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Applications of defect engineering in InP-based structures
Materials Science and Engineering B75 (2000) 103 – 109 www.elsevier.com/locate/mseb
Applications of defect engineering in InP-based structures W.M. Chen a,*, I.A. Buyanova a, C.W. Tu b b
a Department of Physics and Measurement Technology, Linko¨ping Uni6ersity, S-581 83 Linko¨ping, Sweden Department of Electrical and Computer Engineering, Uni6ersity of California, La Jolla, CA 92093 -0407, USA
Abstract Recent developments in defect engineered InP-based structures, by grown-in intrinsic defects, are reviewed. We demonstrate that n-type doping or modulation doping in InP-based structures can be realized by an intentional introduction of PIn antisites during off-stoichiometric growth of InP at low temperatures (LT) ( 260 – 350°C) by gas source molecular beam epitaxy (GS-MBE), without requiring an external shallow impurity doping source. We shall first summarize our present understanding of the mechanism responsible for the n-type conductivity of LT-InP, which is attributed to the auto-ionization of PIn antisites via the (0/+) level resonant with the conduction band. The PIn antisites are shown to exhibit properties meeting basic requirements for a dopant: (1) known chemical identification; (2) known electronic structure; (3) a control of doping concentration by varying growth temperature. We shall also provide a review of recent results from defect engineering, by utilizing the intrinsic n-type dopants of PIn antisites for modulation doping in InP-based heterostructures. Important issues such as doping efficiency, electron mobility, thermal stability, etc., will be addressed, in a close comparison with the extrinsically doping method by shallow dopants. © 2000 Elsevier Science S.A. All rights reserved. Keywords: Defects; Engineering; Doping; InP; Heterostructures; MBE
1. Introduction Defects and impurities have played a major role in determining the currently important status of semiconducting materials in modern electronics. This is largely due to their ability in drastically tailoring electronic properties of the materials and in manipulating the performance of devices. While many of the defects in semiconductors are known to be harmful and are responsible for degradation of materials and devices, some have been found to be extremely valuable. The latter include, for example, shallow dopants providing charge carriers and deep-level defects for charge compensation and carrier lifetime control. A full exploitation of defects in altering semiconductor properties requires, however, a precise control and a delicate engineering of defects. In fact defect engineering has been a focal point ever since semiconductor technology revolutionized the electronics industry. Rapid developments in innovative growth techniques and new device structures in recent years have opened new possibilities for defect engineering beyond the * Corresponding author.
reach of traditional methods. Among them, off-stoichiometric growth of semiconductors has gained wide interest due to unique properties provided by such off-stoichiometric materials. For an example, highly As-rich GaAs grown at low temperatures (LT) by molecular beam epitaxy (MBE) exhibits a unique combination of electronic properties, such as extremely high conductivity and ultra-short carrier lifetime, which has found applications for device isolation, ultra-fast optical detectors, etc. [1]. Stimulated by the outcome of LT-GaAs and by the long-sought desire to obtain undoped, semi-insulating InP, an attempt has been made in recent years to grow off-stoichiometric InP at low temperatures by MBE. Though the material is highly P-rich, it is not at all semi-insulating. In fact, LT-InP becomes highly n-type conducting when the growth temperature is below normal temperatures [2–6]. In this paper, we shall first summarize our present understanding of the mechanism responsible for the n-type conductivity of LT-InP, which is attributed to the auto-ionization of PIn antisites preferentially introduced during the off-stoichiometric growth. We shall then review recent results from defect engineering, by utilizing the intrinsic n-type do-
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pants of the PIn antisites for modulation doping in InGaAs/InP heterostructures.
2. PIn-antisite engineered intrinsic doping in InP The first type of samples studied in this work is single InP epilayers, grown by gas source MBE, on semi-insulating Fe-doped InP substrate (Fig. 1a). The InP epilayers are typically 1-mm thick. The growth temperature varied from 480°C down to 265°C. Within this temperature range, the material remains good crystalline quality without a noticeable presence of structural defects. At growth temperatures below 200°C, on the other hand, the crystal quality of the material deteriorates drastically and eventually becomes polycrystalline. The free electron concentration of the LT-InP films was found to monotonically increase with decreasing
Fig. 3. Relative free electron concentration, measured via infrared absorption by free electrons, as a function of the hydrostatic pressure applied to the LT-InP epilayers. The observed carrier freeze-out occurs when the (0/ +) level of the PIn antisite enters the bandgap at a pressure higher than 1.7 GPa.
Fig. 1. Schematic pictures of the design for (a) the thick InP layers and (b) the InGaAs/InP heterostructures reviewed in this paper.
Fig. 2. Free electron concentration as a function of growth temperature in LT-InP epilayers. The insert shows the Fermi level pinning at the (0/+) level of the PIn antisite which leads to the saturation of the free electron concentration at growth temperatures lower than 265°C. C.B. denotes the conduction band.
growth temperatures [2,3] (Fig. 2). At 265°C a saturation electron concentration of 3×1018 cm − 3 was reached when the LT-InP exhibited a metallic n-type conduction. This saturation corresponds to a Fermi level pinning at the responsible donor level, i.e. the (0/+ ) level of the PIn antisite double donor to be discussed below. This donor level is resonant with the conduction band at Ec + 0.12 eV (see the inset in Fig. 2) estimated from the filling of the conduction band states [3]. The presence of this donor level was further confirmed by a freeze-out of the free electrons when the donor level enters the forbidden bandgap as a consequence of a widening bandgap upon the application of a hydrostatic pressure (Fig. 3). An analysis of the pressure coefficients gives an estimate of the resonant donor level at Ec + 0.11 eV, agreeing within the experimental error with the value of Ec + 0.12 eV given above [3]. The chemical nature of the donor giving rise to the n-type conductivity in LT-InP was revealed by the microscopically informative optically detected magnetic resonance (ODMR) technique. The ODMR studies have unambiguously shown the presence of the PIn antisite in the material, from its hyperfine fingerprint of the nuclear spin I=1/2 of the 31P atom with 100% natural abundance [3,5]. The presence of the PIn antisite was shown to correlate with a lowering of the growth temperature. If the PIn antisite is responsible for the n-type conductivity of LT-InP, the number of the ionized PIn antisites (P+ In ) should equal the free electron
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concentration. This close correlation as a function of growth temperature has, in fact, been provided experimentally as shown in Fig. 4. Here the concentration of P+ In can be scaled with the intensity of the magnetic circular dichroism (MCD) in absorption (MCDA) originating from the ground state of the singly ionized P+ In antisite. Such a direct correlation provides a convincing piece of evidence that the PIn antisite is indeed responsible for the n-type conductivity of LT-InP, via the auto-ionization of the (0/ + ) level resonant with the conduction band at Ec +0.12 eV. These experimental observations proved that the ntype conductivity of LT-InP is an intrinsic property of this off-stoichiometric material, and the PIn antisite is not at all suitable to make the material semi-insulating. On the contrary, it provides a new mechanism for n-type doping in InP, i.e. intrinsic doping without requiring an external doping source. In fact, the PIn
Fig. 4. Correlation between the free-electron concentration and the MCDA intensity (scaling with the concentration) of the P+ In antisite in LT-InP as a function of the growth temperature Tg.
