|
Volume 9, Article 1
M. Sumiya
Shizuoka Univeristy
and
CREST-JST
Nitride films have conventionally been grown on sapphire substrates by using a number of growth techniques, such as two-step metallorganic chemical vapor deposition (MOCVD) [1], molecular beam epitaxy (MBE), pulsed laser deposition (PLD) and hydride vapor phase epitaxy (HVPE). Figure 2 shows optical microscope images for our MOCVD-GaN films with both smooth and hexagonal-facetted surfaces, which are due to +c and –c polarity, respectively. The results of our study compare well with the standard frameworks for the polarity provided by Hellman [2] and the resulting GaN samples are acceptable within this context.
Some researchers have classified the polarity of GaN films by the growth method used to obtain them, in that films grown by MOCVD and MBE have +c and –c polarity, respectively. In addition, nitridation of sapphire substrate has often been regarded as a way to obtain –c GaN films. However, different groups using the same growth technique have sometimes found that there are conflicts between the polarities that result from their experiments. When GaN films are grown by MBE (in which the nitridation process is commonly used) +c GaN films are the likely outcome under III-rich conditions. On the other hand, although the sapphire substrate was not intentionally nitrided, a hexagonal-facetted surface (indicating –c polarity) was observed for a GaN film on a thin GaN buffer layer grown by MOCVD and reported by Nakamura [3]. In another report by Uchida et al., a smooth surface (indicating +c polarity) was obtained for a GaN film grown on a thicker buffer layer by MOCVD, even though the substrate was intentionally nitrided [4]. These confusions in controlling the polarity are a feature of the growth experience of many research groups.
We have studied the role of the LT-buffer layer and the implications of each part of the process for the growth of GaN films by two-step MOCVD. The conditions used to treat the substrate and the deposition and annealing of the LT-buffer layer have been found to correlate with the polarity of the grown layers. Through these studies, ‘recipes’ to control the polarity of the layers have been suggested. Indeed, the polarity can be managed from +c through to –c, including mixed polarity layers, by systematically varying the conditions used in the MOCVD process (Sec. 4). This paper will critically review the status of polarity control for the MBE, PLD and HVPE growth methods, as published in the literature. General conditions that are decisive for the determination of polarity are itemized for each growth method. Comparing the recipes used in MOCVD growth with the features identified for the other methods, we will examine whether or not the implications derived from our evaluation of the MOCVD process are equivalent to those noted for the other methods. The correlation between polarity and the growth conditions will be discussed in order to clarify the various confusions experienced in the determination of polarity during GaN growth.
In Sec. 2, techniques for evaluating the polarity are summarized in chronological order, and their specific features are identified. We used coaxial impact collision ion scattering spectroscopy [5] (CAICISS) to characterize the polarity. The important features and the advantages of CAICISS are mentioned. The case of InGaN multi-quantum wells is examined to demonstrate the potential of CAICISS analysis. In Sec. 3, the correlation between the polarity and the growth conditions in terms of substrate treatment and buffer layer preparation is summarized for GaN films deposited by MBE, PLD, and HVPE. In Sec. 4, variations in the LT-buffer layer depending on the growth conditions in MOCVD are examined with respect to controlling the polarity. While discussing the implications of each growth process in MOCVD, ‘recipes’ to manage the polarity of the GaN films are proposed. By comparing these recipes with the relevant details of other techniques, it will be concluded that the polar structures at the interfaces of both the sapphire substrate and the annealed LT-buffer layer are the most important aspects of polarity management. Our focus moves on in Sec. 5 to a study of the polarity-dependence of the properties of GaN films and of device performance. The dependence of the properties of GaN on the polarity properties is elucidated with respect to both impurity incorporation and defect formation. The effect on device performance of the internal electrical field due to the polarity in the material is reviewed, focusing on the interface between the metal and the III-V nitride semiconductors. Our work is finally summarized in Sec. 6.
There is a positive or a negative polarization charge at each interface of the multi-quantum-wells (MQWs) that are used as the active layer of LEDs and LDs. The band decline of a very thin well and barrier layer should become mutually opposite, taking into account the flatness of the Fermi level and the depth of the depletion layer. The strength of the electric field within the InGaN active layer has been estimated experimentally to be 0.35 MV/cm from the peak shift of the PL spectra of MQWs with various thickness of wells [9]. Since the spontaneous polarizations of InN and GaN are very close, the internal electric field in the InGaN/GaN system is mainly caused by the difference in their piezoelectric polarizations; e33 (GaN: ~0.4C/m2, InN: ~1C/m2 [10]). In contrast, the formation of a 2DEG of up to 1013 cm-2 at the interface of the AlGaN/GaN system cannot be explained by piezoelectric polarization alone. The effects of spontaneous polarization must be taken into consideration, owing to the large difference in the spontaneous polarizations between AlN (-0.081C/m2) and GaN [11].
Thus, a band profile of a hetero-structure of III-V nitrides can be modified by a combination of both the piezoelectric and spontaneous polarizations of the materials (discussed in Sec. 5). Since the influence of spontaneous polarization is sufficiently large in III-nitride system, the evaluation of the polarity is also important when designing for optimum device performance.
7 years later, Sun et al. found that a GaN film deposited on the Si-face of a 6H-SiC substrate was unstable in an H2 ambient at 600°C [15]. This GaN sample was determined to have +c polarity due to the electron negativity of Si and C, in contrast with Sasaki’s conclusion. We also confirmed the instability of +c GaN under these conditions [16]. However, we felt that the window of the conditions used for annealing temperature, time and gas ambient seemed to be very narrow. Therefore, this technique was difficult to use generally for determining the polarity. In 2001, Koukitsu et al. measured the decomposition rate of +c and –c GaN samples in N2 or H2 gas ambients using the microgravity method [17]. GaN with +c polarity decomposed faster at lower temperatures (800-850°C) than GaN with –c polarity (900-950°C), which was consistent with Sun’s report and Hellman’s standard framework.
1. MOCVD- and MBE-GaN films were likely to have +c and mixed polarity-containing inversion domains (IDs), respectively [20].
2. Nitridation of the sapphire substrate in MOCVD might result in a –c GaN film with a rough surface [21]
3. The dependence of various properties on the polarity was revealed [22] [23].
Claims 1 and 2 indicate the importance of the growth process, as discussed in Sec. 4, and claim 3 is related to the dependence of impurity adsorption on the polar surface, as discussed in Sec. 5
In order to extend the possibilities of techniques based on the TEM, efforts have been made continuously to observe atomic alignment with high-resolution image matching [24], to introduce micro-channeling effects during EDS analysis [25], to detect the N K-edge and Ga L-edge during EELS analysis [26], or to count diffracted electrons quantitatively [27]. The polarity can also be evaluated by these new approaches.
