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Optical characterization of III-nitrides
Materials Science and Engineering B93 (2002) 112 /122 www.elsevier.com/locate/mseb
Optical characterization of III-nitrides B. Monemar a,*, P.P. Paskov a, T. Paskova a, J.P. Bergman a, G. Pozina a, W.M. Chen a, P.N. Hai, I.A. Buyanova a, H. Amano b, I. Akasaki b b
a Material Science Division, Department of Physics and Measurement Technology, Linko¨ping University, S-581 83 Linko¨ping, Sweden Department of Materials Science and Engineering, High Tech Research Center, Meijo University, 1-501 Shiogamaguchi, Tempaku-ku, Nagoya 468, Japan
Abstract Recent developments in material properties of GaN and related heterostructure combinations are reviewed, with emphasis on optical data. We discuss recent polarized photoluminescence (PL) data on the free excitons in GaN, obtained from thick HVPE grown layers. The exchange splitting constant is found to be about 0.6 meV, a more accurate value than previous suggestions. The PL signatures of shallow donors and acceptors, i.e. the bound excitons, are discussed and tentatively identified. Intrinsic point defects are discussed in terms of stability and experimental signatures. Quantum well structures in the InGaN/GaN and GaN/ AlGaN systems are briefly discussed, with emphasis on localization of carriers and excitons. # 2002 Published by Elsevier Science B.V. Keywords: GaN; InGaN; AlGaN; Photoluminescence; Excitons; Defects; Donors; Acceptors
1. Introduction The III-nitride research field has seen a dramatic development in the last decade, mainly due to the earlier breakthroughs in crystal growth [1] and p-type doping [2]. This class of materials is now predicted to play a major role for the development of both optical and electronic devices in the future. These applied aspects have been discussed in recent reviews [3 /5], and will not be treated here. In this paper we will concentrate on the present status of the material properties, which still need to be dramatically improved to meet the requirements of future advanced devices. We shall discuss some basic properties of the bulk materials, with emphasis on GaN, since very little detailed work has been done on AlN and InN. The properties of dopants and defects will also be treated; indeed defects are still not under proper control in these materials. We will discuss the present understanding of bound exciton signatures for the most common donors and acceptors, and briefly summarize the situation on intrinsic point defects. Quantum structures will be discussed as well, they are the building
* Corresponding author. Tel.: 46-13-281-765; fax: 46-13-142337. E-mail address: [email protected] (B. Monemar).
blocks of many devices. We shall mention recent results on both InGaN/GaN and GaN/AlGaN multi-quantum well structures, with emphasis on the influence of polarization fields and localization of excitons or charge carriers on the observed optical spectra.
2. The material quality, a limitation for studies of the physical properties Most physical properties cannot be determined with a satisfactory accuracy unless a well-defined material with a sufficiently low defect density is available. For the case of GaN the present status of the material is inadequate for an accurate determination of many physical parameters. The most suitable bulk material for physical measurements is thick epilayers grown by hydride vapor phase epitaxy (HVPE), or thin homoepitaxial layers grown by metal organic vapor phase epitaxy (MOVPE) or molecular beam epitaxy (MBE). The homoepitaxial material is at the moment not readily available due to the lack of GaN substrates. For the details of the electronic structure the optical data are generally required, these are not very accurate if the spectroscopic line width is too large, however. A desired line width is about 0.1 meV, however, this is not generally obtained
0921-5107/02/$ - see front matter # 2002 Published by Elsevier Science B.V. PII: S 0 9 2 1 - 5 1 0 7 ( 0 2 ) 0 0 0 0 6 - 5
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Fig. 1. (a) Expected exchange splitting for the free excitons in GaN, including the allowed polarizations. (b) PL spectra for an 80 mm HVPE grown GaN layer for s-polarization (solid line) and p-polarization (dotted line).
in GaN layers available today, and a value about 1 meV is more typical in the best cases. This is due to a large dislocation density, 107 /108 cm 2 even in good thick HVPE grown GaN layers, and a high concentration of residual dopants or complex defects (often in the range 1017 cm 3). Such defects cause inhomogeneous strain fields that broaden spectroscopic lines for electronic transitions. Measurements of electrical parameters are also of limited accuracy when the defect density is too high. The quantum structures involving GaN, AlGaN or InGaN materials are strongly influenced by polarization fields, as discussed in more detail below. Interface roughness cause a strong spectroscopic broadening, as well as In segregation effects in the case of structures containing InGaN. The spectral broadening in optical data is due to localization of excitons or carriers in the related potential fluctuations, which are enhanced in case of strong polarization induced electric fields. These effects, which are strong for III-nitride structures, limit the spectroscopic accuracy for these quantum structures.
Fig. 2. (a) Sketch of the bound exciton recombination process for neutral donors and acceptors. (b) Photoluminescence spectrum of a thick HVPE GaN layer showing two donor BE lines and an acceptor BE. The spectrum is upshifted by the compressive strain about 6 meV compared with strain free GaN.
3. Free excitons in GaN There have been very scattered values reported for the free exciton binding energies in GaN, as recently reviewed [6]. The scatter in the data is probably due to measurements on highly strained layers, and in some cases misinterpretations of some features in the optical spectra. The more recent data on both reflectance and photoluminescence (PL) reported for homoepitaxial samples appear convincing. A value of 25 meV is reported for the A exciton, and very similar values for the B and C excitons [7,8]. The n /2 transitions are
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clearly resolved in these spectra, giving confidence that these results will be accurate for pure unstrained material with a precision of about 1 meV. For the internal structure of the excitons, caused by the electronhole exchange interaction (Fig. 1(a)), the situation is much less clear. Previous results in literature give a spread in values for the splitting between the G5 and the G6 components of the A exciton between 0.12 meV [9] and 2.9 meV [10]. These measurements have to be undertaken in very thick samples, due to the selection rules for these transitions (Fig. 1(a)), so that the light can be collected perpendicular to the c -axis, and the s (E /c) and p (E ½½c ) polarization geometries can be realized. In Fig. 1(b) we show our own data from an 80 mm thick GaN layer grown by HVPE. Measurements of PL spectra in both s (E /c) and p (E ½½c ) polarizations exhibit clear shifts between the polarized components. The shift is 0.53 meV for the A exciton, and 0.24 meV for the B exciton [11]. These data can be well described in terms of an exchange interaction constant g /0.58 meV, which is close to what has been previously estimated from reflectance data by Julier et al. [12]. Note that the observed splitting of the B exciton is very strain dependent, while the A-exciton splitting is only weakly dependent on the biaxial strain. The longitudinal-transverse polariton splitting has not been directly determined in these experiments, but values of about 1 meV for unstrained samples were recently reported.