Fig. 5. A schematic picture of the energy band diagram for the n-type modulation doped InGaAs/InP heterostructures. For simplicity only two subbands of the 2DEG are drawn. The filling of the 2DEG is depicted by the shaded region. The arrow across the InGaAs bandgap indicates the PL transition between the 2DEG and photo-excited holes upon optical excitation, shown in Fig. 7.
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antisite possesses properties that are regarded as basic requirements for a dopant, i.e. a known chemical identification: PIn; established electronic properties: the (0/+ ) level at Ec + 0.12 eV; a control of doping concentration: by varying Tg. Defect engineering by the PIn antisite can therefore provide a new doping mechanism by exploring such an intrinsic n-type dopant.
3. PIn-antisite engineered intrinsic modulation doping in InGaAs/InP heterostructures When the intrinsic doping concept is used in InPbased heterostructures, PIn-antisite engineered intrinsic modulation doping can readily be realized without invoking an external doping source. The success of such defect engineering was clearly demonstrated [7] by the first example cases of lattice-matched In0.53Ga0.47As/InP heterostructures (Fig. 1b) designed to resemble high electron mobility transistor (HEMT) structures. Three types of InGaAs/InP structures were grown and studied, all were grown by GS-MBE, on semi-insulating Fe-doped InP substrate: (1) intrinsically doped InGaAs/InP structures where the top InP layer was grown at a low temperature of 265°C (denoted as InGaAs/LT-InP); (2) extrinsically doped InGaAs/InP structures by shallow Si donors grown at normal temperature of 480°C; (3) undoped InGaAs/InP structures grown at normal high temperatures (HT) of 480°C (referred to as InGaAs/HT-InP or the reference sample). The design of these structures is identical except for the top InP layer and is shown in Fig. 1b. For the intrinsically doped InGaAs/InP, the entire structures were intentionally undoped and were grown at a normal temperature of 480°C except the top InP layer which was grown at 265°C. The reference sample was also grown with an identical structure except that the top InP layer was in this case grown at 480°C. In the extrinsically doped InGaAs/InP structures, the top InP was grown at normal temperature (480°C) but was intentionally doped by shallow Si donors. These structures were experimentally investigated by transport and optical techniques, mainly Shubnikov-de Haas (SdH) oscillations, Hall measurements, photoluminescence (PL), PL excitation (PLE) and magneto-optical spectroscopy. In Fig. 5 we first show a schematic energy band diagram of the modulation doped InGaAs/InP structures. Due to electron transfer from the n-type doped (either intrinsically or extrinsically) InP barrier to the InGaAs active layer, a notch potential is formed near the heterointerface where a two-dimensional electron gas (2DEG) is confined (Fig. 5).
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Fig. 6. SdH oscillation spectra taken at 1.5 K in dark from (a) the intrinsically doped InGaAs/LT-InP heterostructure and (b) the undoped reference sample, with the external magnetic field normal to the heterointerface plane. The total sheet concentrations of the 2DEG were determined from the analysis of the SdH oscillation spectra to be 1.15× 1012 and 2.37× 1011 cm − 2, respectively.
Fig. 7. PL spectra at 1.6 K obtained from the InGaAs/LT-InP structure, due to recombination of the 2DEG and photo-excited holes shown by the arrow in Fig. 5. The insert provides a close-up of the PL emission, at zero magnetic field (dashed curve) and 5 T (solid curve).
The first experimental evidence on the success of the intrinsic modulation doping concept was provided by the formation of a 2DEG in the InGaAs channel due to the intrinsic modulation doping, by a straightforward SdH oscillation study (Fig. 6a). From a detailed analysis of the SdH oscillations, two subbands of the 2DEG were shown to be readily occupied. The sheet concentrations of the first and second subbands of the 2D gas were determined to be n1 =6.75 ×1011 cm − 2 and n2 = 4.75× 1011 cm − 2, respectively, yielding a total sheet
concentration of 1.15×1012 cm − 2. This is to be compared with a much lower 2DEG concentration (2.37× 1011 cm − 2) observed in the undoped reference sample (Fig. 6b) taken under the same experimental conditions. Only one subband was partially occupied in this case, and was mainly induced by electron transfer from residual donors in the undoped InP buffer layer [8]. The 2D character of the 2DEG was revealed from an angular dependence study of the SdH oscillations, where the period of the SdH oscillations obeyed a cosine relation of the relative angle between the magnetic field and the growth direction [9]. Additional support for the success of the intrinsic modulation doping was provided from optical studies carried out on the same InGaAs/InP structures. In Fig. 7, a PL spectrum obtained from the intrinsically doped InGaAs/LT-InP structure is shown. The photon energy of the PL emission at around 0.8 eV is consistent with the bandgap of In0.53Ga0.47As. The PL spectrum was attributed to the recombination between the 2DEG and the photo-excited holes in the InGaAs layer, illustrated by an arrow across the InGaAs bandgap in Fig. 5. A similar PL emission has been observed in InGaAs/InP and InGaAs/AlInAs HEMT structures [10,11] where the conventional shallow donors were employed to achieve n-type modulation doping. The Landau level splitting of the PL emission in the magnetic field, shown in the insert of Fig. 7, and the deduced electron mass value m *: 0.05m0 further confirmed that the e recombination involves the 2DEG in the In0.53Ga0.47As channel [12], differing from the electron mass value 0.07m0 in InP. The 2D nature can again be shown m *: e by an angular dependence study of the Landau level splitting. Optically detected quantum oscillations (ODQO) [13] identical to the SdH oscillations were observed by detecting this PL emission at the highest photon energy near the Fermi edge, proving that the PL emission originates from the same 2DEG which give rise to the electrical SdH oscillations. The electron mobility of the 2DEG determined from the Hall measurements is me : 1.3× 104 cm2 V − 1 s − 1, which is much higher than me : 5× 102 cm2 V − 1 s − 1 determined independently for a single LT-InP layer grown at the same temperature (i.e. 265°C) [3]. This again confirmed that the 2DEG is located in the InGaAs channel due to electron transfer from the LT-InP doping region, and a much improved electron mobility is obtained due to a space separation of electrons from their parent donors.