Another technique involves the chemical stability of the –c GaN surface in alkali solution [32]. Apparently, –c GaN films are etched in KOH or NaOH, while +c GaN is inert to these solutions. The polarity that was determined by investigating the chemical stability was consistent with that determined by using hemispherically-scanned x-ray photoelectron diffraction (HSXPD) [6]. We reported the mechanism for the selective etching, in that OH- in solution would promote etching, attacking one back-bond of the Ga that was bonded to the nitrogen on the –c polar surface [33]. The etching effects that originate from side facets or dislocations are still unknown, especially for GaN with rough surface morphology [34]. This chemical stability, however, is the easiest way to determine the polarity of GaN
The methods mentioned above seem not to be suitable for determining the polarity of thin and poor quality GaN such as LT-buffer layers. In ’98, we used CAICISS to determine their polarity [35]. The polarity of LT-GaN buffer layers was successfully evaluated for the first time by CAICISS [36]. This technique has been used for analyzing the polarity of MBE-GaN by Shimizu et al. [37]. The various features and advantages of CAICISS will be discussed in Sec. 2.4.
According to the procedure in Ref. [5] as carried out by Katayama et al., we changed the incident angle α from 90° (normal to sample) towards a lower angle. In CAICISS, He+ and He0 (He particles) are back-scattered along an angle of 180°, due to impact collisions with atoms on the surface, and are detected by a multi channel plate. The intensity of the back-scattered He ions depends largely on the incident angle. Consequently, an angular dependence against an identical atom is obtained, as shown in Figure 4. The better the quality, the deeper the dip ((1) in Figure 4) and the narrower the width of the peak ((2)).
A specific feature of CAICISS is that it is a simple way of quantitatively analyzing the atomic arrangement on the surface, such as the distance and the angle made with neighboring atoms, because the analysis of the scattering orbitals can be extremely simplified by the focusing and shadowing effects and by taking only ions that have impact-collided with atoms into account. Using these features, the atomic structure of the surface (several nm deep) can be non-destructively analyzed in real space with CAICISS. The potential of CAICISS for determining the surface atomic arrangement in real space has been demonstrated for Si surfaces, compound semiconductors [47] [48] and oxide thin films [49].
The polarity can be determined by the positions of the parabolic shadowing dip and the focusing peak of the He+ beam in CAICISS analysis. When the incident angle of a He+ beam that is irradiated from the [112̄0] azimuth is changed, the angular dependence of the cation signal intensity is represented by one or two peaks around 70° for +c and –c polarity GaN, respectively. These angular dependences can be calculated by computer simulation [50]. However, we have experimentally determined the polarity of a GaN film from the angular dependence measured by CAICISS, by comparing it to those shown in Figure 5 for bulk ZnO with Zn- (+c) and O- (-c) face polarity [51]. This was achieved without using the results of simulations. This is possible because ZnO has the same crystalline structure and very similar lattice constants to those of GaN, since they are neighboring elements in the periodic table.
We have to comment here that it is not the termination atoms but the polarity, which can be analyzed by the CAICISS method that we use in this section. The competence of CAICISS for determining the polarity of III-nitride can be best demonstrated when it is applied to very thin films, such as LT-buffer layers and quantum well structures. When the CAICISS technique is to be used for thicker III-nitride samples, cross-checking should be implemented using either CBED or the chemical stability in alkali solution.
The integrated signal for the III-group element (shadow area in Figure 4) from each TOF spectrum with changing angle of incidence was obtained to evaluate their polarity. Surface contaminants such as C and O, which have small cross-sections, would not have any influence on determining the polarity. Figure 7 shows the angular dependence of the intensities of Ga and In in CAICISS-TOF spectra for an In0.5Ga0.5N SQW. It was confirmed from comparison with the results in Figure 5 that the alloyed In has +c polarity, while the variation in the Ga signal also indicates +c polarity. The indium atoms incorporated into the InxGa1-xN SQW were found to occupy substitutional sites for Ga, and they exhibited +c polarity.
Heinlein et al. systematically investigated the time dependence of the nitridation and concluded that it took 200 min of exposure to N-radicals to complete the first monolayer of the surface nitride in an MBE chamber [53]. The nitridation done by Mikroulis et al. (Univ. of Crete) at 200°C caused an improvement in the flatness of the sapphire surface and increased the in-plane lattice constant, typically by a value of approximately 9%. The surface become rougher and the lattice constant could be increased by 6.2-6.8% by nitridation at 750°C [54]. Namkoong et al. (Georgia Inst. of Tech.) confirmed that 6Å of AlN and 23Å of AlN+NO were formed by nitridation at 200°C and 700°C, respectively [55]. N-radicals are so reactive that they can even nitride the surface of a sapphire substrate at a temperature as low as 200°C. This is completely different from nitridation caused by a flow of NH3 gas, as used in MOCVD (refer to Sec. 4.2). In a technique that is analogous to nitridation in MOCVD, a flow of NH3 gas was introduced into the MBE process by Sonoda et al. (AIST) [56] and Held et al. (Univ. of Minnesota) [29]. Although the two groups published no results of this surface nitridation technique, Grandjean et al. pointed out that AlN could also be formed by nitridation using an NH3 flow at 850°C for 10 min [57].
The polarity of the AlN formed by the nitridation process was suggested to be –c polarity from theoretical calculations done by Felice et al. [58] (discussed in Sec. 3.1.4). Indeed, it was reported that initial nitridation of the sapphire substrate favors the growth of GaN films along the –c direction [59].
The polarity could be systematically varied by changing the thickness of the AlN buffer layer. Dimitorov et al. (Cornell Univ.) demonstrated that the polarity could be controlled from –c to +c polarity (through mixed polarity layers) by using buffer layers corresponding to 0 nm, <5nm, and 5-15nm of AlN [60]. GaN films on thinner AlN buffer layers (<12nm) are likely to contain IDs, as reported by Georgia Inst. of Tech. [55]. When GaN films are grown not only on sapphire, but also on GaAs (111) B-face substrates, +c polarity can be realized by increasing the thickness of the AlN buffer layer (>20 nm) [63]. Furthermore, the growth temperature of the AlN buffer layer seems to be important. AlN buffer layers deposited at temperatures higher than those used for depositing HT-GaN layers tend to result in +c polarity, as observed in Georgia Inst. of Tech., the Walter Schottky Inst. and at the Univ. of Minnesota, while the AIST group claim that lower temperatures are more suitable.
Thicker GaN buffer layers lead to +c polarity for MBE-GaN, as reported by Huang et al. Increasing the growth rate of the GaN (increase in the amount of Ga [64]) also seems to be effective (Virginia Commonwealth Univ.) [65]. GaN buffer layers were deposited on NH3-nitrided sapphire by AIST [56] and Univ. of Minnesota [29]. The polarities determined by the two groups were opposite (the former –c, and the latter +c). In addition, the polarity of MBE-GaN was converted from +c to –c by annealing the GaN buffer layer, as was observed in Univ. of Crete [54]. Therefore, the conflict in the use of a GaN buffer layer in MBE was still apparent.
It was very difficult from these papers to find both general conditions that completely covered all of the research groups and any clear relationship between the polarity and the growth conditions. However, we consider through reading many reports that +c GaN would be obtained on an AlN buffer layer deposited under III-rich (higher growth rate) conditions at high temperature in MBE.