4. Bound excitons in GaN The electronic states of the bound excitons (BEs) depend strongly on the semiconductor material, in particular the band structure. For dopants we distinguish the case of donors and acceptors, their BE structure is different. Fig. 2(a) shows schematically the electronic configuration of BEs in GaN, for the two different cases discussed in this work: neutral donors and neutral acceptors. (Excitons bound to charged defects are also possible, but rarely observed to be prominent in semiconductors.) A basic assumption in the description of the principal bound exciton states for neutral donors and acceptors is a dominant coupling of the like particles in the BE states [13,14], as indicated in Fig. 2(a). For a shallow neutral donor for example, the two electrons in the BE state are assumed to pair off into a two-electron state with zero spin. The additional bound hole is then assumed to be weakly bound in the net hole-attractive Coulomb potential set up by this bound two-electron aggregate (Fig. 2(a)).
4.1. Excitons bound at shallow donors We shall restrict ourselves to neutral shallow donors, the case normally observed in donor BE (DBE) spectra. Such DBE spectra are observed in PL from heteroepitaxial as well as homoepitaxial GaN samples. The photon energy region for such DBE spectra is about 3.470 /3.472 eV at 2 K for strain free GaN [7,15]. It is possible that different shallow donors contribute to the spectra observed in this region in different samples, but actually the detailed work to assign the characteristic DBE line for a particular shallow donor has so far not been done properly on a series of backdoped samples. HCVD grown thick layers are the best heteroepitaxial samples for spectroscopy, with an optical linewidth of about 1 meV or less (see Fig. 2(b)). The experience is that there are often two dominant DBE lines, as in Fig. 2(b), with a spectroscopic distance of about 1 meV. A natural suggestion is then that these are connected with what is commonly believed to be the two most dominant residual shallow donors in GaN, oxygen and silicon. We shall discuss the present data and provide a tentative assignment below. In homoepitaxial GaN samples grown by MOCVD on strain free bulk GaN substrates a better spectral linewidth has been achieved, down to 0.2 meV has been observed for the DBEs [7]. In this case one dominant line is observed at 3.4709 eV, and has in recent literature been assigned to Si donors [16]. Both Si and O are known contaminants in HVPE GaN, Si originates from the quarts lining of the growth chamber, the quartz is attacked by HCl during growth. O may originate from insufficient purity of the gases supplied at growth, but also from attack of the sapphire substrate in the startup phase of the growth. If we assume that the lower energy line is due to Si, in analogy with previous work on homoepitaxial material, the high energy line would then be O-related. This means that the shallow DBEs in GaN do not obey the so called Haynes rule, stating a linear relationship between the DBE binding energy and the corresponding donor binding energy. Although it has been previously suggested that this rule is obeyed for GaN [17], for most direct bandgap materials this rule is not obeyed. Important additional spectroscopic information is available from the two-electron satellites, where the DBE recombines leaving the remaining neutral donor electron in an excited n /2 state. This gives direct spectroscopic information about the binding energy of the donor electron, which can be extrapolated as :/4/3 times the distance between the two-electron replica and the principal DBE line, assuming ideal effective mass behavior. An example of such a spectrum is given in Fig. 3 for a 400 mm thick HVPE GaN layer grown on sapphire. The DBE peak at about 3.471 eV is presumably composed of two unresolved DBEs, like in Fig.
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Fig. 3. Near bandgap photoluminescence spectrum for a 400 mm GaN layer on sapphire, showing both donor- and acceptor-related bound exciton lines.
2(b). A clear replica is observed at about 3.448 meV, i.e. 23 meV below the DBE. This spectrum then gives a binding energy of about 29 meV for the dominant donor in that sample. Comparison with infrared absorption data for Si donors gives a good agreement for the binding energy [16]. These combined data allow a tentative identification of the 3.4709 eV DBE as due to the neutral silicon donor. It is likely that the often observed second DBE peak at 3.4718 eV in unstrained GaN [6] (upshifted due to strain in Fig. 2(b)) is then related to the O donor, but this remains to be proven. In fact a second weaker replica is observed in Fig. 3 art 3.445 eV, i.e. about 26 meV below the principal DBE peak. This could then be the two-electron replica of the O DBE, but it might also be a deeper acceptor-related BE [18]. This agrees approximately with IR absorption data for similar samples, where a binding energy of 34.5 meV for the O donor was deduced [17]. The magnetic behavior of the DBEs is in this simple model (Fig. 2(a)) expected to be dominated by the magnetic moment of the bound hole in the BE state, which for a shallow donor BE should be effective mass like and related to the properties of free G9 holes at the valence band top, with an anisotropic g-tensor typical of a J /3/2 hole [19]. The magnetic properties of the silicon donor BE have recently been studied [20,21], and found to agree with this model, confirming the G9 character of the hole in the silicon-related DBE. Another important property of the DBEs is the recombination dynamics. The low temperature decay rate for the DBE PL line should give the radiative lifetime tBE of the DBE, which may also be theoretically predicted by the theory of Rashba et al. [22]. In the simplified form given by Henry et al. [23] tBE /constant l 2/nfBE, where l is the photon wavelength, n the refractive index, and fBE the BE oscillator strength.
Fig. 4. (a) Photoluminescence transients for a 400 mm thick GaN layer (a) and for a thin strained layer (b) both grown on sapphire.
For thick HVPE GaN layers /100 mm) of low dislocation density (in the 107 cm 2 range) a low intensity decay time about 200 ps is observed at 2 K (Fig. 4(a)), which is believed to be a typical radiative lifetime for DBEs related to shallow neutral donors. This corresponds to a value of fBE as about 10, very similar to the case of shallow donors in CdS [23]. In thin strained heteroepitaxial layers with a higher defect density the observed lifetime is typically shorter [24,25] (Fig. 4(b)), indicating an influence of excitation transfer from the DBE to lower energy states before the radiative recombination takes place. A shorter value of the decay time is also often observed in thick GaN layers on sapphire, if the point defect density is not sufficiently low [26], explained by the same argument. Similarly,
Fig. 5. The recombination lifetimes of the free exciton and of the donor bound exciton are shown vs. the energies of the free A exciton. The upper axis shows in-plane stress calculated for the GaN layers.
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homoepitaxial GaN layers also often show a fast DBE decay at low excitation density, due to excitation transfer to point defects [27]. A slow component is often observed in the DBE decay [27,28] indicating a slower radiative process overlapping the DBE transition (Fig. 4(b)). An interesting observation on the DBE decay time has recently been made in n-type layers grown on SiC substrates [24]. The DBE observed in these samples (presumably Si-related) shows a clear trend versus the energy position of the DBE line, which in turn correlated with the strain in the layer (Fig. 5). The biaxial strain in the GaN layer strongly affects the top valence band states, which in turn is reflected in the DBE wave function [29], affecting the oscillator strength and hence the observed radiative lifetime [24].
Fig. 6. (a) Low temperature PL spectrum (solid line) of a HVPE GaN layer shown in semi-logarithmic scale. The reflectivity spectrum (dashed) of the same sample taken at normal incidence geometry is plotted in linear scale; (b) Temperature dependence of the recombination lifetimes of the FXA, DBE, ABE1 and ABE2 transitions measured at their spectral peak position.