4. Important issues in defect engineering by intrinsic doping Next, we shall review recent results from studies of a number of important issues such as doping efficiency,
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mobility and thermal stability, which are relevant and in fact determine the value of such defect engineering in practical device applications. A close comparison with the properties of the extrinsically modulation doped structures has been made, and the results are summarized in Table 1.
4.1. Doping efficiency As discussed above and also given in Table 1, a dense 2DEG is formed in the intrinsically modulation doped structures, with the total sheet concentration of ns = 1.15× 1012 cm − 2. This value is noticeably higher than ns =8.4×1011 cm − 2 in the extrinsically modulation doped structure with a Si-doping concentration of 1 × 1018 cm − 3, but comparable to that with a Si-doping concentration of 3× 1018 cm − 3 (ns =1.28× 1012 cm − 2). This agrees with a similar concentration of the PIn-antisite in the LT-InP layer grown at 265°C, determined in earlier studies of 1-mm-thick LT-InP thin films [2,3]. Our earlier self-consistent theoretical calculations of subband populations in the intrinsically modulation doped structure predicted the 2DEG concentration to be about a factor of two higher than that in the Si-doped structure. The present work seems to indicate that this earlier prediction has underestimated the 2DEG concentration in the extrinsically doped structures due to the assumption that the Fermi level is pinned at the donor level in thermal equilibrium. A much lower 2DEG concentration of 2.37× 1011 cm − 2 is formed in the reference structure as expected. The formation of the 2DEG in this nominally undoped structure is caused by an n-type background doping. The corresponding residual donor concentration is estimated to be about 3 ×1016 cm − 3 from our self-consistent theoretical calculations of the reference sample [8].
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It should be noted that the Hall concentration is markedly higher than the 2DEG concentration determined by the SdH oscillations in the intrinsically doped structure (Table 1). Such a deviation suggests the presence of a strong parallel conduction, due to the fact that a parallel conduction of lower mobility contributes much less to the SdH oscillation than to the Hall effect measurements. Possible parallel conduction channels available in the structure are the LT-InP doped region and the InP buffer layer. The SI InP substrate is not expected to contribute to electrical conduction. The contribution from the InP buffer layer in parallel conduction is due to the presence of a residual donor concentration of 3× 1016 cm − 3 [8]. The parallel conduction provided by the LT-InP layer can be attributed to the remaining 90% of the free electrons resulting from the autoionization of the PIn-antisite, after about 10% of them have transferred to the InGaAs channel.
4.2. Electron mobility The strong parallel conduction by the LT-InP layer of a lower mobility, independently determined to be about 500 cm2 V − 1 s − 1 in earlier studies [3], inevitably leads to a lower total Hall mobility as shown in Table 1. A reliable value for the 2DEG mobility requires a detailed analysis of the parallel conduction, by nHmH = Snimi and nHm 2H = Snim 2i , taking into account all the possible parallel conduction channels [14]. Here nH (or mH) and ni (or mi) represent the electron densities (or mobilities) of the total and individual conduction channels, respectively. Apart from the parallel channels by the LT-InP top layer and the buffer layer, the 2DEG by itself consists of two parallel conduction channels due to the occupation of two subbands. Though the concentration of the PIn-antisite and the electron mobility in
Table 1 The 2DEG concentrations (n) and mobilities (m) from the intrinsically and the extrinsically doped InP/InGaAs structures reviewed in this worka Parameters
From SdH
From Hall
Intrinsically doped structure
Extrinsically doped structure ([Si]: 1×1018 cm−3)
n1 (cm−2) m1 (cm2 V−1 s−1) n2 (cm−2) m2 (cm2 V−1 s−1) ntotal (cm−2)
6.75×1011 1.55×104 4.75×1011 3.05×104 1.15×1012
6.2×1011 7.0×103 2.2×1011 1.95×104 0.84×1012
nH (cm−2) mH (cm2 V−1 s−1)
5.7×1012 1.3×104 (3.1×104)b
0.8×1012 2.4×104
Extrinsically doped structure ([Si]: 3×1018 cm−3) 7.6×1011 5.2×1011 1.28×1012
a For the latter structures, two sets of samples were investigated with different doping concentrations of the shallow Si donor. Both the quantum mobilities m1 and m2 deduced from SdH and the Hall mobilities are given. The subscripts 1 and 2 denote the 1st and the 2nd subband of the 2DEG, respectively. b From Ref. [15].
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remote dopants is reduced in the intrinsically doped structure, which can be attributed to a different defect potential and a stronger screening of the defect potential by a degenerate electron gas present in the doped LT-InP barrier due to auto-ionization of the PIn antisite.
4.3. Thermal stability
Fig. 8. Optically detected quantum oscillation spectra taken at 1.5 K from the intrinsically doped InGaAs/LT-InP heterostructures, showing the effect of thermal annealing on the 2DEG concentration with or without surface protection.
the 1-mm-thick LT-InP films are known from earlier studies [3], they may be altered to some extent due to the increasing importance of surface states. Consequently, there are more unknown parameters than the number of the equations for parallel conduction that rules out a unique and reliable determination of the Hall mobility for each of the 2DEG subbands. Independent determination of some of the unknown parameters or a reduction in the number of parallel conduction channels is required to resolve the uncertainty. One approach of avoiding parallel conduction is to prevent the presence of free electrons in the LT-InP layer by applying hydrostatic pressure so that the (+ / 0) level of the PIn-antisite enters the bandgap and freezes electrons [15]. The Hall mobility deduced by this approach yields a value of 3.1×104 cm2 V − 1 s − 1, higher than 2.4× 104 cm2 V − 1 s − 1 deduced in this work from the extrinsically doped structure. In addition, the quantum mobility of the 2DEG deduced from an analysis of fast Fourier transformation of the SdH oscillations [16] is about a factor of 1.5 – 2 higher in the intrinsically doped structure compared to the extrinsically doped structure (Table 1). The difference in the Hall and quantum mobility between the intrinsically and extrinsically doped structures can only be due to the difference in scattering by remote dopants, since that is the only difference between the two structures. Scattering by alloy disorder, interface and background impurities and interband scattering should be nearly identical for both structures, and should not be responsible for the difference. This thus provides evidence that scattering of the 2DEG by
One of the other important issues which is relevant to practical device applications is the thermal stability of the intrinsic doping. This was addressed by examining the corresponding change in the 2DEG concentration upon annealing of the intrinsically modulation doped InGaAs/LT-InP heterostructures at 400–500°C. The annealing was performed both with and without a proper surface protection. For the surface protection, two methods have been used. The first method is proximity annealing, when the sample surface was protected by another piece of InP substrate. The second method is to overgrow an InP cap layer at 480°C on top of the LT-InP, which can serve both to anneal the underlying structure at 480°C and to protect the LTInP surface. A summary of the results is shown in Fig. 8. As can be seen, the annealing at 420°C for the intrinsically doped structure without surface protection caused a significant reduction (more than half) of the 2DEG concentration. With the surface protection, either by proximity annealing or overgrowth, the reduction of the 2DEG concentration has been markedly improved. This experimental finding demonstrated that the instability of the InP surface was the main reason for the thermal degradation of the structures. This is somewhat expected since an InP surface has been known to be unstable at T\ 400°C. It should be pointed out that the extrinsically modulation doped InGaAs/InP heterostructures have been found to exhibit a higher thermal stability, where the 2DEG concentration was virtually unaffected by annealing. This seems to indicate that the surface of the P-rich LT-InP is less stable than that of the Si-doped InP. We should point out, however, that even with the surface protection a noticeable decrease was observed in the 2DEG concentration after annealing of the intrinsically doped InGaAs/LT-InP heterostructures. Thus, it should be related to the thermal induced change in the bulk electrical properties of the LT-InP layer. This could be attributed either to a reduction in the number of the PIn antisites due to complex formation and/or precipitation or to a thermally induced formation/activation of compensating acceptors in the LT-InP layer, upon annealing. A decrease in conductivity of the thick LT grown InP epilayer after annealing has indeed been previously observed and has been attributed [17] to the thermal activation of acceptor impurities, which are passivated in the as-grown mate-
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rial by hydrogen, known to incorporate efficiently in the InP material during the LT MBE growth.
