- Email: [email protected]
Investigation of the structural properties of MBE grown ZnSeGaAs heterostructures
ELSEVIER
Journal of Crystal Growth 175/176 (1997) 571-576
Investigation of the structural properties of MBE grown ZnSe/GaAs heterostructures I. Hernhdez-CaldercW*, E. Lbpez-Luna”, J. Luyo”, M. Mekndez-Lira”, 0. de Meloa,‘, P. Diaz”,‘, L. Hernhndeza”, J. Fuente?l, R. Lebn”,‘, H. Sitterb aPhysics Department, CINVESTAV, Apdo. Postal 14-740, 07000 Mbxico. D.F., Mexico b Institut ftir Experimentalphysik, Universitiit Linz, Linz, Austria
Abstract The results of photoluminescence and Raman spectroscopies, high resolution X-ray diffraction, and Auger electron spectroscopy are analyzed in terms of the structural properties of the ZnSe/GaAs(l 0 0) system as a function of film thickness and substrate temperature. The results of Raman spectroscopy and X-ray diffraction clearly show that the strain in the film is inhomogeneous and depends only on film thickness and not on growth temperature in the 285-325°C range. From these experiments a value of N 0.17 pm is inferred for the critical thickness of ZnSe on GaAs. Photoluminescence experiments sensitive to the ZnSe/GaAs interface reveal the presence of strain in the GaAs substrate. Analysis of the intensities of the LMM Auger transitions of Zn and Se indicate the formation of an interfacial layer with excess of Se, suggesting the formation of a pseudomorphic Ga-Se compound mixed with ZnSe at the interfacial region. PACS: 61.10. -i;
68.35.Ct; 68.55.Vk; 78.30.F~; 78.66. - w; 78.66.Hf
ZnSe; ZnSe/GaAs heterostructures; luminescence; Critical thickness
Keywords:
Semiconductor
1. Introduction
The study of ZnSe/GaAs interfaces has attracted the attention of researchers since the last decade [l]. However, the recent increase in the investigation of the physical properties of ZnSe/GaAs het-
* Corresponding author. E-mail: [email protected]. ’Permanent address: Faculty of Physics, University of Havana, Cuba. 0022-0248/97/$17.00 Copyright PII SOO22-0248(96)01215-S
0
interfaces;
Auger;
X-ray
diffraction;
Raman;
Photo-
erostructures is noteworthy. One of the main reasons is that the improvement of their structural and chemical properties is of fundamental relevance for the elaboration of green-blue emission lasers based on ZnSe [2]. Recently, Sony reported a green laser with a lifetime larger than 100 h [3]. The noticeable improvement was attributed to the reduction of defect density from lo6 to around lo4 cme2. It needs to be further reduced to inhibit defect propagation during laser operation [4]. Wu and coworkers very recently reported that low defect
1997 Elsevier Science B.V. All rights reserved
512
I. Hernirndez-Calderbn
et al. /Journal
density (less than lo4 cm- 2, pseudomorphic ZnSSe and ZnSe layers can be grown on an As-rich GaAs buffer layer initially exposed to Zn [S]. Some of the intrinsic defects at the interface are caused by the chemical and structural differences between the GaAs substrate and the ZnSe film. The lattice mismatch produces biaxially compressed films which release the strain after reaching the critical thickness through the formation of misfit dislocations. It has also been shown that the charge imbalance at the ZnSe/GaAs interface can lead to 3D nucleation and interface roughening [6]. Interdiffusion and crystalline degradation can occur as a result of variations in the ZnSe stoichiometry near the interface [7]. The quality of the epilayers is strongly affected by initial growth conditions such as substrate preparation, growth temperature, and Zn to Se flux ratios [S]. Here, we report results on the investigation of the structural properties of ZnSe/ GaAs(1 0 0) heterostructures grown by molecular beam epitaxy (MBE) by means of photoluminescence (PL) and Raman (RS) spectroscopies, high resolution X-ray diffraction (HRXRD) experiments, (AES). and Auger electron spectroscopy
2. Experimental
details
The films were grown in a RIBER 32P MBE system. A series of epitaxial ZnSe films were deposited on semi-insulating GaAs(1 0 0) substrates at temperatures between 28.5 and 325°C. Prior to their introduction to the MBE chamber the substrates were chemically treated as described in Ref. [l 11. The native oxide layer was removed by heating the indium glued substrate at about 550°C in ultrahigh vacuum (in the low lo-” Torr range). The reflection high energy electron diffraction (RHEED) experiment of the annealed surface showed a (2 x 1) reconstruction pattern, indicating that residual Se from the background atmosphere reacted with the GaAs(1 0 0) surface [S, 9, lo]. The typical Zn/ Se beam pressure ratio was -0.3. The RHEED patterns confirmed the Se rich growth through the ZnSe(2 x 1) surface reconstruction. A typical growth rate of - 1 pm/h was employed [ 111. The PL experiments were performed in a standard setup equipped with He-Ne and Kr-Ar lasers,
of Crystal Growth 175/l 76 (1997) 571~ 576
a 0.5 m monochromator and a closed cycle He refrigerator. The Raman experiments were done with a double monochromator equipped with CCD detection and the 4880 A line of an Ar laser. To verify the selection rules of RS we used a perfect backscattering x(z, z)X configuration; x, y, and z correspond to the (1 0 0) crystal directions. The HRXRD experiments were done using CuK,, radiation and a four-crystal Bartels monochromator [12]. The AES spectra were taken in an analysis chamber connected through UHV to the growth chamber. The Auger spectra were measured immediately after growth of the films employing a MAC3 analyzer set to 2 eV resolution; the energy of the primary electrons was 3 keV.
