)n superlattices were grown by ALMBE on (100) GaAs substrates. In each of the series the thick=

Elsevier Science Publishers B.V. (North-Holland)

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/ ALMBE of InAs/A1As

heterostructures

ness of one of the constituent materials (InAs or AlAs, respectively) was kept constant to 15 mono-

recording system. The evolution of the in-plane lattice parameter, a11, deduced from the separation

layers (ml), while the thickness of the other material was varied up to 1, 2, 3, 5, and 10 ml. In order to get some insight into the process of lattice relaxation when growing InAs on AlAs and vice versa, thick layers of InAs (AlAs) were grown on sufficiently thick layers of AlAs (InAs), to remove the influence of the substrate under the same growth conditions used for the superlattices. These conditions, necessarily a compromise, have previously resulted in good quality layers of InAs [6] and AlAs [81grown directly on GaAs. The evolution of the electron beam diffraction (RHEED) patterns, corresponding to the growth of these thick layers, has been carefully analyzed, and the results compared with the data obtained from X-ray diffraction 0—20 scans and Raman spectra for the AIAs/InAs SLs. Our RHEED results show that InAs grows on AlAs in a 2D fashion, reaching its lattice parameter in few monolayers. In contrast, a slow lattice relaxation process and a rough growth front is observed for AlAs grown on InAs. The X-ray and Raman results for the (AIA5)m(InAs)n superlattices show the enormous impact that only a few monolayers of InAs have on the growth mechanism of the strained AlAs layers.

between rods, can be followed during the growth over consecutive video images with a minimum time interval of 40 ms. The superlattices were characterized by X-ray diffractometry and Raman spectroscopy. A conventional powder X-ray diffractometer (Siemens D500 Kristalloflex) with a Cu source was used in a 6—20 geometry symmetric about the growth direction. Raman spectra of the SLs were obtained at RT in the backscattering configuration. The scattered light was analyzed by means of a computer controlled 0.85 m double monochromator.

2. Experimental The samples were grown by ALMBE on undoped (100) GaAs substrates, as described elsewhere [4]. The As4 beam equivalent pressures used were in the range (3—5) X 6 Torr. After oxide removal, the As4 cell was shuttered and the substrate temperature was lowered to 400°C. The indium and aluminum fluxes were previously adjusted by means of RHEED specular beam oscillations to yield a growth rate of one monolayer per second. The ALMBE procedure used in this work consisted of the continuous supply of group III element (In or Al) and periodic short pulses (0.3 s) of As4, in coincidence with the deposition completion of one atomic layer of In or Al. The RHEED pattern is followed during growth with a high sensitivity CCD camera and a video

3. InAs/AIAs thick layers Due to the large difference in lattice parameters between InAs and AlAs (7.4%), a rapid misfit relaxation process is expected during the growth of one material onto the other. Fig. la shows the lattice mismatch during the first stages of the growth of InAs on a 0.3 ~tm thick AlAs layer grown on a (100) GaAs substrate, as measured from the evolution of the distance between RHEED rods, d (inversely proportional to the change of the in-plane lattice parameter a11). The fast variation of d observed indicates that that lattice parameter of the InAs epilayer breaks away from that of the AlAs buffer layer as soon as 3 ml of InAs are deposited. The InAs bulk lattice parameter is reached in 6 ml indicating, within the experimental error (— 0.3%) of these measurements, a complete lattice relaxation. This behaviour is the same as that observed for the ALMBE growth of InAs on GaAs [6], a system with essentially the same amount of lattice mismatch. Fig. lb shows the evolution of the full width at half maximum (FWHM) of the (00) RHEED streak during the same growth run of fig. la. We observe that the (00) rod broadens suddenly after the deposition of very few monolayers of InAs reaching a maximum width for 6 ml. Thereafter RHEED streak becomes progressively narrower, showing a FWHM comparable to the initial of the AlAs (00) rod after the growth of 10 ml. Notice that the starting AlAs surface, grown without mis-