the U.S. Air Force Office of Scientific Research (AFOSR, F49620-93-1-0367TU).
5. Summary
References
A review is given on a successful example of defect engineering in InP-based heterostructures. Here an intrinsic defect, i.e. the PIn antisite, has been shown to possess the fundamental properties required for an n-type dopant, that is: (1) known chemical identification; (2) known electronic structure; (3) a control of doping concentration by varying growth temperature. Engineering of the PIn antisite in off-stoichiometrically grown InP at low growth temperatures has provided an intrinsic n-type doping source without invoking an external doping source. The success of the defect engineering has been clearly demonstrated by recent results from both the thick LT-InP epilayers and also the InGaAs/LT-InP heterostructures. Important physical properties of the intrinsic modulation doping have been discussed, such as doping efficiency, mobility and thermal stability, as compared to the extrinsically doped structures. Though it has only been demonstrated for InGaAs/InP heterostructures, the principle of defect engineering within the intrinsic doping concept is in fact rather general and can be extended to applications in other electronic material systems.
[1] G.L. Witt, R. Calawa, U. Mishra, E. Weber (Eds.), Low Temperature (LT) GaAs and Related Materials, MRS Symposium Proceedings, Vol. 241, Materials Research Society, Pittsburgh, PA, 1992. [2] B.W. Liang, P.Z. Lee, D.W. Shih, C.W. Tu, Appl. Phys. Lett. 60 (1992) 2014. [3] P. Dreszer, W.M. Chen, K. Seendripu, J.A. Wolk, W. Walukiewicz, B.W. Liang, C.W. Tu, E.R. Weber, Phys. Rev. B47 (1993) 4111. [4] W.M. Chen, P. Dreszer, A. Prasad, A. Kurpiewski, W. Walukiewicz, E.R. Weber, E. So¨rman, B. Monemar, B.W. Liang, C.W. Tu, J. Appl. Phys. 76 (1994) 600. [5] W.M. Chen, P. Dreszer, E.R. Weber, E. So¨rman, B. Monemar, B.W. Liang, C.W. Tu, J. Electron. Mater. 22 (1993) 1491. [6] P. Dreszer, W.M. Chen, D. Wasik, R. Leon, W. Wulukiewicz, B.W. Liang, C.W. Tu, E.R. Weber, J. Electron. Mater. 22 (1993) 1487. [7] W.M. Chen, I.A. Buyanova, A.V. Buyanov, T. Lundstro¨m, W.G. Bi, C.W. Tu, Phys. Rev. Lett. 77 (1996) 2734. [8] I.A. Buyanova, T. Lundstro¨m, A.V. Buyanov, W.M. Chen, W.G. Bi, C.W. Tu, Phys. Rev. B55 (1997) 7052. [9] W.M. Chen, I.A. Buyanova, W.G. Bi, C.W. Tu, Mater. Sci. Forum 258-263 (1997) 805. [10] D.G. Hayes, M.S. Skolnick, D.M. Whittaker, P.E. Simmonds, L.L. Taylor, S.J. Bass, L. Eaves, Phys. Rev. B44 (1991) 3436. [11] Y.-H. Zhang, D.-S. Jiang, R. Cingolani, K. Ploog, Appl. Phys. Lett. 56 (1990) 2195. [12] C.K. Sarkar, R.J. Nicholas, J.C. Portal, M. Razeghi, J. Chevrier, J. Massies, J. Phys. C18 (1985) 2667. [13] I.A. Buyanova, W.M. Chen, A. Buyanov, W.G. Bi, C.W. Tu, Appl. Phys. Lett. 69 (1996) 809. [14] R.L. Pertritz, Phys. Rev. 110 (1958) 1254. [15] D. Wasik, L. Dmowski, J. Mikucki, J. Lusakowski, L. Hsu, W. Walukiewicz, W.G. Bi, C.W. Tu, Mater. Sci. Forum 258-263 (1997) 813. [16] E. Skuras, R. Kumar, R.L. Williams, R.A. Stradling, J.E. Dmochowski, E.A. Johnson, A. Mackinnon, J.J. Harris, R.B. Beall, C. Skierbeszewski, J. Singleton, P.J. van der Wel, P. Wisniewski, Semicond. Sci. Technol. 6 (1991) 535. [17] B.W. Liang, Y. He, C.W. Tu, in: G.L. Witt, R. Calawa, U. Mishara, E. Weber (Eds.), Low Temperature (LT) GaAs and Related Materials, MRS Symposia Proceedings 241, Materials Research Society, Pittsburgh, PA, 1992, p. 297.
Acknowledgements We are grateful to W.G. Bi, A.V. Buyanov, T. Lundstro¨m, P. Dreszer, E.R. Weber, W. Walukiewicz, B.W. Liang, B. Monemar, E. So¨rman, J.A. Wolk, D. Wasik, H.J. Sun and G.D. Watkins for valuable contributions in connection to this work. This work has been supported by the Swedish Council for Engineering Sciences (TFR), the Swedish Natural Science Research Council (NFR) and an AASERT grant from
.