3. Results and discussion At room temperature (RT) the 0.26% lattice mismatch between GaAs (a = 5.6532 & RT) and ZnSe (a = 5.6676 A, RT) gives place to biaxial stress at the interface. Most of the elastic deformation is accumulated in the ZnSe film, and it is customary to neglect the substrate deformation. A precise analysis of the strain must take into account the differences between the thermal expansion coefficients of ZnSe and GaAs, since in most cases the temperature of the sample during characterization or device operation is different from the growth temperature. A systematic analysis of the deformation, both in the substrate and in the films, is needed to have a better understanding of the interface strain. With the purpose of directly identifying the presence of substrate strain, we performed PL measurements sensitive to the GaAs surface. Fig. la shows the PL of covered and clean GaAs substrates. Since ZnSe is transparent to the 4880 A photons (2.54 eV), we can easily reach the GaAs/ ZnSe interfacial region. We observe a broad transition around 1.52 eV associated to excitonic transitions [13] and carbon impurities [14]. Transitions around 1.45 eV have been attributed to GaAs antisites [15] and those around 1.41 eV to Mn impurities [16]. There are additional transitions at lower energies, but here we are only interested in their energy shifts with film thickness. It is worth mentioning that the PL spectra taken
et al. I Journal of‘Crystal Growth 175/I 76 (1997) 571-576
I. Hermindez-Calderbn
, ZnSelGaAs(100) Photolum~nescence br 4880A. 17 K
ating that the uppermost interfacial region detected with the 4880 A line has structural imperfections. At first glance the spectra shown in Fig. la could indicate that the substrate is not affected by the epilayer. However, a careful analysis of the PL peaks for samples with different thickness shows very small but noticeable and consistent energy shifts (see Fig. lb). The general trend is an increase in energy with ZnSe film thickness up to around 0.2 urn and then a slight decrease. Considering the reduced magnitude of these changes, we realize this description may appear a bit imprecise, but as we will see later it can be understood in terms of the evolution of the strain in the substrate and the film before and after the critical thickness is reached and of the difference in the thermal expansion coefficients of ZnSe and GaAs. As expected from high quality films, the RS experiments in the backscattering configuration, x(z, z)X, did not show the TO phonon of the ZnSe film but a very weak TO signal from GaAs, indicating some degree of disorder at the interface. Fig. 2a shows the spectra of the films grown at 325°C; the
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Fig. 1. (a) PL spectra of the clean and ZnSe covered substrates. (b) Energy shifts of the PL peaks of GaAs as a function of thickness of the ZnSe epilayer.
with the 6328 A line of a HeeNe ples a larger GaAs volume than exhibited sharper, stronger and solved excitonic peaks than those
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Fig. 2. (a) Raman spectra of ZnSe/GaAs heterostructures in a perfect backscattering configuration.(b) phonon ofZnSe as a function offilm thickness.(c) Changes in the FWHM of the LO phonon ofZnSe. the eye.