A monolayer of Mg deposits on the Ga-terminated surface of the +c GaN film. Since the local structure of the Mg3N2 is more favorable than that of the bulk GaN, the Mg is likely to bond with the N atoms. Consequently, the displacement of the Ga and the N atoms in the outermost layer should occur, forming the configuration Mg-N-Ga/Ga-N from Mg-Ga-N-Ga-N on the +c polar surface. The sign Ga/Ga indicates IDB, consisting of a plane of Ga-Ga bonds. The N atoms are six-fold coordinated with the outermost Mg and the underlying Ga. This surface structure has been theoretically calculated to be the most stable [31] [74]. Recently, a model of pyramidal (zigzag) inversion domains originating from the Mg on the (0001) segment of the boundary was theoretically performed by first-principles pseudopotential density functional calculations [77]. The most favorable structure of the Mg boundary inserted into the GaN was evaluated as abcab stacking across the (0001) segment, corresponding to the atomic sequence GaNMgNGa, where the side of the boundary lies along the {1113} direction, corresponding to the zigzag inversion domain boundary. In this structure, the concentration of Mg in the boundary layer was calculated to be 3/4 monolayer, occupying H3 sites.
The dependence of the polarity on growth rate is also observed for GaN and AlN buffer layers, shown in Table II (a), as well as in ZnO films with the same wurtzite crystal structure [81]. Takahashi et al. deposited GaN films on GaAs (111) A and B-face substrates by MOMBE. The deposition was carried out at 700°C by changing the Ga flux (beam equivalent pressure (BEP); 2~8x10-8Torr) under a constant supply of DMHy as the N source. The growth rate of the GaN on both substrates increased up to 5x10-8Torr of Ga supply, and then saturated at 400nm/h above that level. This indicates that the growth was promoted with a supply-limit (N-rich) for the lower BEP, while it occurred with a surface kinetic limit (Ga-rich) for the higher BEP. GaN films on A-face GaAs (111) were found to have +c polarity, and were independent of the growth rate. On the other hand, the polarity on the GaAs (111) B-face was –c for the supply-limited condition, and it was +c when it was kinetic-limited. GaN grown under III-rich conditions (kinetic limited, Ga- stable or limited growth) is considered to predominantly display +c polarity [82].
The vapor pressure of GaN is much higher than that of AlN, as shown in Figure 8 [83]. AlN buffer layers, with their lower vapor pressure, should be suitable to act as nucleation layers for MBE-GaN grown under high vacuum conditions. Felice et al. calculated theoretically the atomic structure of films consisting of approximately 1 bilayer of AlN on c-plane sapphire substrates [58]. Under equilibrium conditions, the Al layer in the H3 sites lying between the last O plane (blue region) and the first N plane (yellow region) maintain the stoichiometry of bulk sapphire for 2/3 monolayers, as shown in Figure 9 (a) and Figure 9(b). This favorable structure could be changed by the amount of Al in the initial growth, from Al-rich in Figure 9 (a) to Al-poor in Figure 9 (b). Both of the geometries of the AlN on the outermost layer (between the green brackets) corresponded to +c polarities (from previous calculations) [84]. In contrast, the alignment of the AlN (between the red brackets) in Figure 9(c) and Figure 9(d) corresponded to –c polarity. This structural difference is observed in the complete interface of Al adatoms lying in the T4 sites between the last O plane and the N plane. The calculations predicted that the polarity of these very thin films on sapphire substrates would be attributed to that of the III-V nitrides. Moreover, since the two structures in Figure 9 (b) and Figure 9 (c) have very similar formation energies, independent of Al abundance, it is assumed that a slight fluctuation in conditions in the initial stages of deposition can alternately switch the structure. The situation that is shown in Figure 9 (c) must occur under non-equilibrium conditions, such as nitridation of the sapphire under Al-deficient conditions. Consequently, the nitridation of sapphire is regarded as resulting in –c GaN, though moderate nitridation as carried out by the AIST group [56] might form the interface structure shown in Figure 9 (b). Increasing the supply of group III implies that the growth of nitride materials would approach the equilibrium state, forming the structures with +c polarity shown in Figure 9 (a) and Figure 9 (b).
The orientation of GaN films deposited by PLD was GaN [112̄0] // Al2O3 [11̄00], which is similar to that in the cases of MOCVD and MBE. Ohta and Zhou et al. reported that –c GaN was obtained by PLD when deposition occurred directly onto a sapphire substrate without any buffer layer. By introducing 19nm of AlN [87] or 10nm of Al0.8Ga0.2N [88] (as employed by the former and the latter, respectively) the films were reported to be converted to +c polarity. Furthermore, an AlN buffer layer deposited under Al-rich conditions made the GaN film grow along the +c polar direction, and the initial growth on the sapphire substrate was decisive in determining the polarity [89]. GaN grown on a Si-face 6H-SiC substrate had +c polarity. Better-quality GaN films were obtained under Ga-rich conditions [90]. Thus, the trends for GaN films deposited by PLD seem to be similar to those deposited by MBE, as mentioned in Sec. 3.1.
In Table III, reports of GaN growth on sapphire substrates by HVPE are listed with respect to the relationship between initial growth conditions and surface morphology. The first smooth HVPE-GaN was achieved by Naniwae et al. in 1990 [95]. The key technology was the treatment of the sapphire substrate by Ga+HCl at 1030°C. In order to obtain smooth GaN, Ga+HCl treatment should be carried out for more than 20min, as shown in the Table. Two years later, Detchprochm et al. found that the insertion of a ZnO layer on the sapphire substrate made it possible to grow HVPE-GaN that was transparent, with a smooth surface [96]. A 10-300nm thick ZnO buffer layer was deposited on the sapphire substrate by sputtering at room temperature, and then the sapphire substrate that was covered with the ZnO layer was introduced into the HVPE system.
The vapor pressure of ZnO is so high at 1000°C in an HVPE reactor that there was no evidence of ZnO at the GaN/sapphire interface, as reported by Molnar et al. [97]. Ga atoms left on the sapphire by the GaCl treatment should be desorbed as well. However, it is supposed that a very small fraction of the Zn or the Ga could form compositions that would play a role in converting the polarity at the interface of the sapphire.
Wagner et al. reported that HVPE-GaN grown on a GaN buffer layer was smoother than that on an AlN buffer layer deposited by MOCVD [100]. Here, it is worth noting that HVPE-GaN on an LT-GaN buffer layer less than 10nm thick had a hexagonal-facetted surface (discussed in Sec. 4.3 Recipe 1-(2)). The GaN buffer layer could be desorbed during the ramping of the substrate temperature under an NH3 ambient, and then the unintentional direct-growth of HVPE-GaN on the sapphire substrate was supposed to take place. In contrast, smooth HVPE-GaN was deposited on AlN buffer layers. This is why AlN buffer layers, which have a lower desorption rate (as shown in Figure 8) were expected to work as nucleation layers even at the high temperatures used for HVPE-GaN growth.
The importance of variations in the LT-buffer layer caused by annealing can be found in the report by Tavernier et al., who applied a similar buffer layer technology to that used in MOCVD to HVPE in a single chamber [101]. Their studies of buffer layers revealed that layer thickness and annealing conditions are crucial to obtaining HVPE-GaN of high quality.