Applications of defect engineering in InP-based structures W.M. Chen a,*, I.A. Buyanova a, C.W. Tu b b
a Department of Physics and Measurement Technology, Linko¨ping Uni6ersity, S-581 83 Linko¨ping, Sweden Department of Electrical and Computer Engineering, Uni6ersity of California, La Jolla, CA 92093 -0407, USA
Abstract Recent developments in defect engineered InP-based structures, by grown-in intrinsic defects, are reviewed. We demonstrate that n-type doping or modulation doping in InP-based structures can be realized by an intentional introduction of PIn antisites during off-stoichiometric growth of InP at low temperatures (LT) ( 260 – 350°C) by gas source molecular beam epitaxy (GS-MBE), without requiring an external shallow impurity doping source. We shall first summarize our present understanding of the mechanism responsible for the n-type conductivity of LT-InP, which is attributed to the auto-ionization of PIn antisites via the (0/+) level resonant with the conduction band. The PIn antisites are shown to exhibit properties meeting basic requirements for a dopant: (1) known chemical identification; (2) known electronic structure; (3) a control of doping concentration by varying growth temperature. We shall also provide a review of recent results from defect engineering, by utilizing the intrinsic n-type dopants of PIn antisites for modulation doping in InP-based heterostructures. Important issues such as doping efficiency, electron mobility, thermal stability, etc., will be addressed, in a close comparison with the extrinsically doping method by shallow dopants. © 2000 Elsevier Science S.A. All rights reserved. Keywords: Defects; Engineering; Doping; InP; Heterostructures; MBE
1. Introduction Defects and impurities have played a major role in determining the currently important status of semiconducting materials in modern electronics. This is largely due to their ability in drastically tailoring electronic properties of the materials and in manipulating the performance of devices. While many of the defects in semiconductors are known to be harmful and are responsible for degradation of materials and devices, some have been found to be extremely valuable. The latter include, for example, shallow dopants providing charge carriers and deep-level defects for charge compensation and carrier lifetime control. A full exploitation of defects in altering semiconductor properties requires, however, a precise control and a delicate engineering of defects. In fact defect engineering has been a focal point ever since semiconductor technology revolutionized the electronics industry. Rapid developments in innovative growth techniques and new device structures in recent years have opened new possibilities for defect engineering beyond the * Corresponding author.
reach of traditional methods. Among them, off-stoichiometric growth of semiconductors has gained wide interest due to unique properties provided by such off-stoichiometric materials. For an example, highly As-rich GaAs grown at low temperatures (LT) by molecular beam epitaxy (MBE) exhibits a unique combination of electronic properties, such as extremely high conductivity and ultra-short carrier lifetime, which has found applications for device isolation, ultra-fast optical detectors, etc. [1]. Stimulated by the outcome of LT-GaAs and by the long-sought desire to obtain undoped, semi-insulating InP, an attempt has been made in recent years to grow off-stoichiometric InP at low temperatures by MBE. Though the material is highly P-rich, it is not at all semi-insulating. In fact, LT-InP becomes highly n-type conducting when the growth temperature is below normal temperatures [2–6]. In this paper, we shall first summarize our present understanding of the mechanism responsible for the n-type conductivity of LT-InP, which is attributed to the auto-ionization of PIn antisites preferentially introduced during the off-stoichiometric growth. We shall then review recent results from defect engineering, by utilizing the intrinsic n-type do-
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pants of the PIn antisites for modulation doping in InGaAs/InP heterostructures.
2. PIn-antisite engineered intrinsic doping in InP The first type of samples studied in this work is single InP epilayers, grown by gas source MBE, on semi-insulating Fe-doped InP substrate (Fig. 1a). The InP epilayers are typically 1-mm thick. The growth temperature varied from 480°C down to 265°C. Within this temperature range, the material remains good crystalline quality without a noticeable presence of structural defects. At growth temperatures below 200°C, on the other hand, the crystal quality of the material deteriorates drastically and eventually becomes polycrystalline. The free electron concentration of the LT-InP films was found to monotonically increase with decreasing
Fig. 3. Relative free electron concentration, measured via infrared absorption by free electrons, as a function of the hydrostatic pressure applied to the LT-InP epilayers. The observed carrier freeze-out occurs when the (0/ +) level of the PIn antisite enters the bandgap at a pressure higher than 1.7 GPa.
Fig. 1. Schematic pictures of the design for (a) the thick InP layers and (b) the InGaAs/InP heterostructures reviewed in this paper.
Fig. 2. Free electron concentration as a function of growth temperature in LT-InP epilayers. The insert shows the Fermi level pinning at the (0/+) level of the PIn antisite which leads to the saturation of the free electron concentration at growth temperatures lower than 265°C. C.B. denotes the conduction band.
growth temperatures [2,3] (Fig. 2). At 265°C a saturation electron concentration of 3×1018 cm − 3 was reached when the LT-InP exhibited a metallic n-type conduction. This saturation corresponds to a Fermi level pinning at the responsible donor level, i.e. the (0/+ ) level of the PIn antisite double donor to be discussed below. This donor level is resonant with the conduction band at Ec + 0.12 eV (see the inset in Fig. 2) estimated from the filling of the conduction band states [3]. The presence of this donor level was further confirmed by a freeze-out of the free electrons when the donor level enters the forbidden bandgap as a consequence of a widening bandgap upon the application of a hydrostatic pressure (Fig. 3). An analysis of the pressure coefficients gives an estimate of the resonant donor level at Ec + 0.11 eV, agreeing within the experimental error with the value of Ec + 0.12 eV given above [3]. The chemical nature of the donor giving rise to the n-type conductivity in LT-InP was revealed by the microscopically informative optically detected magnetic resonance (ODMR) technique. The ODMR studies have unambiguously shown the presence of the PIn antisite in the material, from its hyperfine fingerprint of the nuclear spin I=1/2 of the 31P atom with 100% natural abundance [3,5]. The presence of the PIn antisite was shown to correlate with a lowering of the growth temperature. If the PIn antisite is responsible for the n-type conductivity of LT-InP, the number of the ionized PIn antisites (P+ In ) should equal the free electron
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concentration. This close correlation as a function of growth temperature has, in fact, been provided experimentally as shown in Fig. 4. Here the concentration of P+ In can be scaled with the intensity of the magnetic circular dichroism (MCD) in absorption (MCDA) originating from the ground state of the singly ionized P+ In antisite. Such a direct correlation provides a convincing piece of evidence that the PIn antisite is indeed responsible for the n-type conductivity of LT-InP, via the auto-ionization of the (0/ + ) level resonant with the conduction band at Ec +0.12 eV. These experimental observations proved that the ntype conductivity of LT-InP is an intrinsic property of this off-stoichiometric material, and the PIn antisite is not at all suitable to make the material semi-insulating. On the contrary, it provides a new mechanism for n-type doping in InP, i.e. intrinsic doping without requiring an external doping source. In fact, the PIn
Fig. 4. Correlation between the free-electron concentration and the MCDA intensity (scaling with the concentration) of the P+ In antisite in LT-InP as a function of the growth temperature Tg.