2.5 3.0 (pm)
Changes in the energy of the LO The straight lines arejust guides to
514
I. Hernbndez-Calderbn et al. I Journal of Crystal Growth 175/l 76 (I 997) 571-576
variations in intensity between the spectra are due to interference effects caused by different film thicknesses. Fig. 2b shows the observed shift in the energy of the ZnSe LO phonon and Fig. 2c the change in its full width at half maximum (FWHM) as a function of the ZnSe thickness. We observe that the LO peak moves towards lower energies up to around 0.17 urn and then remains at a value of 250.8 cm- r. The total shift is around 1 cm-‘. A similar shift was observed between thin and thick ZnSe films grown by chemical vapor deposition [17]. In Fig. 2c we observe that the FWHM of the LO phonon increases from around 4 cm-i for the thinnest film to around 8 cm- 1 for the 0.25 ym thick ZnSe film; afterwards it remains at around 7cmP’. Looking at the straight lines (which are drawn just as a guide to the eye), we can describe the general behavior of the FWHM as an increase to around 0.18 urn and constant afterwards. Then, from the evolution of the LO energy shift and its FWHM, we can conclude that the changes depend on the thickness but not on the substrate temperature. These changes also indicate that the critical thickness (h,) of ZnSe in GaAs(1 0 0) is around 0.17 urn. This value is consistent with previously reported values [lS, 191. The fact that the FWHM shows an important increase around h, is an indication of the presence of structural disorder produced by the appearance of misfit dislocations which release the strain caused by the lattice mismatch. The observation in Fig. 2 that the average compressive biaxial strain of the ZnSe film is continuously reduced with increasing film thickness up to h,, is a direct indication of nonuniform (or inhomogeneous) stress in the films, in agreement with previous results of Olego et al. [20]. In Fig. 3 we have summarized the results of HRXRD experiments. Fig. 3a presents the difference between the angle of the (0 0 4) reflection of ZnSe with respect to that of GaAs as a function of film thickness. We can see a very close similarity in the behavior of d@ and that of the LO phonon energy shown in Fig. 2a. Again, no dependence in the substrate temperature is observed. However, for HRXRD the changes reveal directly the modification of the (1 0 0) interplanar distance parallel to the substrate, indicating a value of h, around 0.17 urn, consistent with the Raman measurements.
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Fig. 3. (a) Difference between the angle of the (0 0 4) reflection of ZnSe and GaAs as a function of film thickness. (b) Changes in the FWHM of the (0 0 4) reflection of ZnSe as a function of film thickness. The lines are just guides to the eye.
The presence of residual strain is observed even for the thickest film. Fig. 3b shows the modification of the FWHM of the (0 0 4) reflection of the HRXRD patterns of the heterostructures. Analogous to the previous findings, we can see two types of behavior: a fast decrease up to around h, and then a very slow decrease with thickness. Recent similar experiments by other groups show a variety of behaviors of the FWHM of the (0 0 4) reflection with ZnSe film thickness. Reichow et al. [21] found a continuous decrease of the FWHM with thickness of MBE deposited films, very close to our results. The report of Sou et al. [19] does not show a clear trend of the FWHM with film thickness in MBE deposited films. Lee et al. [17] found a reduction of the FWHM with film thickness of chemical vapor deposited films. Then, apparently the modification of the FWHM of the (0 0 4) reflection depends on the type and conditions of growth. In our case, we attribute the continuous decrease of the FWHM to
I. Hernirndez-Calderh
515
et al. 1 Journal ofCyysta1 Growth 175/l 76 (1997) 571-576
the improvement of the overall film structural quality with thickness. We start with a relatively rough substrate and the film becomes smooth with increasing ZnSe thickness. The apparent discrepancy between the conclusions from Fig. 2c and Fig. 3b is easily understood when we consider that Raman spectroscopy is very sensitive to the microscopic structure, while HRXRD is more sensitive to the average macroscopic structure of the films. Based on the previous results it is easier to explain the shifts observed in PL in Fig. la. First, we have to consider the fact that the measurements were done at low temperature. Since the thermal expansion coefficient of ZnSe is larger than that of GaAs [22], after cooling to 17 K the situation will be reversed in relation to the one at growth temperature, the ZnSe film will be under biaxial tensile stress and the GaAs interfacial region will suffer biaxial compressive strain. The result is an increase in the band gap of GaAs with a maximum value just around h,. The transitions observed in the PL spectra of Fig. 1 indicate this behavior. After h, is reached most of the strain is released through the misfit dislocations and then the band gap energy of GaAs tends to decrease. Then, from the PL measurements we have obtained a direct indication of substrate strain. We performed Auger experiments of the ZnSe/ GaAs heterostructure as a function of ZnSe deposition of very thin films. Fig. 4 illustrates the ratio of the peak to peak intensities of the LMM transitions of Zn and Se from Auger derivative spectra. The intensities were measured at 5, 15, 30, 60 and 300 s deposition time. The typicalOgrowth rate of the films was - 1 urn/h (- 2.8 A/s); however, since growth conditions can change with film thickness we prefer to indicate a deposition time scale. From this figure is clear that the initial growth produces films with excess Se (or deficient in Zn). The stoichiometric regime is only reached around 30 s (around 60 monolayers, ML). The Zn/Se beam pressure ratio was maintained at -0.3 for all the films during the whole growth, so the excess in Se with respect to Zn must be attributed to the chemical composition of the interfacial region. It is well known that the heating process employed to eliminate the native oxides of the GaAs substrate produces surfaces rich in Ga; even Ga clusters can be
0.6 0 .Z E
LMM
-m-O-
0.2 ’ 0
I 100 deposition
transitions
fn / (Zn + Se) Se ! (Zn + Se)
I I 200 3oc time (set)
Fig. 4. Ratio of the peak to peak intensities of the LMM transitions of Zn and Se from Auger derivative spectra. The intensities were measured after 5, 15,30,60and 300 s deposition time.