In HVPE-GaN, the coalescence of the GaN islands occurs rapidly (within 10 sec of their growth) corresponding to 0.4-0.5μm of film thickness [105]. The quality of the HVPE-GaN can be divided into two regimes, thinner than 0.4-0.5μm and thicker than that, which correspond to destructive and better quality, respectively. Gu et al. obtained better quality HVPE-GaN by changing the optimum conditions for the two regions [102]. In addition, in a similar way to the case of MEE [69] given in Sec. 3.1, flow modulated growth (FMG) with a periodically interrupted HCl flow under a constant flow of NH3 can also improve HVPE-GaN, as reported by Zhang et al. [106].
The first GaN films deposited by MOCVD in 1984 had hexagonal-facetted [107] or granular surfaces [108]. By using an LT-AlN buffer layer, GaN with a smooth surface was achieved for the first time in 1986 [1]. It has been thought that this variation could be related to the polarity. Hwang et al. investigated the influence of sapphire nitridation on GaN films in 1995 [109]. Their nitridation was carried out at 900°C for 5min. GaN films using 50nm of LT-GaN buffer layer deposited on nitrided and non-nitrided sapphire substrates had hexagonal-facetted or smooth surfaces, respectively. The carrier density for GaN films on nitrided sapphire was higher by one order of magnitude than that for films on non-nitrided substrates. In addition, the mobility of the former was 1/4 of that of the latter. That is, the substrate nitridation not only had an influence on the surface morphology, but also on the properties of the GaN. Van Der Stricht et al. systematically changed the surface morphology of GaN films deposited on nitrided sapphire substrates from hexagonal-facetted to smooth by lowering the deposition temperature of the GaN buffer layer from 550 to 450°C [110]. The variations (re-crystallization or sublimation [111] [112]) of the LT-buffer layer induced by the gas ambient [113] and the annealing time [114] [115] were investigated intensively to improve the quality of GaN films. Thus, the key technologies for controlling the polarity in MOCVD can usually be found in thesis reports. The implications of these conditions will be discussed with respect to polarity-control in the subsequent sections.
Although the conditions can be independently controlled during each process step, each step has an influence on the subsequent processes. There are infinite combinations of conditions until deposition of the HT-GaN film. The polarities of samples prepared under a representative condition for each process were analyzed by each of CAICISS, X-TEM and chemical stability.
Although the surface of the sapphire substrate was sometimes covered with undesirable contamination (Ga, N, Al, Si etc.) due to hysteresis in the MOCVD apparatus [116], oxygen was conventionally removed from the sapphire surface during H2 cleaning. Consequently, the surface was slightly rougher, and an Al-rich surface was formed, which was confirmed to be Al:O = 50:50% by XPS. When it was subsequently nitrided under flowing NH3 at various temperatures between 600°C and 1080°C, the surface compositions of the Al, O and N were changed, as shown in Fig. 10. Nitrogen was detected for sapphire nitrided at even 600°C, and the nitrogen composition increased with higher temperature, while the oxygen decreased.
Thus, an Al-rich surface was formed on the sapphire substrate by the removal of oxygen during H2 cleaning. In contrast, AlOxN1-x was induced by nitridation, depending on the temperature used. These chemical states at the surface of the substrate play a decisive role in the polar structure of the buffer layer and also the evaporation behavior. (Refer to Recipes 1 and 2).
Our concern was focused on buffer layers on nitrided sapphire. The non-stoichiometric AlOxN1-x layer formed by the nitridation is likely to have –c polarity, as indicated by theoretical calculations [58]. Both +c and –c polarity nucleate simultaneously on the nitrided sapphire, probably due to either inhomogeneous nitridation [117] or favorable formation energy. The surface of the as-deposited buffer layer would be covered with a +c layer with higher growth rate [118]. Subsequent annealing of the film would make the –c domains rise to the surface due to the sublimation of the film. Based on these considerations, buffer layers on nitrided sapphire substrate are assumed to be covered with a +c layer grown laterally over the –c domains.
To confirm this assumption, thicker buffer layers (210nm) were prepared on nitrided sapphire and annealed for various times at 1040°C. An LT-GaN buffer layer on a nitrided substrate was found to evaporate with a layer-by-layer mode due to the AlOxN1-x on its surface [119]. Figure 12 shows the angular dependence of the Ga signal intensity in CAICISS for the thicker buffer layer annealed for 0, 10, 20 and 30min [120]. The CAICISS result for the as-deposited sample shows predominantly +c polarity [Figure 12 (a)]. The sample continues to have +c polarity after annealing for 10min. Sharpening of the peaks is also observed in the result, suggesting an improvement in the crystal quality near the surface [Figure 12 (b)]. The peak at 72° splits into two peaks and the peak at 35° intensifies with further annealing, as shown in Figs. 12 (c) and (d). This indicates that the film surface is transforming from +c polarity (triangles) to –c polarity (squares). The lines in Figs. 12 (c) and (d) present the weight ratio of +c: -c polarity at 5:5 and 2:8, respectively, with the assumption that they share the same crystal quality (i.e., the same intensity of the CAICISSS signal for +c and –c domains).
Since CAICISS analysis detects the atomic arrangement of only the surface region (as discussed in Sec. 2.4), the films were further investigated by TEM [120]. Figure 13 shows X-TEM images for the same samples that are shown in Figs. 12 (c) and 12 (d). Columnar IDs were found to extend to the surface, and dome-shaped domains were found near the interface, similar to the observations reported by Wu et al. [121]. With increasing annealing time (sublimation of the film), the dome-shaped domains are exposed to the surface. The TEM images are consistent with the CAICISS spectra. Since the –c signal component of the CAICISS result increased after annealing, the inverted domains (dome shaped) can be considered to have –c polarity.
Recently, we have deposited a GaN film by MOCVD on an H2 cleaned sapphire substrate that had once been exposed to air. Although the surface should have been terminated with oxygen, a +c GaN film with sufficient quality was obtained on the sapphire by two-step MOCVD without the need for a second H2 cleaning process [122]. The sapphire surface becomes rougher by H2 cleaning, represented by the weak RHEED pattern. In addition, a comparable GaN film with high quality was even grown on a sapphire substrate cleaned in an N2 ambient at more than 1000°C [123]. It is supposed from these facts that the thermal roughening of sapphire at higher temperature might be important as a nucleation site for the growth of the LT-buffer layer.
The effect of the thickness of the buffer layer can be seen in the data for MBE-GaN films shown in Table II (a). In spite of the presence of an AlN or GaN buffer layer, MBE-GaN films on thicker buffer layers are likely to have +c polarity. When the group at Georgia Inst. of Tech. [62] increased the thickness of the AlN from 12 nm to 30 nm, the polarity was converted from mixed polarity to +c polarity. Similar results were obtained by both the Virginia Commonwealth University, who used GaN buffer layers more than 60 nm thick [65] and the Walter Schottky Institute, who used AlN buffer layers that were more than 5 nm in thickness [60]. These examples of polarity conversion depending on the thickness of the buffer layer are exactly similar to our case for MOCVD-GaN.