M. Vazquez et a!. / ALMBE of InAs/A1As heterostructures

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match on GaAs shows a sharp RHEED pattern corresponding to the very flat surface morphology, achieve by ALMBE. The broad and diffuse spots observed during growth of InAs could be originated by a loss of diffraction coherence: note that a rapid misfit relaxation process, as that shown on fig. 1, involves the formation of a high density of defects and consequently a loss of coherence for diffraction. As for the system InAs—GaAs [6], even during growth of the first monolayers of InAs on AlAs an alternation between (2 X 4) and (4 X 2) InAs surface reconstructions is clearly seen. However, fractional order streaks show a small and practically constant halfwidth. This indicates that in spite of the large mismatch InAs grows on AlAs in a 2D fashion by ALMBE. Finally, fig. lb shows that when InAs thickness reaches 6 ml, in coincidence with the fully relaxed lattice situation, halfwidth for the integral order streaks rapidly improves to achieve instrumental values in about 10 ml more.

width increases progressively from the first monolayers (see fig. 2b), and does not show any clear recovery as growth proceeds. This lack of recovery contrasts greatly with the fast recovery observed when growing InAs on AlAs (see fig. la), and seems to indicate that the lattice relaxation is no longer confined to two or three monolayers near the interface, but it extends to about 30 ml. The broad RHEED streaks, holding up to a thicknesses where the lattice is completely relaxed, can be considered an evidence of AlAs surface roughening. It is obvious that low substrate temperature can be more detrimental for the growth of a lower surface mobility and higher chemical reactivity material such as AlAs. However, previ-

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ous experiments [6,8] have shown that monolayer flat 2D growth of AlAs on GaAs (100) can be obtained by ALMBE under the same experimental conditions as those used in the present experiments. The differences in the growth behaviour of InAs and AlAs could also be due to the asymmetry of strain for both materials: AlAs grown on InAs is subjected to a tensile strain, while InAs grown on AlAs suffers a compressive strain. This effect has already been observed by Lievin and Fonstad [2] when growing In5Al1 ~As on InP by conventional MBE. These authors obtain better 2D growth for those compositions where the lattice mismatch is positive: the higher the In composition, the better the 2D growth. This was attributed to the influence of strain on the surface bond strength of cations, so that the migration rate may be quite different, affecting to the growth mechanism. More experimental work is needed in order to clarify the role of the sign of mismatch on the growth kinetics. On the other hand, it is also important to keep in mind that the details of the relaxation process of a mismatched epilayer involves a balance between the built in deformation energy in the layer, and the energy necessary for the formation of a certain type of defects (dislocations). If, as expected from the degree of bond covalence [9], the formation of dislocations in AlAs implies overcoming a higher threshold energy than in InAs, the differences in the relaxation process of both epilayers could be explained.

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lated to the AlAs LO phonon (broader than for the (AIAs))5(InAs)2 SL), shifted further in energy, and a peak at lower energy related to the TO phonon. Notice that this phonon should be forbidden in our geometry, and its appearance evidences a loss of crystalline quality of the AlAs layers caused by the introduction of only two monolayers of InAs. The Raman spectra for the superlattices with more than two monolayers of InAs in each sublayer consists of a single broad peak, as shown for the (AIAs)15(InAs)5 SL in fig. 3c. The obvious loss of crystallimty of sample c causes a relaxation of the k 0 selection rules and allows different kinds of phonons to participate in the observed broad scattering peak. In this case the shift of the LO(F) phonons caused by strain cannot be determined. Similar experiments have been performed for the symmetric case, InAs rich superlattices. Figs. 4a—4d show Raman spectra for the (AIAS)m (InAs)15 SLs, m 1, 2, 5, and 10 ml, respectively, As for the AlAs rich superlattices (figs. 3a—3c), the larger the amount of AlAs in every sublayer of the superlattice, the larger the shift of the InAs LO phonon related peak, indicating an increase of strain in the InAs layers. Obviously all of the SLs of this set are completely decoupled as a whole from the GaAs substrate. From n 5 on, the InAs phonon related peak does not shift significantly any further, as can be observed in fig. 4 for the (A1As)10(InAs)15 SL, and simultaneously a drastic increase in the TO phonon related peak is observed. So above critical thickness for AlAs grown on InAs (only 2 ml as shown by the RHEED results), misfit dislocations must be formed within each superlattice period. Theoretical estimations of critical thicknesses for a lattice mismatch of 7% would supply values in this range [10,11]. A direct comparison of spectra in figs. 3 and 4 is no possible because of the quite different absorption depth for the laser light in the two sets of samples. In particular in the case of (AlAs)m(InAs)is SLs we obtain information of a depth of only about 300 A near the surface while to the (AlAs)15(InAs)~ SLs the whole sample is analyzed. In order to get a deeper insight into the relaxation process, the (AIAS)m(InAs)n SLs have been further characterized by X-ray diffraction. Fig. 5 =