4.2. Excitons bound at acceptors The situation about acceptor BEs (ABEs) in GaN is somewhat less clear than for the donors. The most prominent neutral acceptor BE is found at about 3.466 eV in strain free GaN (Fig. 2(b), Fig. 3, Fig. 6(a)). There are several arguments that this should be an ABE for a neutral acceptor, one argument involves the magnetooptical properties which will be discussed below. Further this BE has been found to be dominant in slightly Mg-doped GaN samples [30 /32], it has a low energy acoustic phonon wing which is very characteristic for ABEs [33,34]. Also, it has a rather strong LO phonon coupling, much stronger than for the DBE, a property which has also been found in previous studies of shallow ABEs in CdS [33]. An alternative interpretation of this BE line as being related to charged donor bound excitons has recently been proposed by several authors [35 /38]. This idea conflicts with the observed behavior of the BE line in magnetooptical data [21], as further discussed below. As stated above many workers have tentatively associated this ABE with MgGa [28,30,39]. A difficulty is that most experiments with doped crystals up to now have been performed on heteroepitaxial material, where the ABE peak position strongly depends on the strain. While the most shallow BEs (such as the DBEs) tend to have a nearly constant distance to the FEA position in strained layers, the ABEs, being deeper, may not follow such a simple behavior. For example in Mg-doped samples with an A exciton position at 3.499 eV at 2 K the Mg ABE is observed at 3.480 eV [40], a binding energy of 19 meV. In unstrained samples the same distance is observed to be 11/12 meV [7,28,41], under the assumption that the 3.466 eV BE is Mg-related. The fact that this 3.466 eV ABE line is observed in almost all PL spectra of GaN in literature has then to be explained by the assumption that Mg is a common contaminant in GaN. Recent SIMS data on HVPE GaN have confirmed this situation, i. e. Mg is often present in HVPE GaN, at concentrations in the 1016 cm 3 range. It cannot be ruled out that the 3.466 eV BE is due to another commonly occurring but as yet unidentified acceptor, however. There are other deeper BE PL lines in GaN which are probably acceptor related. One example is given in Fig. 6(a), where additional ABE peaks are observed at 3.461 eV (a doublet structure) [26]. Since there is a strain shift upwards in the spectrum in Fig. 6(a), this is probably the same ABE as observed in some homoepitaxial layers at 3.455 eV [27], indicating a BE binding energy of about 21 meV. The observed splitting is tentatively related to a splitting in the neutral acceptor ground state, as supported by the lack of thermalization [26]. The identity of this acceptor is not established, but it is close to the position observed from the dominant ABE in Zn-
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doped samples, if strain shifts are considered [42]. Therefore, it might be due to residual Zn acceptors, present as contaminants in many samples. Neutral shallow acceptor bound excitons (ABEs) are expected to have a two-hole state derived from two G9 holes from the topmost valence band. Only one such state with J /0 is allowed by the Pauli principle (in contrast to the case of zinc blend semiconductors, where J /2 is possible as well). For deeper acceptors with a spin-like hole a similar J /0 two-hole state is expected. The additional electron in the ABE state then contributes its unpaired spin, so that the ABE state has J /1/ 2. The ABE state is then expected to have a nearly isotropic g-tensor, ideally reflecting the shallow donor g -value in GaN, g /1.95 [43]. A single J /1/2 ABE ground state is consequently expected in this picture. The recent magneto-optical data for homoepitaxial GaN seem to confirm the identification of the 3.466 eV BE as an ABE [21], as already mentioned above. The ABE state has an effective spin of the electron (Fig. 2(a)), while the acceptor ground state has the effective hole spin. We note that from independent magnetic resonance studies of the Mg acceptor it has been established that the Mg acceptor hole is essentially spin-like, with a g-factor close to 2 and nearly isotropic [44,45]. This may be expected even for a degenerate bound hole state with a very localized wave function. The magnetic field splitting of the ABE PL line should, therefore, show essentially an isotropic splitting pattern into three lines, as schematically indicated in Fig. 7. The experimental data from Ref. 21 may be interpreted along these lines. For the case of B /c a splitting into three lines is clearly seen, in agreement with the expected pattern in Fig. 7. (The interpretation in Ref. 21 is different, since the authors assume a G9 hole for the acceptor). For the alternative interpretation of a charged donor BE [35 /38] there is no particle in the final state of the PL transition. Therefore, the magnetic splitting should be similar to the case of the free excitons
Fig. 7. Sketch of the expected photoluminescence lines of a spin-like acceptor in a magnetic field. In this case both the ABE state and the acceptor ground state will have an approximately isotropic splitting corresponding to g : 2, which leads to essentially a three line structure of the PL spectrum (right).
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[8]. For instance, in the case B /c a splitting into a doublet would be expected, which is not observed. The ABE recombination dynamics has been studied in transient PL data [25 /28]. The decay curves are usually clean exponential for the ABEs, Fig. 4(b), (in contrast to the case of DBEs), and presumable reflect the radiative lifetimes of the BEs at the lowest temperatures, before thermalization sets in (Fig. 6(b)) [26]. The observed values of the radiative lifetimes are about 0.7 ns for the 3.466 eV ABE [46] and much longer, 3.6 ns, for the deeper 3.455 /3.46 eV ABE (Fig. 6(b)) [26]. This corresponds to an oscillator strength of the order 1, very similar to the shallow acceptor BEs in CdS [23].
5. Intrinsic point defects in GaN By point defects we here understand the isolated intrinsic vacancies and interstitials of the material, i.e. VGa and VN, as well as Gai and Ni . Theoretically the equilibrium formation enthalpies of these defects have been estimated, and it has been concluded that the vacancies are expected to form at suitable Fermi level positions, while the interstitials or antisites are less likely to be abundant from growth [47]. Any of these defects will be possible to produce via ionizing particle irradiation, however. The key question is then the stability of these defects, which is difficult to predict from theory. A general trend among different materials is that the hard wide bandgap materials exhibit a higher stability for the point defects than do the lower bandgap softer materials. So e.g. it is well known that for Si both the vacancy and the interstitial become unstable already at cryogenic temperatures, and are therefore, only expected as parts in complex defects after growth and/or processing. SiC is a relevant material for comparison with GaN, SiC is hard and has higher growth and processing temperatures than GaN. For SiC the C vacancy is suggested to be stable up to 150 8C [48], while the Si vacancy survives up to 750 8C [49]. The stability of the interstitials in SiC is not known, but it is believed they may be less stable than the vacancies. In the light of this comparison some claims in literature about the observation of VGa and VN in asgrown GaN appear questionable. In the case of VGa positron annihilation experiments have been interpreted as evidence for the presence of Ga vacancies in material cooled down from a growth temperature above 1000 8C [50]. It is probable that complexes involving VGa are detected in these experiments, since VGa is not expected to be stable at the growth temperature (1050 8C or even higher). Recent studies of positron annihilation in HVPE GaN reveals a rather high concentration of VGa complexes in nominally undoped material (in the 1017 cm 2 range) [51], and therefore, such defects may well influence the electrical properties and compensation
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Fig. 8. (a) The ODMR spectrum observed in the as-grown Zn-doped GaN sample, at 9.21686 GHz and 5 K. The known ODMR signals from a shallow donor and Zn are denoted by SD and Zn, respectively. (b) The Ga-1 ODMR spectrum after subtracting the other known ODMR signals. (c) The simulated Ga-1 ODMR spectrum taking into account the contribution from the two isotopes of Ga.