The FWHM of the Ga 3d spectrum detected by XPS for the former as-deposited sample was 2.2 eV wider than that (1.5-1.7eV by our analysis) for the latter layers [119], suggesting the existence of Ga metal or Ga that was weakly bonded with the N. Thus, the III-rich condition is suggested as being effective for obtaining a +c interface, as discussed in section 3.1.4 for MBE, as well as in MOCVD. Furthermore, the effect of the III-rich condition on the +c polarity might be extended to the HVPE technique using the GaCl treatment mentioned in Sec. 3.3.
The annealing of the LT-buffer layer in MOCVD is a unique process. To add to the importance of the V/III ratio, the annealing conditions are also crucial to the polarity in terms of controlling the thickness of the buffer layer. An H2 ambient should be used for thicker or III-poor buffer layers, while an N2 ambient is appropriate for a thin or III-rich layer in order to obtain a buffer layer of the optimum thickness. These correlations are very complex, depending on the individual growth apparatus. Our case is referred to in detail elsewhere [119].
Two recipes for obtaining +c GaN on nitrided sapphire are; 1) the deposition of a thicker buffer layer under III-rich conditions, and 2) annealing of the layer for a short time under an N2 ambient. We suppose that Uchida et al. in Ref. 4, who deposited +c GaN on nitrided sapphire substrates, managed to use these correlations in their MOCVD apparatus.
It was confirmed that nitridation using NH3 gas at temperatures higher than 700°C resulted in –c GaN films in our deposition system. In addition, the introduction of a flow of NH3 into the reactor, even for several seconds at 1080°C, resulted in the growth of –c GaN films. Therefore, unintentional nitridation probably takes place in the following circumstances; 1) when NH3 is introduced into the reactor during the decrease in substrate temperature for the deposition of the LT-buffer layer after H2 cleaning of the sapphire, and 2) when a longer time is used and annealing is carried out in an H2 ambient for a thin LT buffer layer (exposure of the sapphire surface). These factors indicate the importance of correct timing of the switching of the source gases when shifting to a subsequent part of the growth process.
Seelmann-Eggebert et al. have already pointed out that the occurrence of inversion domains in the films could mainly be attributed to poor process control during substrate cleaning and in the very initial stages of the nucleation process, preceding buffer growth [6]. It is supposed that the GaN with the hexagonal facets on a thin buffer layer that was reported in Ref. [3] might originate from the unintentional nitridation of the sapphire substrate.
Figure 18 shows the relationship between the FWHM and the annealing time for 20 nm thick AlN and GaN buffer layers. The annealing time means the time interval from the end of the buffer layer deposition at 600°C till the start of the HT-GaN deposition at 1040°C. After taking 7 min to reach 1040°C, the annealing was maintained at 1040°C for the remainder of the time. All of the samples exhibited +c polarity. Although the best quality GaN was obtained on an AlN buffer layer, the values changed drastically in the narrow window of the annealing-time conditions. Figure 19 shows the FWHM and surface morphology (polarity) of HT-GaN films on nitrided sapphire substrates using AlN and GaN buffer layers of various thicknesses. The buffer layers were annealed for 10min. As discussed in Recipe 1, the material changed from –c to +c polarity (through a mixed polarity condition) for both types of buffer layer with increasing buffer layer thickness. However, the window for the AlN buffer layer is narrower, as in the case shown in Figure 18. Therefore, LT-GaN buffer layers may have an advantage for controlling the polarity of GaN films grown by MOCVD.
The TMA precursor was introduced into the reactor 10 seconds before the NH3 gas during the deposition of an HT-AlN buffer layer at 1040°C. As expected, a +c GaN film was achieved. Judging from the deposition rate, the thickness of the Al metal layer was estimated to be 4 Å during this 10 second period. In another experiment, a lower V/III ratio of less than 1800 was used for the AlN deposition at 1040°C, which resulted in +c GaN with a smooth surface, as shown in Figure 20. These two recipes, 1) the deposition of Al metal and 2) the use of a lower V/III ratio, seem to be consistent with the features seen in MBE.
We have investigated the implications in the exact growth sequence and conditions in MOCVD. When summarizing them in Articles 1-5 and Recipes 1-3 with respect to the polarity, the relationship between the growth conditions and the polarity can be represented as a branching road map on a timing chart of the MOCVD process, as shown in Figure 21 [130]. The route is divided by the nitridation of the sapphire substrate. In order to achieve +c GaN in MOCVD, the sapphire substrate must not be nitrided after the H2 cleaning. In the case where nitridation occurs, however, the polarity can still be controlled by modifying the preparation of the LT-buffer layer.
By modifying the timing of the introduction of the source gases according to Recipe 2, growth of GaN was achieved on a Si (111) substrate using only an AlN buffer layer. The key point of this was that the TMA and the NH3 gas should arrive simultaneously at the Si substrate in order to prevent the Al and N from alloying with the Si [132].
GaN films were deposited on (La0.29, Sr0.71)(Al0.65, Ta0.35) O3 (LSAT) (111) substrates, which have a lattice constant that corresponds to the 3x3 structure of GaN (0001) and a thermal expansion coefficient close to that of GaN. Since the LSAT substrate was deteriorated by NH3 and TMG gases at high temperature, an AlN layer was used as a blocking layer to protect the surface. The GaN film on the LSAT had +c polarity and its GaN [11̄00] // LSAT [11̄0] orientation was rotated in-plane by 30° against the expected orientation (GaN [21̄ 1̄0] // LSAT [11̄0]) [133]. This was probably caused by the bond configuration of the surface of the LSAT substrate. In addition, GaN films on metal-face and O-face LiGaO2 (001) substrates had +c and –c polarity, respectively [134].
In this section, the polarity-dependence of the optical properties of GaN films is discussed with respect to impurity-incorporation and defect formation. Furthermore, the structures of interfaces are summarized with respect to the contact formed between the metal and GaN polar surface, which are not fully discussed in these excellent reviews.
Li et al. found that the doping behavior of Mg and the resulting conductivity of the doped layers in MBE-GaN strongly depends on the polarity of the GaN [139]. In fact, Mg was incorporated by up to a factor of 30 times more into +c GaN [140]. Activation of the p-type dopants was achieved for +c GaN. One of the reasons why this was possible is thought to be that C and O impurities are incorporated more into –c GaN [141]. GaN films on bulk GaN single crystals with –c polarity grown by MOCVD [142] and films on ZnO with O-face polarity grown by MBE [143] were also found to contain more oxygen impurities in the –c GaN, respectively, which was confirmed by IR and SIMS. This impurity-dependence on the polarity is consistent with our experiments with MOCVD-GaN, except for the case of Si [137]. Comparable Si incorporation into both types of GaN films has been suggested by both Ng et al. and ourselves, while Li et al. reported higher incorporation of Si in –c GaN. The issue of Si impurity incorporation is still therefore controversial.