expected. We infer from the Auger data that an interfacial layer, which is a mixture of ZnSe and a Ga-Se compound, is formed during the first 30 s of growth. The possibility of formation of a pseudomorphic GazSe3 film has been mentioned frequently [l, 23-251 but it is still a matter of active discussion; however, our results point towards this direction. Another possibility could be the formation of a more complex pseudomorphic Zn-Ga-Se interfacial compound. The RHEED patterns show a clear transition between a spotty pattern at the beginning of the growth and a streaky pattern after around 2&30 s deposition. Additionally, the Auger spectra show the Ga and As transitions before 30 s deposition, afterwards no signal from the substrate can be detected. From all the previous discussions we can conclude that, under the growth conditions employed, the interfacial region is formed during the first 30 s of growth, which is the time needed to cover the whole substrate and also time needed to go from a 3D to 2D growth mode. It is unclear if this 30 s deposition time represents -60 ML. The
516
I. Hernbndez-Calderbn
et al. ; Journal
initial growth can proceed at probably lower deposition rates, and could also vary as a function of substrate preparation. These results indicate that additional and systematic studies of the interfacial region are required to understand its nature and influence on the performance of ZnSe based devices.
4. Conclusions We employed PL, RS, HRXRD and AES to investigate the structural properties of MBE grown ZnSe/ GaAs heterostructures as a function of growth temperature and ZnSe film thickness. The presence of stress in the GaAs substrate was directly observed by PL. The results of RS and HRXRD show very good agreement and clearly indicate that the strain in the ZnSe film is inhomogeneous and depends on the film thickness, but not on the substrate temperature in the 285-325°C range. From these experiments we infer a value of 0.17 urn for h,. From the analysis of the Auger spectra it is concluded that the interfacial region is composed of a mixture of a GaaSe compound (probably Ga,Se,) and ZnSe. After the formation of this interfacial layer a stoichiometric regime is reached and the deposition proceeds in a 2D growth mode. We expect that these results will contribute to a better understanding of the ZnSe/GaAs interface. It is well known that important mechanisms of defect generation and propagation, that severely affect the performance of ZnSe-based greenblue emitting devices, take place in this interfacial region.
Acknowledgements This work was partially supported by CONACyT (Mexico). M.M.L. thanks the OAS and NSF (grant DMR-9521507) for partial support. We thank A. Guillen, H. Silva, Z. Rivera and M. Guerrero for their kind technical assistance.
References [l]
D.-W. Tu and A. Kahn, J. Vat. Sci. Technol. 922, and references therein.
A 3 (1985)
of Cystal Grow~th I75!176 (1997) 571-576 c21 M.A. Hasse, J. Qiu. J.M. DePuydt
and H. Cheng. Appl. Phys. Lett. 59 (1991) 1272. Semiconductors 2 (1996) 7. 131 Compound [41 K. Nakano, S. Tomiya, M. Ukita, H. Yoshida, S. Itoh, E. Morita, M. Ikeda and A. Ishibashi, J. Electron. Mater. 25 (1996) 213. r51 B.J. Wu, G.M. Haugen, J.M. Depuydt, L.H. Kuo and L. Salamanca-Riba, Appl. Phys. Lett. 68 (1996) 2828. C61 M.C. Tamargo, R.E. Nahory, B.J. Skromme, SM. Shibh. A.L. Weaver, R.J. Martin and H.H. Farrel, J. Crystal Growth 111(1991) 741. [71 G. Bratina, L. Vanzetti, A. Bonani, L. Sorba, J.J. Paggel, A. Franciosi, T. Peluso and L. Tapfer, J. Crystal Growth 159 (1996) 703. 181 K. Menda, I. Takayasu, T. Minato and M. Kawashima, Jpn. J. Appl. Phys. 26 (1987) Ll326; J.M. DePuydt, H. Cheng, J.E. Potts, T.L. Smith and SK. Mohapatra, J. Appl. Phys. 62 (1987) 4756. 191 H. Shigekawa, H. Oigawa, K. Miyake, Y. Aiso, Y. Nannichi, T. Hashizume and T. Sakurai, Appl. Phys. Lett. 65 (1994) 607. Cl01 Y. Hirose, S.R. Forrest and A. Kahn. Appl. Phys. Lett. 66 (1995) 944. M. Melendez-Lira, 0. de Melo, Z. Rivera [Ill L. Hernandez, and I. Hernandez-Calderon, J. Vat. Sci. Technol. A 14 (1996) 2269. W. Faschinger, P. Juza. A. Pesek, K. [I21 F. Heuzenberger, Lischka and H. Sitter, Thin Solid Films 225 (1993) 265. Cl31 V. Heim and P. Hiesinger. Phys. Status Solidi (b) 66 (1974) 461. Cl41 G.B. Stringfellow. W. Koschel and F. Briones, Appl. Phys. Lett. 39 (1981) 581. [I51 L. Pavesi and M. Guzzi, J. Appl. Phys. 75 (1994) 4780. [I61 D.J. Ashen and P.J. Dean, J. Phys. Chem. Solids 36 (1975) 1041. [I71 CD. Lee, B.K. Kim, J.W. Kim, S.K. Chang and S.H. Suh, J. Appl. Phys. 76 (1994) 928. Cl81 Y.R. Lee, A.K. Ramdas, L.A. Kolodziejski and R.L. Gunshor. Phys. Rev. B 38 (1988) 13143. 