Fig. 5. A schematic picture of the energy band diagram for the n-type modulation doped InGaAs/InP heterostructures. For simplicity only two subbands of the 2DEG are drawn. The filling of the 2DEG is depicted by the shaded region. The arrow across the InGaAs bandgap indicates the PL transition between the 2DEG and photo-excited holes upon optical excitation, shown in Fig. 7.
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antisite possesses properties that are regarded as basic requirements for a dopant, i.e. a known chemical identification: PIn; established electronic properties: the (0/+ ) level at Ec + 0.12 eV; a control of doping concentration: by varying Tg. Defect engineering by the PIn antisite can therefore provide a new doping mechanism by exploring such an intrinsic n-type dopant.
3. PIn-antisite engineered intrinsic modulation doping in InGaAs/InP heterostructures When the intrinsic doping concept is used in InPbased heterostructures, PIn-antisite engineered intrinsic modulation doping can readily be realized without invoking an external doping source. The success of such defect engineering was clearly demonstrated [7] by the first example cases of lattice-matched In0.53Ga0.47As/InP heterostructures (Fig. 1b) designed to resemble high electron mobility transistor (HEMT) structures. Three types of InGaAs/InP structures were grown and studied, all were grown by GS-MBE, on semi-insulating Fe-doped InP substrate: (1) intrinsically doped InGaAs/InP structures where the top InP layer was grown at a low temperature of 265°C (denoted as InGaAs/LT-InP); (2) extrinsically doped InGaAs/InP structures by shallow Si donors grown at normal temperature of 480°C; (3) undoped InGaAs/InP structures grown at normal high temperatures (HT) of 480°C (referred to as InGaAs/HT-InP or the reference sample). The design of these structures is identical except for the top InP layer and is shown in Fig. 1b. For the intrinsically doped InGaAs/InP, the entire structures were intentionally undoped and were grown at a normal temperature of 480°C except the top InP layer which was grown at 265°C. The reference sample was also grown with an identical structure except that the top InP layer was in this case grown at 480°C. In the extrinsically doped InGaAs/InP structures, the top InP was grown at normal temperature (480°C) but was intentionally doped by shallow Si donors. These structures were experimentally investigated by transport and optical techniques, mainly Shubnikov-de Haas (SdH) oscillations, Hall measurements, photoluminescence (PL), PL excitation (PLE) and magneto-optical spectroscopy. In Fig. 5 we first show a schematic energy band diagram of the modulation doped InGaAs/InP structures. Due to electron transfer from the n-type doped (either intrinsically or extrinsically) InP barrier to the InGaAs active layer, a notch potential is formed near the heterointerface where a two-dimensional electron gas (2DEG) is confined (Fig. 5).
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Fig. 6. SdH oscillation spectra taken at 1.5 K in dark from (a) the intrinsically doped InGaAs/LT-InP heterostructure and (b) the undoped reference sample, with the external magnetic field normal to the heterointerface plane. The total sheet concentrations of the 2DEG were determined from the analysis of the SdH oscillation spectra to be 1.15× 1012 and 2.37× 1011 cm − 2, respectively.
Fig. 7. PL spectra at 1.6 K obtained from the InGaAs/LT-InP structure, due to recombination of the 2DEG and photo-excited holes shown by the arrow in Fig. 5. The insert provides a close-up of the PL emission, at zero magnetic field (dashed curve) and 5 T (solid curve).
The first experimental evidence on the success of the intrinsic modulation doping concept was provided by the formation of a 2DEG in the InGaAs channel due to the intrinsic modulation doping, by a straightforward SdH oscillation study (Fig. 6a). From a detailed analysis of the SdH oscillations, two subbands of the 2DEG were shown to be readily occupied. The sheet concentrations of the first and second subbands of the 2D gas were determined to be n1 =6.75 ×1011 cm − 2 and n2 = 4.75× 1011 cm − 2, respectively, yielding a total sheet
concentration of 1.15×1012 cm − 2. This is to be compared with a much lower 2DEG concentration (2.37× 1011 cm − 2) observed in the undoped reference sample (Fig. 6b) taken under the same experimental conditions. Only one subband was partially occupied in this case, and was mainly induced by electron transfer from residual donors in the undoped InP buffer layer [8]. The 2D character of the 2DEG was revealed from an angular dependence study of the SdH oscillations, where the period of the SdH oscillations obeyed a cosine relation of the relative angle between the magnetic field and the growth direction [9]. Additional support for the success of the intrinsic modulation doping was provided from optical studies carried out on the same InGaAs/InP structures. In Fig. 7, a PL spectrum obtained from the intrinsically doped InGaAs/LT-InP structure is shown. The photon energy of the PL emission at around 0.8 eV is consistent with the bandgap of In0.53Ga0.47As. The PL spectrum was attributed to the recombination between the 2DEG and the photo-excited holes in the InGaAs layer, illustrated by an arrow across the InGaAs bandgap in Fig. 5. A similar PL emission has been observed in InGaAs/InP and InGaAs/AlInAs HEMT structures [10,11] where the conventional shallow donors were employed to achieve n-type modulation doping. The Landau level splitting of the PL emission in the magnetic field, shown in the insert of Fig. 7, and the deduced electron mass value m *: 0.05m0 further confirmed that the e recombination involves the 2DEG in the In0.53Ga0.47As channel [12], differing from the electron mass value 0.07m0 in InP. The 2D nature can again be shown m *: e by an angular dependence study of the Landau level splitting. Optically detected quantum oscillations (ODQO) [13] identical to the SdH oscillations were observed by detecting this PL emission at the highest photon energy near the Fermi edge, proving that the PL emission originates from the same 2DEG which give rise to the electrical SdH oscillations. The electron mobility of the 2DEG determined from the Hall measurements is me : 1.3× 104 cm2 V − 1 s − 1, which is much higher than me : 5× 102 cm2 V − 1 s − 1 determined independently for a single LT-InP layer grown at the same temperature (i.e. 265°C) [3]. This again confirmed that the 2DEG is located in the InGaAs channel due to electron transfer from the LT-InP doping region, and a much improved electron mobility is obtained due to a space separation of electrons from their parent donors.