The dependence of impurity incorporation on the polarity resulted in differences in the IR [142] and Raman spectra [41] as in Sec. 2.1. The dependence of the optical properties and of defect formation on the polarity was investigated on our +c and –c GaN films. The PL spectrum of +c GaN at 8K [144] exhibits free A and B exciton emission [FE(A) and FE(B)] at 3.492 and 3.499eV, respectively, as shown in Figure 23 (a). The recombination of the first excited states of the A exciton [FE(An=2)] is also found at 3.510eV. These assignments are based on the position of the energies of the respective exciton absorption peaks. The sample also exhibits exciton recombination at 3.486 and 3.464 eV, which are due to the bound-to-neutral donor (D0, X) and neutral deep acceptor (A0d, X) transitions, respectively. The PL peak at 300K is assigned to the free exciton emission, as revealed from Figure 23 (c). Conversely, -c GaN exhibits a rather broader PL band and an absorption tail, as in Figure 23 (b) and (d). The spectrum at 8K exhibits a peak at 3.475eV and a shoulder around 3.45 eV. Since the residual electron density of the –c GaN is as high as 3.5x1018cm-3, the disappearance of excitonic absorption due to Coulomb screening by the increase in temperature above 75K is reasonable. However, -c GaN exhibits a Stokes shift of nearly 20 meV at 300K, although high quality Si-doped +c GaN (with nearly the same electron density of 2.2 x1018cm-3) did not show any Stokes shift at 300K. Therefore, the formation of an impurity-induced band tail is probable in –c GaN, which subsequently causes band gap narrowing due to potential fluctuations and inhomogeneous distribution of the fixed charges. The donor impurity is considered to be O, which is more readily incorporated into –c GaN, as revealed by SIMS in Figure 22. To determine the possibility for the incorporation of acceptor-type defects such as Ga vacancies [145], mono-energetic slow positron annihilation measurements were carried out. Indeed, the S parameter of –c GaN [146] is greater than that of +c GaN, as shown in Figure 24. The fitting results of the relationship between E and S are also shown by the solid lines, and the diffusion lengths of positrons are derived to be 23nm for +c GaN and 4.8nm for –c GaN. These results imply that –c GaN contains a higher density of vacancy-type defects or defect-complexes than +c GaN. Therefore, the formation of extended band-tail states in –c GaN is considered to be due to the simultaneous distribution of donors and acceptor-type vacancy defects, which occurred during the growth of the –c polarity GaN material.
Rickert et al. studied the SBH for thin metal over-layers of Au, Al, Ni, Ti, Pt and Pd on n- and p-type +c GaN samples using synchrotron radiation-based x-ray photoemission spectroscopy [147]. Figure 25 shows the values of the SBH of the six kinds of metals as a function of the work function of the metals. The relationship between the SBH and the work function did not obey the perfect Schottky barrier model. The change in the barrier height was smaller than the value expected from the change in the metal work function. These results indicate the importance of the effect of the surface state and the interface state (refer to Sec. 5.4). To explain this appropriately, a model that takes into account the pinning of the Fermi level at the surface would be required. The position of the Fermi level on the surface of the GaN can be changed by the kind of metal, the conduction type of the GaN, or the chemical treatment of the surface. In addition, the SBH for a Pt/n-GaN structure was reported to be about 1.6 eV, which is comparable to that obtained by C-V measurement and HRPES.
Jang et al. compared the characteristics of an ohmic contact using Ti/Al/Ni/Au metal on +c GaN with that on a –c GaN sample [44]. The contact resistivities, as determined by the transfer length method (TLM), were 8.3×10-4Ωcm2 for the +c sample and 7.0×10-2Ωcm2 for the -c sample. These samples were annealed at 700°C for 1 min. Although the net carrier concentration of the -c sample was higher than that of +c sample, the contact resistivity on the +c GaN was lower by two orders of magnitude than that on the –c GaN.
It is well known that a thin AlN layer can be formed at the interface between the metal and the GaN after annealing a Ti/Al contact at a temperature higher than 400°C. Luther et al. explained the lower contact resistivity for a +c sample as being due to polarization, that is, a 2DEG could be induced by polarization at the +c AlN/GaN interface, but not at the –c interface [148]. Therefore, the effective SBH for the tunneling probability of an electron through the AlN epilayer in a +c sample is considered to be much higher than that for a –c sample. Kwak et al. investigated the effects of polarity on the electrical properties of Ti/Al contacts for n-type GaN [149]. They reported that Ti/Al contacts on n-type +c GaN become ohmic with a contact resistivity of 5×10-5Ω cm2, while a Schottky contact with a barrier height of over 1 eV was formed for contacts on n-type –c GaN.
The effects of polarity on ohmic contacts to p-type GaN were also investigated. Band bending, which is caused by the polarization charge of a thin InGaN capping layer on a p-GaN layer, is utilized to realize the ohmic contact. Gessmann et al. discussed the effect of thin strained cap layers on the contact resistance of p-type +c GaN [150] [151]. Two cases were investigated; a GaN cap/AlGaN structure and an InGaN cap/GaN structure. The electric field in the strained cap layers can reduce the thickness of the tunneling barrier at the metal/semiconductor interface. Since band bending due to the capping layer can be induced by an internal electric field, the thickness of the capping layer must be optimized so that two conditions, the formation of the 2DEG and sufficient tunneling probability for holes, could be satisfied simultaneously. The specific contact resistances were obtained experimentally using the TLM-method. These were 6×10-3Ω cm2 for InGaN (2nm) /p-GaN (Ni/Au contact: annealing at 500°C) and 7×10-4 Ω cm2 for GaN (10nm)/AlGaN (Pd/Au contact: annealing at 500°C). These results indicated the advantageous effect of polarization fields in the cap layer on the reduction of ohmic contact resistance.
Several methods of reducing the surface states have been proposed. The treatment with (NH4)2Sx carried out by Lin et al. [153] caused the band bending of a p-type +c GaN surface to release by 0.25eV. The interface states of an n-type +c GaN surface were decreased to 1×1012 cm-2eV-1 by treatment with an N2 plasma, which was about 1/5 of the level for a non-treated surface (experiment carried out by Hashizume et al. [154]). Furthermore, a decrease in the electron affinity of Cs-adsorbed AlN [155] and GaN [156] surfaces was observed due to the effect of the Cs-surface dipole, which was formed by an initial interaction between the Cs and the empty surface states. Unfortunately, these experimental data were not discussed in connection with spontaneous polarization, which induced sheet charges. Although a method of controlling the surface states has not yet been established, the control of charges caused by surface states and spontaneous polarization is considered to be one of the new topics that could realize new functionality from nitride semiconductors.
GaN films grown by MBE, PLD, HVPE and MOCVD have been reviewed, mainly with respect to the control of polarity. The first three are based on the published literature, while the last is based on our studies. Focusing on the growth conditions in the substrate treatment and/or the preparation of LT-buffer layers, a set of conditions for obtaining +c GaN films is provided for each growth method, as follows;
(1) MBE and PLD (Sec. 3.1 and 3.2)
CAICISS can be used to evaluate the polarity of thin films such as buffer layers (Sec. 2.4). The polarity of an LT-buffer layer was detected for the first time by using CAICISS (Sec. 4.3). We have studied the role of the LT-buffer layer and its implications in each process for GaN film growth by two-step MOCVD, not only in terms of controlling the polarity but also for improving material quality. A road map representing the correlation between the polarity and the growth conditions in MOCVD has been made, as shown in Figure 21. Indeed, the polarity can be managed from +c to –c though mixed polarity by systematically varying the conditions used in MOCVD-GaN. Although the initial growth on sapphire would determine the polarity of the epitaxial materials, the polarity can be controlled by understanding the structure of the LT-buffer layer and its implications for the MOCVD process. It can be concluded through these studies that the polar structure at the interface of an annealed LT-buffer layer is the most crucial factor in determining the polarity of a GaN film.