in the material. In similar experiments a strong increase of VGa complexes has been observed in Be doped MBE GaN [52]. This is an indication of the introduction of donors during the growth of Be doped GaN, increasing the solubility of VGa. Be is a shallow acceptor in GaN (binding energy about 100 meV), but it has so far not been possible to achieve a sufficiently high p-doping level, indicating a complex behavior of Be in GaN. Similarly, the VN has long been blamed for the ndoping frequently observed in undoped GaN films [53]. In this case theory predicts a shallow donor state for the isolated VN defect [47,54], but the stability at the growth temperature is in doubt. It has been claimed that the VN donor is observed in electrical measurements at room temperature in electron irradiated GaN [55]. In these cases it has not yet been proven that these defects are not complexes involving the VN as one component, however [56]. Recent studies on implantation induced defects indicate a substantial mobility of the intrinsic point defects at room temperature [57]. So far no conclusive magnetic resonance data are available for the vacancy defects in GaN. The Ga interstitial has recently been studied by ODMR in irradiated GaN [58]. It has been found to be unstable already below room temperature, unlike earlier suggestions that the isolated Gai can actually be present in GaN at room temperature [59]. Complexes involving Gai are of course possible as recently reported
Fig. 9. (a) Sketch of the In0.11Ga0.89N/GaN MQW structure grown with mass transport MOVPE. (b) Spatially integrated PL spectrum at 2 K at a low excitation density of 10 mW cm 2.
[60], another example of this has recently been observed by ODMR in Zn-doped GaN [61], see Fig. 8. A Garelated hyperfine structure is clearly seen in Fig. 8, demonstrating together with a trigonal symmetry a defect likely to be related to a Gai complex. Whether the other component in the complex is Zn or another impurity is not clear at the moment. The N interstitial has so far not been positively identified, although it was suggested to be observed in electron irradiated material at room temperature [55].
6. Quantum well structures The III-nitride system offers several interesting material combinations for heterostructures, with a type I band alignment. The InGaN/GaN system has been studied intensely, due to the interest for visible LEDs and lasers [4]. The AlGaN/GaN system is also of great interest, both for HEMT transistor structures and for UV LEDs. We shall first discuss some properties of the InGaN/GaN system.
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Fig. 10. (a) Set of micro-PL spectra of InGaN/GaN MQW structure. (b) PL integrated intensity as a function of the spot position.
6.1. InGaN/GaN QWs One question of prime importance in this materials system is the presence of In segregation during growth of InGaN at elevated temperatures. Theoretically it has been predicted that for compositions above about 10% In the alloy has a strong tendency to segregate so that a second phase rich in In occurs [62]. The presence of such an In rich phase has also been claimed in optical data for In compositions about 10% [63], via a second low energy peak in PL spectra. It appears that this phenomenon depends on the growth conditions, it may be possible to avoid in many cases. In Fig. 9(b) we show PL data from an InGaN/GaN MQW with an In composition in the QWs of 11% (Fig. 9(a)), where the entire structure was
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Fig. 11. (a) Schematic drawing of an AlGaN/GaN MQW LED structure. (b) Panchromatic CL image of the n-AlGaN buffer layer in the structure shown in (a).
grown at one temperature (800 8C). The observed linewidth of the QW excitons is just about 40 meV at 2 K, which is what may be expected from the added contributions to the linewidth of the alloy broadening (about 15 meV) and the interface roughness broadening (about 25 meV). No low energy peak with the properties expected for In rich islands is observed, indicating the absence of segregation. Also, the dislocation density does not seem to influence this linewidth. In Fig. 10(a) is shown a set of micro-PL spectra for this structure, which has stripes of mass transport overgrown areas in the GaN buffer layer, reducing the dislocation density by about an order of magnitude down to the 107 cm 2 range. A scan across both the overgrown and the standard area shows no difference in
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Fig. 12. (a /d) PL decay curves measured at 2 K for different AlGaN/GaN MQWs with different doping. The Si concentration in the barriers is indicated in the figure. The evaluated PL lifetime is about 360 ps for all MQWs.
PL linewidth, only a small difference in PL intensity (Fig. 10(b)). The reason for different results on InGaN/GaN MQW properties between samples grown in different laboratories is not clarified, but is certainly to be found in differences in the growth conditions. We speculate that the procedure of a temperature ramping between the growth of GaN and InGaN layers, respectively, employed by many workers, might have adverse effects on the quality of the InGaN layers, since the surface is exposed to an excessive temperature during the ramping after the growth of the InGaN layer. Evidence for the absence of In segregation even for higher In compositions has recently been reported for MBE grown layers [64]. Such a segregation may even be beneficial in the case of LEDs, since it allows efficient radiative emission from localized excitons at room temperature. For lasers, however, phase separation is expected to reduce the gain, and should be avoided. 6.2. AlGaN/GaN MQWs AlGaN/GaN MQWs are different, since the well material is a binary. The localization of excitons or carriers in this case is expected to be due to interface roughness and the alloy fluctuations in the AlGaN barrier. We have studied such MQW structures grown by MOVPE, where lateral overgrowth was applied in the AlGaN buffer layer (see Fig. 11(a)). CL topographs of the AlGaN buffer layer shows well resolved contrast
from threading dislocations of different types (Fig. 11(b)). The edge dislocations in the coalescence zone show only a weak contrast, in agreement to previous reports [65]. The PL linewidth of the MQW excitons in this case is about 40 meV, i.e. about a factor two larger than the best reported values in MBE samples. This is already a sign of strong localization potentials. Since this MQW was present at the surface of the structure we cannot rule out a broadening due to a potential gradient across the MQW part, so that all QWs are not equivalent. A remarkable result is obtained in a set of MQW (5 QWs) samples with different Si doping, from nominally undoped up to the 1020 cm 2 range. Ideally such a set would be expected to exhibit a strong variation on radiative lifetime t with doping, i.e. t should vary as 1/n where n is the electron density. Instead no variation of the decay time is observed, as shown in Fig. 12. This is most probably explained as due to a strong localization of the holes at all doping levels. Such localization is expected to be particularly strong for rather low Al composition in the AlGaN barriers, as recently suggested [66].