4. Important issues in defect engineering by intrinsic doping Next, we shall review recent results from studies of a number of important issues such as doping efficiency,
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mobility and thermal stability, which are relevant and in fact determine the value of such defect engineering in practical device applications. A close comparison with the properties of the extrinsically modulation doped structures has been made, and the results are summarized in Table 1.
4.1. Doping efficiency As discussed above and also given in Table 1, a dense 2DEG is formed in the intrinsically modulation doped structures, with the total sheet concentration of ns = 1.15× 1012 cm − 2. This value is noticeably higher than ns =8.4×1011 cm − 2 in the extrinsically modulation doped structure with a Si-doping concentration of 1 × 1018 cm − 3, but comparable to that with a Si-doping concentration of 3× 1018 cm − 3 (ns =1.28× 1012 cm − 2). This agrees with a similar concentration of the PIn-antisite in the LT-InP layer grown at 265°C, determined in earlier studies of 1-mm-thick LT-InP thin films [2,3]. Our earlier self-consistent theoretical calculations of subband populations in the intrinsically modulation doped structure predicted the 2DEG concentration to be about a factor of two higher than that in the Si-doped structure. The present work seems to indicate that this earlier prediction has underestimated the 2DEG concentration in the extrinsically doped structures due to the assumption that the Fermi level is pinned at the donor level in thermal equilibrium. A much lower 2DEG concentration of 2.37× 1011 cm − 2 is formed in the reference structure as expected. The formation of the 2DEG in this nominally undoped structure is caused by an n-type background doping. The corresponding residual donor concentration is estimated to be about 3 ×1016 cm − 3 from our self-consistent theoretical calculations of the reference sample [8].
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It should be noted that the Hall concentration is markedly higher than the 2DEG concentration determined by the SdH oscillations in the intrinsically doped structure (Table 1). Such a deviation suggests the presence of a strong parallel conduction, due to the fact that a parallel conduction of lower mobility contributes much less to the SdH oscillation than to the Hall effect measurements. Possible parallel conduction channels available in the structure are the LT-InP doped region and the InP buffer layer. The SI InP substrate is not expected to contribute to electrical conduction. The contribution from the InP buffer layer in parallel conduction is due to the presence of a residual donor concentration of 3× 1016 cm − 3 [8]. The parallel conduction provided by the LT-InP layer can be attributed to the remaining 90% of the free electrons resulting from the autoionization of the PIn-antisite, after about 10% of them have transferred to the InGaAs channel.
4.2. Electron mobility The strong parallel conduction by the LT-InP layer of a lower mobility, independently determined to be about 500 cm2 V − 1 s − 1 in earlier studies [3], inevitably leads to a lower total Hall mobility as shown in Table 1. A reliable value for the 2DEG mobility requires a detailed analysis of the parallel conduction, by nHmH = Snimi and nHm 2H = Snim 2i , taking into account all the possible parallel conduction channels [14]. Here nH (or mH) and ni (or mi) represent the electron densities (or mobilities) of the total and individual conduction channels, respectively. Apart from the parallel channels by the LT-InP top layer and the buffer layer, the 2DEG by itself consists of two parallel conduction channels due to the occupation of two subbands. Though the concentration of the PIn-antisite and the electron mobility in
Table 1 The 2DEG concentrations (n) and mobilities (m) from the intrinsically and the extrinsically doped InP/InGaAs structures reviewed in this worka Parameters
From SdH
From Hall
Intrinsically doped structure
Extrinsically doped structure ([Si]: 1×1018 cm−3)
n1 (cm−2) m1 (cm2 V−1 s−1) n2 (cm−2) m2 (cm2 V−1 s−1) ntotal (cm−2)
6.75×1011 1.55×104 4.75×1011 3.05×104 1.15×1012
6.2×1011 7.0×103 2.2×1011 1.95×104 0.84×1012
nH (cm−2) mH (cm2 V−1 s−1)
5.7×1012 1.3×104 (3.1×104)b
0.8×1012 2.4×104
Extrinsically doped structure ([Si]: 3×1018 cm−3) 7.6×1011 5.2×1011 1.28×1012
a For the latter structures, two sets of samples were investigated with different doping concentrations of the shallow Si donor. Both the quantum mobilities m1 and m2 deduced from SdH and the Hall mobilities are given. The subscripts 1 and 2 denote the 1st and the 2nd subband of the 2DEG, respectively. b From Ref. [15].
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remote dopants is reduced in the intrinsically doped structure, which can be attributed to a different defect potential and a stronger screening of the defect potential by a degenerate electron gas present in the doped LT-InP barrier due to auto-ionization of the PIn antisite.
4.3. Thermal stability
Fig. 8. Optically detected quantum oscillation spectra taken at 1.5 K from the intrinsically doped InGaAs/LT-InP heterostructures, showing the effect of thermal annealing on the 2DEG concentration with or without surface protection.
the 1-mm-thick LT-InP films are known from earlier studies [3], they may be altered to some extent due to the increasing importance of surface states. Consequently, there are more unknown parameters than the number of the equations for parallel conduction that rules out a unique and reliable determination of the Hall mobility for each of the 2DEG subbands. Independent determination of some of the unknown parameters or a reduction in the number of parallel conduction channels is required to resolve the uncertainty. One approach of avoiding parallel conduction is to prevent the presence of free electrons in the LT-InP layer by applying hydrostatic pressure so that the (+ / 0) level of the PIn-antisite enters the bandgap and freezes electrons [15]. The Hall mobility deduced by this approach yields a value of 3.1×104 cm2 V − 1 s − 1, higher than 2.4× 104 cm2 V − 1 s − 1 deduced in this work from the extrinsically doped structure. In addition, the quantum mobility of the 2DEG deduced from an analysis of fast Fourier transformation of the SdH oscillations [16] is about a factor of 1.5 – 2 higher in the intrinsically doped structure compared to the extrinsically doped structure (Table 1). The difference in the Hall and quantum mobility between the intrinsically and extrinsically doped structures can only be due to the difference in scattering by remote dopants, since that is the only difference between the two structures. Scattering by alloy disorder, interface and background impurities and interband scattering should be nearly identical for both structures, and should not be responsible for the difference. This thus provides evidence that scattering of the 2DEG by
One of the other important issues which is relevant to practical device applications is the thermal stability of the intrinsic doping. This was addressed by examining the corresponding change in the 2DEG concentration upon annealing of the intrinsically modulation doped InGaAs/LT-InP heterostructures at 400–500°C. The annealing was performed both with and without a proper surface protection. For the surface protection, two methods have been used. The first method is proximity annealing, when the sample surface was protected by another piece of InP substrate. The second method is to overgrow an InP cap layer at 480°C on top of the LT-InP, which can serve both to anneal the underlying structure at 480°C and to protect the LTInP surface. A summary of the results is shown in Fig. 8. As can be seen, the annealing at 420°C for the intrinsically doped structure without surface protection caused a significant reduction (more than half) of the 2DEG concentration. With the surface protection, either by proximity annealing or overgrowth, the reduction of the 2DEG concentration has been markedly improved. This experimental finding demonstrated that the instability of the InP surface was the main reason for the thermal degradation of the structures. This is somewhat expected since an InP surface has been known to be unstable at T\ 400°C. It should be pointed out that the extrinsically modulation doped InGaAs/InP heterostructures have been found to exhibit a higher thermal stability, where the 2DEG concentration was virtually unaffected by annealing. This seems to indicate that the surface of the P-rich LT-InP is less stable than that of the Si-doped InP. We should point out, however, that even with the surface protection a noticeable decrease was observed in the 2DEG concentration after annealing of the intrinsically doped InGaAs/LT-InP heterostructures. Thus, it should be related to the thermal induced change in the bulk electrical properties of the LT-InP layer. This could be attributed either to a reduction in the number of the PIn antisites due to complex formation and/or precipitation or to a thermally induced formation/activation of compensating acceptors in the LT-InP layer, upon annealing. A decrease in conductivity of the thick LT grown InP epilayer after annealing has indeed been previously observed and has been attributed [17] to the thermal activation of acceptor impurities, which are passivated in the as-grown mate-
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rial by hydrogen, known to incorporate efficiently in the InP material during the LT MBE growth.