Recently, the importance of polarity was highlighted in the state-of-the-art growth of InN films [157] [158], and–c GaN with a smooth surface has apparently been obtained by MOCVD [159]. Moreover, we have discovered that an HNO3 treatment of H2-cleaned sapphire substrates resulted in –c GaN film growth in MOCVD [160]. Thus, even more interesting results relating to film polarity have been reported. We intend to further promote studies of the polarity of III-nitrides for applications that utilize the effects of polarity.
[1] H. Amano, N. Sawaki, I. Akasaki , Y. Toyoda , Appl. Phys. Lett. 48, 353-355 (1986).
[2] E. S. Hellman, MRS Internet J. Nitride Semicond. Res. 3, 11 (1998).
[3] S. Nakamura, Jpn. J. Appl. Phys. 30, L1705-L1707 (1991).
[4] K. Uchida, A. Watanabe, F. Yano, M. Kouguchi, T. Tanaka, S. Minagawa , J. Appl. Phys. 79, 3487-3491 (1996).
[5] M. Katayama, E. Nomura, N. Kanekawa, H. Soejima, M. Aono, Nucl. Instrum. Methods B 33, 857 (1988).
[6] M. Seelmann-Eggebert, J. L. Weyher, H. Obloh, H. Zimmermann, A. Rar, S. Porowski, Appl. Phys. Lett. 71, 2635-2637 (1997).
[7] F. Bernardini, V. Fiorentini, D. Vanderbilt, Phys. Rev. B 56, R10024 (1997).
[8] U. Karrer, O. Ambacher, M. Stutzmann, Appl. Phys. Lett. 77, 2012 (2000).
[9] S. F. Chichibu, A. C. Abare, M. S. Minsky, S. Keller, S. B. Fleischer, J. E. Bowers, E. Hu, U. K. Mishra, L. A. Coldren, S. P. DenBaars, T. Sota, Appl. Phys. Lett. 73, 2006 (1998).
[10]F. Bernardini, and V. Fiorentini, http://xxx.lanl.gov/abs/condmat/9808098
[11] N. Maeda, T. Nishida, N. Kobayashi, M. Tomizawa, Appl. Phys. Lett, 73 (1998).
[12] T. Sasaki, T. Matsuoka , J. Appl. Phys. 64, 4531-4535 (1988).
[13] SY Ren, JD Dow, Appl. Phys. Lett. 69, 251-253 (1996).
[14] SY Ren, JD Dow, J. Electron. Mater. 26, 341-346 (1997).
[15] C. J. Sun, P. Kung, A. Saxler, H. Ohsato, E. Bigan, M. Razeghi , D. K. Gaskill , J. Appl. Phys. 76, 236-241 (1994).
[16] S. Fuke, H. Teshigawara, K. Kuwahara, Y. Takano, T. Ito, M. Yanagihara, K. Ohtsuka, J. Appl. Phys. 83, 764 (1998).
[17] A. Koukitu, M. Mayumi, Y. Kumagai, J. Cryst. Growth 246, 230 (2002).
[18] F.A. Ponce, D.P. Bour, W.T. Young, M. Saunders, J.W. Steeds, Appl. Phys. Lett. 69, 337-339 (1996).
[19] Z. Liliental-Weber, C. Kisielowski, S. Ruvimov, Y. Chen, J. Washburn, I. Grzegory, M. Bockowski, J. Jun, S. Porowski, J. Electron. Mater. 25, 1545 (1996).
[20] LT Romano, JE Northrup, MA O'Keefe, Appl. Phys. Lett. 69, 2394-2396 (1996).
[21] J. L. Rouviere, M. Arlery, R. Niebuhr, K. H. Bachem, Olivier Briot, MRS Internet J. Nitride Semicond. Res. 1, 33 (1996).
[22] S. Keller, B. P. Keller, Y.-F. Wu, B. Heying, D. Kapolnek, J. S. Speck, U. K. Mishra, S. P.DenBaars , Appl. Phys. Lett. 68, 1525-1527 (1996).
[23] F. A. Ponce, D. P. Bour, W. Gotz , P. J. Wright , Appl. Phys. Lett. 68, 57-59 (1996).
[24] J. N. Stirman, F. A. Ponce, A. Pavloska, I. S. T. Tsong, D. J. Smith, Appl. Phys. Lett. 76, 822 (2000).
[25] N. Jiang, T. J. Eustis, J. Cai, F. A. Ponce, J. C. H. Spence, J. Silox, Appl. Phys. Lett. 80, 389 (2002).
[26] X. Kong, G. Q. Hu, X. F. Duan, Y. Lu, X. Liu, Appl. Phys. Lett. 81, 1990 (2002).
[27] H. W. Zandbergen, J. Janzen, A. R. A. Zauner, J. L. Weyher, J. Cryst. Growth 210, 167 (2000).
[28] A. R. Smith, R. M. Feenstra, D. W. Greve, J. Neugebauer, J. E. Northrup, Phys. Rev. Lett. 79, 3934 (1997).
[29] R. Held, G. Nowak, B.E. Ishaug, S.M. Seutter, A. Parkhomovsky, A.M. Dabiran, P.I. Cohen, I. Grzegory, S. Porowski, J. Appl. Phys. 85, 7697-7704 (1999).
[30] AR Smith, RM Feenstra, DW Greve, M-S Shin, M Skowronski, J Neugebauer, J Northrup, J. Vac. Sci. Technol. B 16, 2242-2249 (1998).
[31] R. M. Feenstra, J. E. Northrup, Jörg Neugebauer, MRS Internet J. Nitride Semicond. Res. 7, 3 (2002).
[32] J. L. Weyher, S. Müller, I. Grzegory, S. Porowski, J. Cryst. Growth 182, 17-22 (1997).
[33] D. Li, M. Sumiya, S. Fuke, D. Yang, D. Que, Y. Suzuki, Y. Fukuda, J. Appl. Phys. 90, 4219 (2001).
[34] M. Losurdo, M. Giangregorio, P. Capezzuto, G. Bruno, G. Namkoong, W. A. Doolittle, A. S. Brown, Mater. Res. Soc. Symp. Proc. 722, K3.4.1 (2002).
[35]"CAICISS analysis of GaN gilms grown on sapphire substrate by MOCVD method", M. Sumiya, T. Ohnishi, H. Teshigawara, M. Tanaka, I. Ohkubo, M. Kawasaki, M. Yoshimoto, K. Ohtsuka, H. Koinuma, and S. Fuke, Proc. of the 2nd Intern. Symp. On Blue Laser and Light Emitting Diodes, Chiba, Japan, 339 (1998)
[36] M. Sumiya, M. Tanaka, K. Ohtsuka, S. Fuke, T. Ohnishi, I. Ohkubo, M. Yoshimoto, H. Koinuma, M. Kawasaki, Appl. Phys. Lett. 75, 674 (1999).