Acknowledgements This work was partially supported by the European Community via the CLERMONT project PHRN-CT1999-00132.
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Optical characterization of III-nitrides B. Monemar a,*, P.P. Paskov a, T. Paskova a, J.P. Bergman a, G. Pozina a, W.M. Chen a, P.N. Hai, I.A. Buyanova a, H. Amano b, I. Akasaki b b
a Material Science Division, Department of Physics and Measurement Technology, Linko¨ping University, S-581 83 Linko¨ping, Sweden Department of Materials Science and Engineering, High Tech Research Center, Meijo University, 1-501 Shiogamaguchi, Tempaku-ku, Nagoya 468, Japan
Abstract Recent developments in material properties of GaN and related heterostructure combinations are reviewed, with emphasis on optical data. We discuss recent polarized photoluminescence (PL) data on the free excitons in GaN, obtained from thick HVPE grown layers. The exchange splitting constant is found to be about 0.6 meV, a more accurate value than previous suggestions. The PL signatures of shallow donors and acceptors, i.e. the bound excitons, are discussed and tentatively identified. Intrinsic point defects are discussed in terms of stability and experimental signatures. Quantum well structures in the InGaN/GaN and GaN/ AlGaN systems are briefly discussed, with emphasis on localization of carriers and excitons. # 2002 Published by Elsevier Science B.V. Keywords: GaN; InGaN; AlGaN; Photoluminescence; Excitons; Defects; Donors; Acceptors
1. Introduction The III-nitride research field has seen a dramatic development in the last decade, mainly due to the earlier breakthroughs in crystal growth [1] and p-type doping [2]. This class of materials is now predicted to play a major role for the development of both optical and electronic devices in the future. These applied aspects have been discussed in recent reviews [3 /5], and will not be treated here. In this paper we will concentrate on the present status of the material properties, which still need to be dramatically improved to meet the requirements of future advanced devices. We shall discuss some basic properties of the bulk materials, with emphasis on GaN, since very little detailed work has been done on AlN and InN. The properties of dopants and defects will also be treated; indeed defects are still not under proper control in these materials. We will discuss the present understanding of bound exciton signatures for the most common donors and acceptors, and briefly summarize the situation on intrinsic point defects. Quantum structures will be discussed as well, they are the building
* Corresponding author. Tel.: 46-13-281-765; fax: 46-13-142337. E-mail address: [email protected] (B. Monemar).
blocks of many devices. We shall mention recent results on both InGaN/GaN and GaN/AlGaN multi-quantum well structures, with emphasis on the influence of polarization fields and localization of excitons or charge carriers on the observed optical spectra.
2. The material quality, a limitation for studies of the physical properties Most physical properties cannot be determined with a satisfactory accuracy unless a well-defined material with a sufficiently low defect density is available. For the case of GaN the present status of the material is inadequate for an accurate determination of many physical parameters. The most suitable bulk material for physical measurements is thick epilayers grown by hydride vapor phase epitaxy (HVPE), or thin homoepitaxial layers grown by metal organic vapor phase epitaxy (MOVPE) or molecular beam epitaxy (MBE). The homoepitaxial material is at the moment not readily available due to the lack of GaN substrates. For the details of the electronic structure the optical data are generally required, these are not very accurate if the spectroscopic line width is too large, however. A desired line width is about 0.1 meV, however, this is not generally obtained
0921-5107/02/$ - see front matter # 2002 Published by Elsevier Science B.V. PII: S 0 9 2 1 - 5 1 0 7 ( 0 2 ) 0 0 0 0 6 - 5
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Fig. 1. (a) Expected exchange splitting for the free excitons in GaN, including the allowed polarizations. (b) PL spectra for an 80 mm HVPE grown GaN layer for s-polarization (solid line) and p-polarization (dotted line).
in GaN layers available today, and a value about 1 meV is more typical in the best cases. This is due to a large dislocation density, 107 /108 cm 2 even in good thick HVPE grown GaN layers, and a high concentration of residual dopants or complex defects (often in the range 1017 cm 3). Such defects cause inhomogeneous strain fields that broaden spectroscopic lines for electronic transitions. Measurements of electrical parameters are also of limited accuracy when the defect density is too high. The quantum structures involving GaN, AlGaN or InGaN materials are strongly influenced by polarization fields, as discussed in more detail below. Interface roughness cause a strong spectroscopic broadening, as well as In segregation effects in the case of structures containing InGaN. The spectral broadening in optical data is due to localization of excitons or carriers in the related potential fluctuations, which are enhanced in case of strong polarization induced electric fields. These effects, which are strong for III-nitride structures, limit the spectroscopic accuracy for these quantum structures.
Fig. 2. (a) Sketch of the bound exciton recombination process for neutral donors and acceptors. (b) Photoluminescence spectrum of a thick HVPE GaN layer showing two donor BE lines and an acceptor BE. The spectrum is upshifted by the compressive strain about 6 meV compared with strain free GaN.
3. Free excitons in GaN There have been very scattered values reported for the free exciton binding energies in GaN, as recently reviewed [6]. The scatter in the data is probably due to measurements on highly strained layers, and in some cases misinterpretations of some features in the optical spectra. The more recent data on both reflectance and photoluminescence (PL) reported for homoepitaxial samples appear convincing. A value of 25 meV is reported for the A exciton, and very similar values for the B and C excitons [7,8]. The n /2 transitions are
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clearly resolved in these spectra, giving confidence that these results will be accurate for pure unstrained material with a precision of about 1 meV. For the internal structure of the excitons, caused by the electronhole exchange interaction (Fig. 1(a)), the situation is much less clear. Previous results in literature give a spread in values for the splitting between the G5 and the G6 components of the A exciton between 0.12 meV [9] and 2.9 meV [10]. These measurements have to be undertaken in very thick samples, due to the selection rules for these transitions (Fig. 1(a)), so that the light can be collected perpendicular to the c -axis, and the s (E /c) and p (E ½½c ) polarization geometries can be realized. In Fig. 1(b) we show our own data from an 80 mm thick GaN layer grown by HVPE. Measurements of PL spectra in both s (E /c) and p (E ½½c ) polarizations exhibit clear shifts between the polarized components. The shift is 0.53 meV for the A exciton, and 0.24 meV for the B exciton [11]. These data can be well described in terms of an exchange interaction constant g /0.58 meV, which is close to what has been previously estimated from reflectance data by Julier et al. [12]. Note that the observed splitting of the B exciton is very strain dependent, while the A-exciton splitting is only weakly dependent on the biaxial strain. The longitudinal-transverse polariton splitting has not been directly determined in these experiments, but values of about 1 meV for unstrained samples were recently reported.