the U.S. Air Force Office of Scientific Research (AFOSR, F49620-93-1-0367TU).
5. Summary
References
A review is given on a successful example of defect engineering in InP-based heterostructures. Here an intrinsic defect, i.e. the PIn antisite, has been shown to possess the fundamental properties required for an n-type dopant, that is: (1) known chemical identification; (2) known electronic structure; (3) a control of doping concentration by varying growth temperature. Engineering of the PIn antisite in off-stoichiometrically grown InP at low growth temperatures has provided an intrinsic n-type doping source without invoking an external doping source. The success of the defect engineering has been clearly demonstrated by recent results from both the thick LT-InP epilayers and also the InGaAs/LT-InP heterostructures. Important physical properties of the intrinsic modulation doping have been discussed, such as doping efficiency, mobility and thermal stability, as compared to the extrinsically doped structures. Though it has only been demonstrated for InGaAs/InP heterostructures, the principle of defect engineering within the intrinsic doping concept is in fact rather general and can be extended to applications in other electronic material systems.
[1] G.L. Witt, R. Calawa, U. Mishra, E. Weber (Eds.), Low Temperature (LT) GaAs and Related Materials, MRS Symposium Proceedings, Vol. 241, Materials Research Society, Pittsburgh, PA, 1992. [2] B.W. Liang, P.Z. Lee, D.W. Shih, C.W. Tu, Appl. Phys. Lett. 60 (1992) 2014. [3] P. Dreszer, W.M. Chen, K. Seendripu, J.A. Wolk, W. Walukiewicz, B.W. Liang, C.W. Tu, E.R. Weber, Phys. Rev. B47 (1993) 4111. [4] W.M. Chen, P. Dreszer, A. Prasad, A. Kurpiewski, W. Walukiewicz, E.R. Weber, E. So¨rman, B. Monemar, B.W. Liang, C.W. Tu, J. Appl. Phys. 76 (1994) 600. [5] W.M. Chen, P. Dreszer, E.R. Weber, E. So¨rman, B. Monemar, B.W. Liang, C.W. Tu, J. Electron. Mater. 22 (1993) 1491. [6] P. Dreszer, W.M. Chen, D. Wasik, R. Leon, W. Wulukiewicz, B.W. Liang, C.W. Tu, E.R. Weber, J. Electron. Mater. 22 (1993) 1487. [7] W.M. Chen, I.A. Buyanova, A.V. Buyanov, T. Lundstro¨m, W.G. Bi, C.W. Tu, Phys. Rev. Lett. 77 (1996) 2734. [8] I.A. Buyanova, T. Lundstro¨m, A.V. Buyanov, W.M. Chen, W.G. Bi, C.W. Tu, Phys. Rev. B55 (1997) 7052. [9] W.M. Chen, I.A. Buyanova, W.G. Bi, C.W. Tu, Mater. Sci. Forum 258-263 (1997) 805. [10] D.G. Hayes, M.S. Skolnick, D.M. Whittaker, P.E. Simmonds, L.L. Taylor, S.J. Bass, L. Eaves, Phys. Rev. B44 (1991) 3436. [11] Y.-H. Zhang, D.-S. Jiang, R. Cingolani, K. Ploog, Appl. Phys. Lett. 56 (1990) 2195. [12] C.K. Sarkar, R.J. Nicholas, J.C. Portal, M. Razeghi, J. Chevrier, J. Massies, J. Phys. C18 (1985) 2667. [13] I.A. Buyanova, W.M. Chen, A. Buyanov, W.G. Bi, C.W. Tu, Appl. Phys. Lett. 69 (1996) 809. [14] R.L. Pertritz, Phys. Rev. 110 (1958) 1254. [15] D. Wasik, L. Dmowski, J. Mikucki, J. Lusakowski, L. Hsu, W. Walukiewicz, W.G. Bi, C.W. Tu, Mater. Sci. Forum 258-263 (1997) 813. [16] E. Skuras, R. Kumar, R.L. Williams, R.A. Stradling, J.E. Dmochowski, E.A. Johnson, A. Mackinnon, J.J. Harris, R.B. Beall, C. Skierbeszewski, J. Singleton, P.J. van der Wel, P. Wisniewski, Semicond. Sci. Technol. 6 (1991) 535. [17] B.W. Liang, Y. He, C.W. Tu, in: G.L. Witt, R. Calawa, U. Mishara, E. Weber (Eds.), Low Temperature (LT) GaAs and Related Materials, MRS Symposia Proceedings 241, Materials Research Society, Pittsburgh, PA, 1992, p. 297.
Acknowledgements We are grateful to W.G. Bi, A.V. Buyanov, T. Lundstro¨m, P. Dreszer, E.R. Weber, W. Walukiewicz, B.W. Liang, B. Monemar, E. So¨rman, J.A. Wolk, D. Wasik, H.J. Sun and G.D. Watkins for valuable contributions in connection to this work. This work has been supported by the Swedish Council for Engineering Sciences (TFR), the Swedish Natural Science Research Council (NFR) and an AASERT grant from
.