1191 I.K. Sou, S.M. Mou, Y.W. Chan, G.C. Xu and G.K.L. Wong, J. Crystal Growth 147 (1995) 39. 1201 D.J. Olego. K. Shahzad, J. Petruzzello and D. Cammack, Phys. Rev. B 36 (1987) 7674. H. c211 J. Reichow, J. Griesche, N. Hoffmann, C. Muggelberg, Rossmann, L. Wilde, F. Hennenberg and K. Jacobs, J. Crystal Growth 13 I (1993) 277. lx1 T. Yao, Y. Okada, S. Matsui, K. Ishida and I. Fujimoto, J. Crystal Growth 81 (1987) 518. cl31 I. Suemune, K. Ohmi, T. Kanda, K. Yatake. Y. Kan and M. Yamanishi, Jpn. J. Appl. Phys. 25 (1986) 1827. ~241 D. Li, M. Gonsalves, N. Otsuka, J. Qiu, M. Kobayashi and R.L. Gunshor, Appl. Phys. Lett. 57 (1990) 449. I251 A. Krost, W. Richter, D.R.T. Zahn, K. Hingerl and H. Sitter, Appl. Phys. Lett. 57 (1990) 1981.
Journal of Crystal Growth 175/176 (1997) 571-576
Investigation of the structural properties of MBE grown ZnSe/GaAs heterostructures I. Hernhdez-CaldercW*, E. Lbpez-Luna”, J. Luyo”, M. Mekndez-Lira”, 0. de Meloa,‘, P. Diaz”,‘, L. Hernhndeza”, J. Fuente?l, R. Lebn”,‘, H. Sitterb aPhysics Department, CINVESTAV, Apdo. Postal 14-740, 07000 Mbxico. D.F., Mexico b Institut ftir Experimentalphysik, Universitiit Linz, Linz, Austria
Abstract The results of photoluminescence and Raman spectroscopies, high resolution X-ray diffraction, and Auger electron spectroscopy are analyzed in terms of the structural properties of the ZnSe/GaAs(l 0 0) system as a function of film thickness and substrate temperature. The results of Raman spectroscopy and X-ray diffraction clearly show that the strain in the film is inhomogeneous and depends only on film thickness and not on growth temperature in the 285-325°C range. From these experiments a value of N 0.17 pm is inferred for the critical thickness of ZnSe on GaAs. Photoluminescence experiments sensitive to the ZnSe/GaAs interface reveal the presence of strain in the GaAs substrate. Analysis of the intensities of the LMM Auger transitions of Zn and Se indicate the formation of an interfacial layer with excess of Se, suggesting the formation of a pseudomorphic Ga-Se compound mixed with ZnSe at the interfacial region. PACS: 61.10. -i;
68.35.Ct; 68.55.Vk; 78.30.F~; 78.66. - w; 78.66.Hf
ZnSe; ZnSe/GaAs heterostructures; luminescence; Critical thickness
Keywords:
Semiconductor
1. Introduction
The study of ZnSe/GaAs interfaces has attracted the attention of researchers since the last decade [l]. However, the recent increase in the investigation of the physical properties of ZnSe/GaAs het-
* Corresponding author. E-mail: [email protected]. ’Permanent address: Faculty of Physics, University of Havana, Cuba. 0022-0248/97/$17.00 Copyright PII SOO22-0248(96)01215-S
0
interfaces;
Auger;
X-ray
diffraction;
Raman;
Photo-
erostructures is noteworthy. One of the main reasons is that the improvement of their structural and chemical properties is of fundamental relevance for the elaboration of green-blue emission lasers based on ZnSe [2]. Recently, Sony reported a green laser with a lifetime larger than 100 h [3]. The noticeable improvement was attributed to the reduction of defect density from lo6 to around lo4 cme2. It needs to be further reduced to inhibit defect propagation during laser operation [4]. Wu and coworkers very recently reported that low defect
1997 Elsevier Science B.V. All rights reserved
512
I. Hernirndez-Calderbn
et al. /Journal
density (less than lo4 cm- 2, pseudomorphic ZnSSe and ZnSe layers can be grown on an As-rich GaAs buffer layer initially exposed to Zn [S]. Some of the intrinsic defects at the interface are caused by the chemical and structural differences between the GaAs substrate and the ZnSe film. The lattice mismatch produces biaxially compressed films which release the strain after reaching the critical thickness through the formation of misfit dislocations. It has also been shown that the charge imbalance at the ZnSe/GaAs interface can lead to 3D nucleation and interface roughening [6]. Interdiffusion and crystalline degradation can occur as a result of variations in the ZnSe stoichiometry near the interface [7]. The quality of the epilayers is strongly affected by initial growth conditions such as substrate preparation, growth temperature, and Zn to Se flux ratios [S]. Here, we report results on the investigation of the structural properties of ZnSe/ GaAs(1 0 0) heterostructures grown by molecular beam epitaxy (MBE) by means of photoluminescence (PL) and Raman (RS) spectroscopies, high resolution X-ray diffraction (HRXRD) experiments, (AES). and Auger electron spectroscopy
2. Experimental
details