4. Bound excitons in GaN The electronic states of the bound excitons (BEs) depend strongly on the semiconductor material, in particular the band structure. For dopants we distinguish the case of donors and acceptors, their BE structure is different. Fig. 2(a) shows schematically the electronic configuration of BEs in GaN, for the two different cases discussed in this work: neutral donors and neutral acceptors. (Excitons bound to charged defects are also possible, but rarely observed to be prominent in semiconductors.) A basic assumption in the description of the principal bound exciton states for neutral donors and acceptors is a dominant coupling of the like particles in the BE states [13,14], as indicated in Fig. 2(a). For a shallow neutral donor for example, the two electrons in the BE state are assumed to pair off into a two-electron state with zero spin. The additional bound hole is then assumed to be weakly bound in the net hole-attractive Coulomb potential set up by this bound two-electron aggregate (Fig. 2(a)).
4.1. Excitons bound at shallow donors We shall restrict ourselves to neutral shallow donors, the case normally observed in donor BE (DBE) spectra. Such DBE spectra are observed in PL from heteroepitaxial as well as homoepitaxial GaN samples. The photon energy region for such DBE spectra is about 3.470 /3.472 eV at 2 K for strain free GaN [7,15]. It is possible that different shallow donors contribute to the spectra observed in this region in different samples, but actually the detailed work to assign the characteristic DBE line for a particular shallow donor has so far not been done properly on a series of backdoped samples. HCVD grown thick layers are the best heteroepitaxial samples for spectroscopy, with an optical linewidth of about 1 meV or less (see Fig. 2(b)). The experience is that there are often two dominant DBE lines, as in Fig. 2(b), with a spectroscopic distance of about 1 meV. A natural suggestion is then that these are connected with what is commonly believed to be the two most dominant residual shallow donors in GaN, oxygen and silicon. We shall discuss the present data and provide a tentative assignment below. In homoepitaxial GaN samples grown by MOCVD on strain free bulk GaN substrates a better spectral linewidth has been achieved, down to 0.2 meV has been observed for the DBEs [7]. In this case one dominant line is observed at 3.4709 eV, and has in recent literature been assigned to Si donors [16]. Both Si and O are known contaminants in HVPE GaN, Si originates from the quarts lining of the growth chamber, the quartz is attacked by HCl during growth. O may originate from insufficient purity of the gases supplied at growth, but also from attack of the sapphire substrate in the startup phase of the growth. If we assume that the lower energy line is due to Si, in analogy with previous work on homoepitaxial material, the high energy line would then be O-related. This means that the shallow DBEs in GaN do not obey the so called Haynes rule, stating a linear relationship between the DBE binding energy and the corresponding donor binding energy. Although it has been previously suggested that this rule is obeyed for GaN [17], for most direct bandgap materials this rule is not obeyed. Important additional spectroscopic information is available from the two-electron satellites, where the DBE recombines leaving the remaining neutral donor electron in an excited n /2 state. This gives direct spectroscopic information about the binding energy of the donor electron, which can be extrapolated as :/4/3 times the distance between the two-electron replica and the principal DBE line, assuming ideal effective mass behavior. An example of such a spectrum is given in Fig. 3 for a 400 mm thick HVPE GaN layer grown on sapphire. The DBE peak at about 3.471 eV is presumably composed of two unresolved DBEs, like in Fig.
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Fig. 3. Near bandgap photoluminescence spectrum for a 400 mm GaN layer on sapphire, showing both donor- and acceptor-related bound exciton lines.
2(b). A clear replica is observed at about 3.448 meV, i.e. 23 meV below the DBE. This spectrum then gives a binding energy of about 29 meV for the dominant donor in that sample. Comparison with infrared absorption data for Si donors gives a good agreement for the binding energy [16]. These combined data allow a tentative identification of the 3.4709 eV DBE as due to the neutral silicon donor. It is likely that the often observed second DBE peak at 3.4718 eV in unstrained GaN [6] (upshifted due to strain in Fig. 2(b)) is then related to the O donor, but this remains to be proven. In fact a second weaker replica is observed in Fig. 3 art 3.445 eV, i.e. about 26 meV below the principal DBE peak. This could then be the two-electron replica of the O DBE, but it might also be a deeper acceptor-related BE [18]. This agrees approximately with IR absorption data for similar samples, where a binding energy of 34.5 meV for the O donor was deduced [17]. The magnetic behavior of the DBEs is in this simple model (Fig. 2(a)) expected to be dominated by the magnetic moment of the bound hole in the BE state, which for a shallow donor BE should be effective mass like and related to the properties of free G9 holes at the valence band top, with an anisotropic g-tensor typical of a J /3/2 hole [19]. The magnetic properties of the silicon donor BE have recently been studied [20,21], and found to agree with this model, confirming the G9 character of the hole in the silicon-related DBE. Another important property of the DBEs is the recombination dynamics. The low temperature decay rate for the DBE PL line should give the radiative lifetime tBE of the DBE, which may also be theoretically predicted by the theory of Rashba et al. [22]. In the simplified form given by Henry et al. [23] tBE /constant l 2/nfBE, where l is the photon wavelength, n the refractive index, and fBE the BE oscillator strength.
Fig. 4. (a) Photoluminescence transients for a 400 mm thick GaN layer (a) and for a thin strained layer (b) both grown on sapphire.
For thick HVPE GaN layers /100 mm) of low dislocation density (in the 107 cm 2 range) a low intensity decay time about 200 ps is observed at 2 K (Fig. 4(a)), which is believed to be a typical radiative lifetime for DBEs related to shallow neutral donors. This corresponds to a value of fBE as about 10, very similar to the case of shallow donors in CdS [23]. In thin strained heteroepitaxial layers with a higher defect density the observed lifetime is typically shorter [24,25] (Fig. 4(b)), indicating an influence of excitation transfer from the DBE to lower energy states before the radiative recombination takes place. A shorter value of the decay time is also often observed in thick GaN layers on sapphire, if the point defect density is not sufficiently low [26], explained by the same argument. Similarly,
Fig. 5. The recombination lifetimes of the free exciton and of the donor bound exciton are shown vs. the energies of the free A exciton. The upper axis shows in-plane stress calculated for the GaN layers.
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homoepitaxial GaN layers also often show a fast DBE decay at low excitation density, due to excitation transfer to point defects [27]. A slow component is often observed in the DBE decay [27,28] indicating a slower radiative process overlapping the DBE transition (Fig. 4(b)). An interesting observation on the DBE decay time has recently been made in n-type layers grown on SiC substrates [24]. The DBE observed in these samples (presumably Si-related) shows a clear trend versus the energy position of the DBE line, which in turn correlated with the strain in the layer (Fig. 5). The biaxial strain in the GaN layer strongly affects the top valence band states, which in turn is reflected in the DBE wave function [29], affecting the oscillator strength and hence the observed radiative lifetime [24].
Fig. 6. (a) Low temperature PL spectrum (solid line) of a HVPE GaN layer shown in semi-logarithmic scale. The reflectivity spectrum (dashed) of the same sample taken at normal incidence geometry is plotted in linear scale; (b) Temperature dependence of the recombination lifetimes of the FXA, DBE, ABE1 and ABE2 transitions measured at their spectral peak position.