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and 6 show X-ray diffraction scans in the vicinity of the (002) GaAs diffraction peak for the two series of (A1As)15(InAs)~ and (AIA5)m(InAs)55 SLs, respectively. The values of the SL period calculated from the satellite peaks position are very close to those expected from layer design parameters. The intensity distribution of the superlattice satellite peaks is not symmetric with respect to the zero order reflection of the superlattice, as expected for the large difference of lattice parameters mismatch between both materials [12]. The three first spectra of both fig. 5 and fig. 6 show that the number of satellites increase with the thickness of the minority component of the superlattice, as would be expected from kinematical theory calculations [13]. However, in order to obtain information about the crystalline quality of these samples, we have to consider the FWHM of the sateffites. Table 1 shows the measured FWHM values obtained for the zero order diffraction peak of both (AlAs)15 (InAs)~and (AlAs)m(InAs)15 SLs. These data evidence a loss of crystalline quality when increasing the number of monolayers of InAs(A1As) in the A1As(InAs) rich superlattices, in agreement with above presented Raman results. However a detailed analysis is rather complicated, taking into account that the mosaic spread is not reflected in the FWHM values obtained in 0—20 scans [14]. More experimental work with a double crystal diffractometer is underway.

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5. Conclusions Thick layers of InAs (AlAs) on AlAs (InAs), and (AIAS)m(IflAS)n superlattices have been grown by ALMBE on(l00) GaAs substrates. According to RHEED measurements InAs grown by ALMBE on AlAs exhibits a layer by layer growth mode, reaching its bulk lattice parameter in few monolayers after rapidly decoupling from the AlAs thick layer. In contrast, the growth of AlAs on InAs is characterized by a slower relaxation process. This behaviour can be attributed to the asymmetry of strain (tensile versus compressive) for both heterostructures and its influence on the growth kinetics and defects formation. Differences in the threshold energy necessary for the formation of relaxing defects (dislocations) need also to be considered. Concerning the possibility of fabricating InAs/AIAs SLs, our results using X-ray and Raman spectroscopy show clearly that they are pseudomorphic only if one of individual layer thicknesses is of the order of 1 or 2 ml, in agreement with our previously described RHEED results and with semiempirical estimations of critical thicknesses.

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[10] I.J. Fritz, App!. Phys. Letters 51(1987) 1080. [11] J.H. van der Merwe and W.A. Jesser, J. App!. Phys. 63 (1988) 1509. [12] M. Quillec, L. Goldstein, G. Le Roux, J. Burgeat and J. Primot, J. App!. Phys. 55 (1984) 2904. [13] J. Kervarec, M. Baudet, J. Caulet, P. Awray, J.Y. Emery and A. Regreny, J. App!. Cryst. 17 (1984) 196. [14] K. Kanngaki, H. Sakashita, H. Kato, N. Nakayawa, N. Sano and H. Terauchi, App!. Phys. Letters 49(1986)1071.