The films were grown in a RIBER 32P MBE system. A series of epitaxial ZnSe films were deposited on semi-insulating GaAs(1 0 0) substrates at temperatures between 28.5 and 325°C. Prior to their introduction to the MBE chamber the substrates were chemically treated as described in Ref. [l 11. The native oxide layer was removed by heating the indium glued substrate at about 550°C in ultrahigh vacuum (in the low lo-” Torr range). The reflection high energy electron diffraction (RHEED) experiment of the annealed surface showed a (2 x 1) reconstruction pattern, indicating that residual Se from the background atmosphere reacted with the GaAs(1 0 0) surface [S, 9, lo]. The typical Zn/ Se beam pressure ratio was -0.3. The RHEED patterns confirmed the Se rich growth through the ZnSe(2 x 1) surface reconstruction. A typical growth rate of - 1 pm/h was employed [ 111. The PL experiments were performed in a standard setup equipped with He-Ne and Kr-Ar lasers,
of Crystal Growth 175/l 76 (1997) 571~ 576
a 0.5 m monochromator and a closed cycle He refrigerator. The Raman experiments were done with a double monochromator equipped with CCD detection and the 4880 A line of an Ar laser. To verify the selection rules of RS we used a perfect backscattering x(z, z)X configuration; x, y, and z correspond to the (1 0 0) crystal directions. The HRXRD experiments were done using CuK,, radiation and a four-crystal Bartels monochromator [12]. The AES spectra were taken in an analysis chamber connected through UHV to the growth chamber. The Auger spectra were measured immediately after growth of the films employing a MAC3 analyzer set to 2 eV resolution; the energy of the primary electrons was 3 keV.
3. Results and discussion At room temperature (RT) the 0.26% lattice mismatch between GaAs (a = 5.6532 & RT) and ZnSe (a = 5.6676 A, RT) gives place to biaxial stress at the interface. Most of the elastic deformation is accumulated in the ZnSe film, and it is customary to neglect the substrate deformation. A precise analysis of the strain must take into account the differences between the thermal expansion coefficients of ZnSe and GaAs, since in most cases the temperature of the sample during characterization or device operation is different from the growth temperature. A systematic analysis of the deformation, both in the substrate and in the films, is needed to have a better understanding of the interface strain. With the purpose of directly identifying the presence of substrate strain, we performed PL measurements sensitive to the GaAs surface. Fig. la shows the PL of covered and clean GaAs substrates. Since ZnSe is transparent to the 4880 A photons (2.54 eV), we can easily reach the GaAs/ ZnSe interfacial region. We observe a broad transition around 1.52 eV associated to excitonic transitions [13] and carbon impurities [14]. Transitions around 1.45 eV have been attributed to GaAs antisites [15] and those around 1.41 eV to Mn impurities [16]. There are additional transitions at lower energies, but here we are only interested in their energy shifts with film thickness. It is worth mentioning that the PL spectra taken
et al. I Journal of‘Crystal Growth 175/I 76 (1997) 571-576
I. Hermindez-Calderbn
, ZnSelGaAs(100) Photolum~nescence br 4880A. 17 K
ating that the uppermost interfacial region detected with the 4880 A line has structural imperfections. At first glance the spectra shown in Fig. la could indicate that the substrate is not affected by the epilayer. However, a careful analysis of the PL peaks for samples with different thickness shows very small but noticeable and consistent energy shifts (see Fig. lb). The general trend is an increase in energy with ZnSe film thickness up to around 0.2 urn and then a slight decrease. Considering the reduced magnitude of these changes, we realize this description may appear a bit imprecise, but as we will see later it can be understood in terms of the evolution of the strain in the substrate and the film before and after the critical thickness is reached and of the difference in the thermal expansion coefficients of ZnSe and GaAs. As expected from high quality films, the RS experiments in the backscattering configuration, x(z, z)X, did not show the TO phonon of the ZnSe film but a very weak TO signal from GaAs, indicating some degree of disorder at the interface. Fig. 2a shows the spectra of the films grown at 325°C; the
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Fig. 1. (a) PL spectra of the clean and ZnSe covered substrates. (b) Energy shifts of the PL peaks of GaAs as a function of thickness of the ZnSe epilayer.