4.2. Excitons bound at acceptors The situation about acceptor BEs (ABEs) in GaN is somewhat less clear than for the donors. The most prominent neutral acceptor BE is found at about 3.466 eV in strain free GaN (Fig. 2(b), Fig. 3, Fig. 6(a)). There are several arguments that this should be an ABE for a neutral acceptor, one argument involves the magnetooptical properties which will be discussed below. Further this BE has been found to be dominant in slightly Mg-doped GaN samples [30 /32], it has a low energy acoustic phonon wing which is very characteristic for ABEs [33,34]. Also, it has a rather strong LO phonon coupling, much stronger than for the DBE, a property which has also been found in previous studies of shallow ABEs in CdS [33]. An alternative interpretation of this BE line as being related to charged donor bound excitons has recently been proposed by several authors [35 /38]. This idea conflicts with the observed behavior of the BE line in magnetooptical data [21], as further discussed below. As stated above many workers have tentatively associated this ABE with MgGa [28,30,39]. A difficulty is that most experiments with doped crystals up to now have been performed on heteroepitaxial material, where the ABE peak position strongly depends on the strain. While the most shallow BEs (such as the DBEs) tend to have a nearly constant distance to the FEA position in strained layers, the ABEs, being deeper, may not follow such a simple behavior. For example in Mg-doped samples with an A exciton position at 3.499 eV at 2 K the Mg ABE is observed at 3.480 eV [40], a binding energy of 19 meV. In unstrained samples the same distance is observed to be 11/12 meV [7,28,41], under the assumption that the 3.466 eV BE is Mg-related. The fact that this 3.466 eV ABE line is observed in almost all PL spectra of GaN in literature has then to be explained by the assumption that Mg is a common contaminant in GaN. Recent SIMS data on HVPE GaN have confirmed this situation, i. e. Mg is often present in HVPE GaN, at concentrations in the 1016 cm 3 range. It cannot be ruled out that the 3.466 eV BE is due to another commonly occurring but as yet unidentified acceptor, however. There are other deeper BE PL lines in GaN which are probably acceptor related. One example is given in Fig. 6(a), where additional ABE peaks are observed at 3.461 eV (a doublet structure) [26]. Since there is a strain shift upwards in the spectrum in Fig. 6(a), this is probably the same ABE as observed in some homoepitaxial layers at 3.455 eV [27], indicating a BE binding energy of about 21 meV. The observed splitting is tentatively related to a splitting in the neutral acceptor ground state, as supported by the lack of thermalization [26]. The identity of this acceptor is not established, but it is close to the position observed from the dominant ABE in Zn-
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doped samples, if strain shifts are considered [42]. Therefore, it might be due to residual Zn acceptors, present as contaminants in many samples. Neutral shallow acceptor bound excitons (ABEs) are expected to have a two-hole state derived from two G9 holes from the topmost valence band. Only one such state with J /0 is allowed by the Pauli principle (in contrast to the case of zinc blend semiconductors, where J /2 is possible as well). For deeper acceptors with a spin-like hole a similar J /0 two-hole state is expected. The additional electron in the ABE state then contributes its unpaired spin, so that the ABE state has J /1/ 2. The ABE state is then expected to have a nearly isotropic g-tensor, ideally reflecting the shallow donor g -value in GaN, g /1.95 [43]. A single J /1/2 ABE ground state is consequently expected in this picture. The recent magneto-optical data for homoepitaxial GaN seem to confirm the identification of the 3.466 eV BE as an ABE [21], as already mentioned above. The ABE state has an effective spin of the electron (Fig. 2(a)), while the acceptor ground state has the effective hole spin. We note that from independent magnetic resonance studies of the Mg acceptor it has been established that the Mg acceptor hole is essentially spin-like, with a g-factor close to 2 and nearly isotropic [44,45]. This may be expected even for a degenerate bound hole state with a very localized wave function. The magnetic field splitting of the ABE PL line should, therefore, show essentially an isotropic splitting pattern into three lines, as schematically indicated in Fig. 7. The experimental data from Ref. 21 may be interpreted along these lines. For the case of B /c a splitting into three lines is clearly seen, in agreement with the expected pattern in Fig. 7. (The interpretation in Ref. 21 is different, since the authors assume a G9 hole for the acceptor). For the alternative interpretation of a charged donor BE [35 /38] there is no particle in the final state of the PL transition. Therefore, the magnetic splitting should be similar to the case of the free excitons
Fig. 7. Sketch of the expected photoluminescence lines of a spin-like acceptor in a magnetic field. In this case both the ABE state and the acceptor ground state will have an approximately isotropic splitting corresponding to g : 2, which leads to essentially a three line structure of the PL spectrum (right).
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[8]. For instance, in the case B /c a splitting into a doublet would be expected, which is not observed. The ABE recombination dynamics has been studied in transient PL data [25 /28]. The decay curves are usually clean exponential for the ABEs, Fig. 4(b), (in contrast to the case of DBEs), and presumable reflect the radiative lifetimes of the BEs at the lowest temperatures, before thermalization sets in (Fig. 6(b)) [26]. The observed values of the radiative lifetimes are about 0.7 ns for the 3.466 eV ABE [46] and much longer, 3.6 ns, for the deeper 3.455 /3.46 eV ABE (Fig. 6(b)) [26]. This corresponds to an oscillator strength of the order 1, very similar to the shallow acceptor BEs in CdS [23].
5. Intrinsic point defects in GaN By point defects we here understand the isolated intrinsic vacancies and interstitials of the material, i.e. VGa and VN, as well as Gai and Ni . Theoretically the equilibrium formation enthalpies of these defects have been estimated, and it has been concluded that the vacancies are expected to form at suitable Fermi level positions, while the interstitials or antisites are less likely to be abundant from growth [47]. Any of these defects will be possible to produce via ionizing particle irradiation, however. The key question is then the stability of these defects, which is difficult to predict from theory. A general trend among different materials is that the hard wide bandgap materials exhibit a higher stability for the point defects than do the lower bandgap softer materials. So e.g. it is well known that for Si both the vacancy and the interstitial become unstable already at cryogenic temperatures, and are therefore, only expected as parts in complex defects after growth and/or processing. SiC is a relevant material for comparison with GaN, SiC is hard and has higher growth and processing temperatures than GaN. For SiC the C vacancy is suggested to be stable up to 150 8C [48], while the Si vacancy survives up to 750 8C [49]. The stability of the interstitials in SiC is not known, but it is believed they may be less stable than the vacancies. In the light of this comparison some claims in literature about the observation of VGa and VN in asgrown GaN appear questionable. In the case of VGa positron annihilation experiments have been interpreted as evidence for the presence of Ga vacancies in material cooled down from a growth temperature above 1000 8C [50]. It is probable that complexes involving VGa are detected in these experiments, since VGa is not expected to be stable at the growth temperature (1050 8C or even higher). Recent studies of positron annihilation in HVPE GaN reveals a rather high concentration of VGa complexes in nominally undoped material (in the 1017 cm 2 range) [51], and therefore, such defects may well influence the electrical properties and compensation
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Fig. 8. (a) The ODMR spectrum observed in the as-grown Zn-doped GaN sample, at 9.21686 GHz and 5 K. The known ODMR signals from a shallow donor and Zn are denoted by SD and Zn, respectively. (b) The Ga-1 ODMR spectrum after subtracting the other known ODMR signals. (c) The simulated Ga-1 ODMR spectrum taking into account the contribution from the two isotopes of Ga.
in the material. In similar experiments a strong increase of VGa complexes has been observed in Be doped MBE GaN [52]. This is an indication of the introduction of donors during the growth of Be doped GaN, increasing the solubility of VGa. Be is a shallow acceptor in GaN (binding energy about 100 meV), but it has so far not been possible to achieve a sufficiently high p-doping level, indicating a complex behavior of Be in GaN. Similarly, the VN has long been blamed for the ndoping frequently observed in undoped GaN films [53]. In this case theory predicts a shallow donor state for the isolated VN defect [47,54], but the stability at the growth temperature is in doubt. It has been claimed that the VN donor is observed in electrical measurements at room temperature in electron irradiated GaN [55]. In these cases it has not yet been proven that these defects are not complexes involving the VN as one component, however [56]. Recent studies on implantation induced defects indicate a substantial mobility of the intrinsic point defects at room temperature [57]. So far no conclusive magnetic resonance data are available for the vacancy defects in GaN. The Ga interstitial has recently been studied by ODMR in irradiated GaN [58]. It has been found to be unstable already below room temperature, unlike earlier suggestions that the isolated Gai can actually be present in GaN at room temperature [59]. Complexes involving Gai are of course possible as recently reported
Fig. 9. (a) Sketch of the In0.11Ga0.89N/GaN MQW structure grown with mass transport MOVPE. (b) Spatially integrated PL spectrum at 2 K at a low excitation density of 10 mW cm 2.
[60], another example of this has recently been observed by ODMR in Zn-doped GaN [61], see Fig. 8. A Garelated hyperfine structure is clearly seen in Fig. 8, demonstrating together with a trigonal symmetry a defect likely to be related to a Gai complex. Whether the other component in the complex is Zn or another impurity is not clear at the moment. The N interstitial has so far not been positively identified, although it was suggested to be observed in electron irradiated material at room temperature [55].
6. Quantum well structures The III-nitride system offers several interesting material combinations for heterostructures, with a type I band alignment. The InGaN/GaN system has been studied intensely, due to the interest for visible LEDs and lasers [4]. The AlGaN/GaN system is also of great interest, both for HEMT transistor structures and for UV LEDs. We shall first discuss some properties of the InGaN/GaN system.
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Fig. 10. (a) Set of micro-PL spectra of InGaN/GaN MQW structure. (b) PL integrated intensity as a function of the spot position.
6.1. InGaN/GaN QWs One question of prime importance in this materials system is the presence of In segregation during growth of InGaN at elevated temperatures. Theoretically it has been predicted that for compositions above about 10% In the alloy has a strong tendency to segregate so that a second phase rich in In occurs [62]. The presence of such an In rich phase has also been claimed in optical data for In compositions about 10% [63], via a second low energy peak in PL spectra. It appears that this phenomenon depends on the growth conditions, it may be possible to avoid in many cases. In Fig. 9(b) we show PL data from an InGaN/GaN MQW with an In composition in the QWs of 11% (Fig. 9(a)), where the entire structure was
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Fig. 11. (a) Schematic drawing of an AlGaN/GaN MQW LED structure. (b) Panchromatic CL image of the n-AlGaN buffer layer in the structure shown in (a).
grown at one temperature (800 8C). The observed linewidth of the QW excitons is just about 40 meV at 2 K, which is what may be expected from the added contributions to the linewidth of the alloy broadening (about 15 meV) and the interface roughness broadening (about 25 meV). No low energy peak with the properties expected for In rich islands is observed, indicating the absence of segregation. Also, the dislocation density does not seem to influence this linewidth. In Fig. 10(a) is shown a set of micro-PL spectra for this structure, which has stripes of mass transport overgrown areas in the GaN buffer layer, reducing the dislocation density by about an order of magnitude down to the 107 cm 2 range. A scan across both the overgrown and the standard area shows no difference in
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Fig. 12. (a /d) PL decay curves measured at 2 K for different AlGaN/GaN MQWs with different doping. The Si concentration in the barriers is indicated in the figure. The evaluated PL lifetime is about 360 ps for all MQWs.
PL linewidth, only a small difference in PL intensity (Fig. 10(b)). The reason for different results on InGaN/GaN MQW properties between samples grown in different laboratories is not clarified, but is certainly to be found in differences in the growth conditions. We speculate that the procedure of a temperature ramping between the growth of GaN and InGaN layers, respectively, employed by many workers, might have adverse effects on the quality of the InGaN layers, since the surface is exposed to an excessive temperature during the ramping after the growth of the InGaN layer. Evidence for the absence of In segregation even for higher In compositions has recently been reported for MBE grown layers [64]. Such a segregation may even be beneficial in the case of LEDs, since it allows efficient radiative emission from localized excitons at room temperature. For lasers, however, phase separation is expected to reduce the gain, and should be avoided. 6.2. AlGaN/GaN MQWs AlGaN/GaN MQWs are different, since the well material is a binary. The localization of excitons or carriers in this case is expected to be due to interface roughness and the alloy fluctuations in the AlGaN barrier. We have studied such MQW structures grown by MOVPE, where lateral overgrowth was applied in the AlGaN buffer layer (see Fig. 11(a)). CL topographs of the AlGaN buffer layer shows well resolved contrast
from threading dislocations of different types (Fig. 11(b)). The edge dislocations in the coalescence zone show only a weak contrast, in agreement to previous reports [65]. The PL linewidth of the MQW excitons in this case is about 40 meV, i.e. about a factor two larger than the best reported values in MBE samples. This is already a sign of strong localization potentials. Since this MQW was present at the surface of the structure we cannot rule out a broadening due to a potential gradient across the MQW part, so that all QWs are not equivalent. A remarkable result is obtained in a set of MQW (5 QWs) samples with different Si doping, from nominally undoped up to the 1020 cm 2 range. Ideally such a set would be expected to exhibit a strong variation on radiative lifetime t with doping, i.e. t should vary as 1/n where n is the electron density. Instead no variation of the decay time is observed, as shown in Fig. 12. This is most probably explained as due to a strong localization of the holes at all doping levels. Such localization is expected to be particularly strong for rather low Al composition in the AlGaN barriers, as recently suggested [66].
Acknowledgements This work was partially supported by the European Community via the CLERMONT project PHRN-CT1999-00132.
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