with the 6328 A line of a HeeNe ples a larger GaAs volume than exhibited sharper, stronger and solved excitonic peaks than those
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Fig. 2. (a) Raman spectra of ZnSe/GaAs heterostructures in a perfect backscattering configuration.(b) phonon ofZnSe as a function offilm thickness.(c) Changes in the FWHM of the LO phonon ofZnSe. the eye.
2.5 3.0 (pm)
Changes in the energy of the LO The straight lines arejust guides to
514
I. Hernbndez-Calderbn et al. I Journal of Crystal Growth 175/l 76 (I 997) 571-576
variations in intensity between the spectra are due to interference effects caused by different film thicknesses. Fig. 2b shows the observed shift in the energy of the ZnSe LO phonon and Fig. 2c the change in its full width at half maximum (FWHM) as a function of the ZnSe thickness. We observe that the LO peak moves towards lower energies up to around 0.17 urn and then remains at a value of 250.8 cm- r. The total shift is around 1 cm-‘. A similar shift was observed between thin and thick ZnSe films grown by chemical vapor deposition [17]. In Fig. 2c we observe that the FWHM of the LO phonon increases from around 4 cm-i for the thinnest film to around 8 cm- 1 for the 0.25 ym thick ZnSe film; afterwards it remains at around 7cmP’. Looking at the straight lines (which are drawn just as a guide to the eye), we can describe the general behavior of the FWHM as an increase to around 0.18 urn and constant afterwards. Then, from the evolution of the LO energy shift and its FWHM, we can conclude that the changes depend on the thickness but not on the substrate temperature. These changes also indicate that the critical thickness (h,) of ZnSe in GaAs(1 0 0) is around 0.17 urn. This value is consistent with previously reported values [lS, 191. The fact that the FWHM shows an important increase around h, is an indication of the presence of structural disorder produced by the appearance of misfit dislocations which release the strain caused by the lattice mismatch. The observation in Fig. 2 that the average compressive biaxial strain of the ZnSe film is continuously reduced with increasing film thickness up to h,, is a direct indication of nonuniform (or inhomogeneous) stress in the films, in agreement with previous results of Olego et al. [20]. In Fig. 3 we have summarized the results of HRXRD experiments. Fig. 3a presents the difference between the angle of the (0 0 4) reflection of ZnSe with respect to that of GaAs as a function of film thickness. We can see a very close similarity in the behavior of d@ and that of the LO phonon energy shown in Fig. 2a. Again, no dependence in the substrate temperature is observed. However, for HRXRD the changes reveal directly the modification of the (1 0 0) interplanar distance parallel to the substrate, indicating a value of h, around 0.17 urn, consistent with the Raman measurements.
0.24
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.
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.
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.
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Fig. 3. (a) Difference between the angle of the (0 0 4) reflection of ZnSe and GaAs as a function of film thickness. (b) Changes in the FWHM of the (0 0 4) reflection of ZnSe as a function of film thickness. The lines are just guides to the eye.
The presence of residual strain is observed even for the thickest film. Fig. 3b shows the modification of the FWHM of the (0 0 4) reflection of the HRXRD patterns of the heterostructures. Analogous to the previous findings, we can see two types of behavior: a fast decrease up to around h, and then a very slow decrease with thickness. Recent similar experiments by other groups show a variety of behaviors of the FWHM of the (0 0 4) reflection with ZnSe film thickness. Reichow et al. [21] found a continuous decrease of the FWHM with thickness of MBE deposited films, very close to our results. The report of Sou et al. [19] does not show a clear trend of the FWHM with film thickness in MBE deposited films. Lee et al. [17] found a reduction of the FWHM with film thickness of chemical vapor deposited films. Then, apparently the modification of the FWHM of the (0 0 4) reflection depends on the type and conditions of growth. In our case, we attribute the continuous decrease of the FWHM to
I. Hernirndez-Calderh
515
et al. 1 Journal ofCyysta1 Growth 175/l 76 (1997) 571-576
the improvement of the overall film structural quality with thickness. We start with a relatively rough substrate and the film becomes smooth with increasing ZnSe thickness. The apparent discrepancy between the conclusions from Fig. 2c and Fig. 3b is easily understood when we consider that Raman spectroscopy is very sensitive to the microscopic structure, while HRXRD is more sensitive to the average macroscopic structure of the films. Based on the previous results it is easier to explain the shifts observed in PL in Fig. la. First, we have to consider the fact that the measurements were done at low temperature. Since the thermal expansion coefficient of ZnSe is larger than that of GaAs [22], after cooling to 17 K the situation will be reversed in relation to the one at growth temperature, the ZnSe film will be under biaxial tensile stress and the GaAs interfacial region will suffer biaxial compressive strain. The result is an increase in the band gap of GaAs with a maximum value just around h,. The transitions observed in the PL spectra of Fig. 1 indicate this behavior. After h, is reached most of the strain is released through the misfit dislocations and then the band gap energy of GaAs tends to decrease. Then, from the PL measurements we have obtained a direct indication of substrate strain. We performed Auger experiments of the ZnSe/ GaAs heterostructure as a function of ZnSe deposition of very thin films. Fig. 4 illustrates the ratio of the peak to peak intensities of the LMM transitions of Zn and Se from Auger derivative spectra. The intensities were measured at 5, 15, 30, 60 and 300 s deposition time. The typicalOgrowth rate of the films was - 1 urn/h (- 2.8 A/s); however, since growth conditions can change with film thickness we prefer to indicate a deposition time scale. From this figure is clear that the initial growth produces films with excess Se (or deficient in Zn). The stoichiometric regime is only reached around 30 s (around 60 monolayers, ML). The Zn/Se beam pressure ratio was maintained at -0.3 for all the films during the whole growth, so the excess in Se with respect to Zn must be attributed to the chemical composition of the interfacial region. It is well known that the heating process employed to eliminate the native oxides of the GaAs substrate produces surfaces rich in Ga; even Ga clusters can be
0.6 0 .Z E
LMM
-m-O-
0.2 ’ 0
I 100 deposition
transitions
fn / (Zn + Se) Se ! (Zn + Se)
I I 200 3oc time (set)
Fig. 4. Ratio of the peak to peak intensities of the LMM transitions of Zn and Se from Auger derivative spectra. The intensities were measured after 5, 15,30,60and 300 s deposition time.
expected. We infer from the Auger data that an interfacial layer, which is a mixture of ZnSe and a Ga-Se compound, is formed during the first 30 s of growth. The possibility of formation of a pseudomorphic GazSe3 film has been mentioned frequently [l, 23-251 but it is still a matter of active discussion; however, our results point towards this direction. Another possibility could be the formation of a more complex pseudomorphic Zn-Ga-Se interfacial compound. The RHEED patterns show a clear transition between a spotty pattern at the beginning of the growth and a streaky pattern after around 2&30 s deposition. Additionally, the Auger spectra show the Ga and As transitions before 30 s deposition, afterwards no signal from the substrate can be detected. From all the previous discussions we can conclude that, under the growth conditions employed, the interfacial region is formed during the first 30 s of growth, which is the time needed to cover the whole substrate and also time needed to go from a 3D to 2D growth mode. It is unclear if this 30 s deposition time represents -60 ML. The
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initial growth can proceed at probably lower deposition rates, and could also vary as a function of substrate preparation. These results indicate that additional and systematic studies of the interfacial region are required to understand its nature and influence on the performance of ZnSe based devices.
4. Conclusions We employed PL, RS, HRXRD and AES to investigate the structural properties of MBE grown ZnSe/ GaAs heterostructures as a function of growth temperature and ZnSe film thickness. The presence of stress in the GaAs substrate was directly observed by PL. The results of RS and HRXRD show very good agreement and clearly indicate that the strain in the ZnSe film is inhomogeneous and depends on the film thickness, but not on the substrate temperature in the 285-325°C range. From these experiments we infer a value of 0.17 urn for h,. From the analysis of the Auger spectra it is concluded that the interfacial region is composed of a mixture of a GaaSe compound (probably Ga,Se,) and ZnSe. After the formation of this interfacial layer a stoichiometric regime is reached and the deposition proceeds in a 2D growth mode. We expect that these results will contribute to a better understanding of the ZnSe/GaAs interface. It is well known that important mechanisms of defect generation and propagation, that severely affect the performance of ZnSe-based greenblue emitting devices, take place in this interfacial region.
Acknowledgements This work was partially supported by CONACyT (Mexico). M.M.L. thanks the OAS and NSF (grant DMR-9521507) for partial support. We thank A. Guillen, H. Silva, Z. Rivera and M. Guerrero for their kind technical assistance.
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