[37] S. Shimizu, Y. Suzuki, T. Nishihara, S. Hayashi, M. Shinohara, Jpn. J. Appl. Phys. 37, L703 (1998).
[38] A. Kazimirov, N. Faleev, H. Temkin, M. J. Bedzyk, V. Dmitriev, Yu. Melnik, J. Appl. Phys. 89, 6092 (2001).
[39] M. Tabuchi, N. Matsumoto, Y. Takeda, T. Takeuchi, H. Amano, I. Akasaki, J. Cryst. Growth 189/190, 291 (1998).
[40] R. Dimitrov, V. Tilak, M. Murphy, W. J. Schaff, L. F. Eastman, A. P. Lima, C. Miskys, O. Ambacher, M. Stutzmann, Mater. Res. Soc. Symp. Proc. 622, T4.6.1 (2000).
[41]"Raman characterizatiion of an intentionally created Inversion domain boudary in GaN", A. Cros, N. V. Joshi, T. Smith, A. Cantarero, G. Martines-Criado, O. Ambacher and M. Stutzmann, the 5th international Conference on Nitrides semiconductors, Nara, Japan. Technical Digest Mo-P1.109 (2003)
[42] K. M. Jones, P. Visconti, F. Yun, A. A. Baski, H. Morkoc, Appl. Phys. Lett. 78, 2497 (2001).
[43] B. J. Rodriguesz, A. Gruverman, A. I. Kingon, R. J. Nemanich, O. Ambacher, Appl. Phys. Lett. 80, 4166 (2002).
[44] H. W. Jang, J. H. Lee, J. L . Lee, Appl. Phys. Lett. 80, 3955 (2002).
[45] B Daudin, JL Rouviere, M Arlery, Appl. Phys. Lett. 69, 2480-2482 (1996).
[46] M. M. Sung, J. Ahn, V. Bykov, J. W. Rabalais, D. D. Koleske, A. E. Wickenden, Phys. Rev. B 54, 14652-14663 (1996).
[47] M. Katayama, R. S. Williams, M. Kato, E. Nomura, M. Aono, Phys. Rev. Lett. 66, 2762 (1991).
[48]Papers related to CAICISS analysis are listed at http://cobalt.ele.eng.osaka-u.ac.jp/~katayama/caiciss.html
[49]T. Ohnishi Doctoral thesis entitled by ‘Atomic scale analysis and engineering of oxide thin films’, Tokyo Inst. of Tech. 1999
[50] S. Sonoda, S. Shimizu, Y. Suzuki, K. Balakrishnan, J. Shirakashi, H. Okumura, T. Nishihara, M. Shonohara, Jpn. J. Appl. Phys. 38, L1219 (1999).
[51] T. Ohnishi, A. Ohtomo, M. Kawasaki, K. Takahashi, M. Yoshimoto, H. Koinuma, Appl. Phys. Lett. 72, 824 (1998).
[52] M. Sumiya, S. Nakamura, S. F. Chichibu, K. Mizuno, M. Furusawa, M. Yoshimoto, Appl. Phys. Lett. 77, 2512 (2000).
[53] C. Heinlein, J. Grepstad, T. Berge, H. Riechert, Appl. Phys. Lett. 71, 341 (1997).
[54] S. Mikroulis, A. Georgakilas, A. Kostopoulos, V. Cimalla, E. Dimakis, Ph. Komniou, Appl. Phys. Lett. 80, 2886 (2002).
[55] G. Namkoong, W. A. Doolittle, A. S. Brown, M. Losurdo, P. Capezzuto, G. Bruno, J. Appl. Phys. 91, 2499 (2001), J. Vac. Sci. Technol. B 20, 1221 (2002).
[56] S. Sonoda, S. Shimizu, X. Q. Shen, S. Hara, H. Okumura, Jpn. J. Appl. Phys. 39, L202 (2000).
[57] N Grandjean, J Massies, M Leroux, Appl. Phys. Lett. 69, 2071-2073 (1996).
[58] R. Di Felice, J. E. Northrup, Appl. Phys. Lett. 73, 936 (1998).
[59] S. Sonoda, S. Shimizu, Y. Suzuki, K. Balakrishnan, J. Sirakashi, H. Okumura, Jpn. J. Appl. Phys. 39, L73 (2000).
[60] R. Dimitrov, M. Murphy, J. Smart, W. Schaff, J. S. Shealy, L. F. Eastman, O. Ambacher, M. Stutzmann, J. Appl. Phys. 87, 3375 (2000).
[61] C. Piquette, P. M. Bridger, R. A. Beach, T. C. McGill, Mater. Res. Soc. Symp. Proc. 537, G3.77 (1999).
[62] Gon Namkoong, W. Alan Doolittle, April S. Brown, Maria Losurdo, Maria M. Giangregorio, Giovanni Bruno, J. Cryst. Growth 252, 159 (2003).
[63] F. Hasegawa, O. Takahashi, T. Nakayama, R. Souda, Phys. Stat. Sol. B 228, 549 (2001).
[64] O. H. Hughes, T. S. Cheng, S. V. Novikov, C. T. Foxon, D. Korakakis, N. J. Jeffs, J. Cryst. Growth 201/201, 388 (1999).
[65] D. Huang, P. Visconti, K. M. Jones, M. A. Reshchikov, F. Yun, A. A. Baski, T. King, H. Morkoc, Appl. Phys. Lett. 78, 4145 (2001).
[66] K. Xu, N. Yano, A. W. Jia, A. Yoshikawa, K. Takahashi, Phys. Stat. Sol. B 228, 523 (2001).
[67] D. H. Lim, K. Xu, Y. Taniyasu, K. Suzuki, S. Arima, B. Liu, K. Takahashi, A. Yoshikawa, Proc. Int. Workshop on Nitride Semicond, 1, 150 (2000).
[68] Y. S. Park, H. S. Lee, J. H. Ha, H. J. Kim, Sang Man Si, Hwa-Mok Kim, T. W. Kang, Jae Eung Oh, Appl. Phys. Lett. 94, 800 (2003).
[69] A. Kikuchi, T. Yamada, K. Kusakabe, D. Sugihara, S. Nakamura, K. Kishino, 1, 154 (2000).
[70] S. Yoshida, J. Appl. Phys. 87, 1673 (1997).
[71]S. Yoshida, Oyo Buturi 68, 787 (1999)
[72] A. Ishida, E. Yamamoto, K. Ishino, K. Ito, H. Fujiyasu , Y. Nakanishi , Appl. Phys. Lett. 67, 665-666 (1995).
[73]K. Okuno, M. Sumiya, S. Fuke, 11a-N-3, Extended Abstract (The 62nd Fall Meeting 2001) The Japan Society of Applied Physics and Related Societies
[74] V. Ramachandran, R. M. Feenstra, W. L. Sarney, L. Salamanca-Riba, J. E. Northrup, L. T. Romano, D. W. Greve, Appl. Phys. Lett. 75, 808 (1999).
[75] L. T. Romano, J. E. Northrup, A. J. Ptak, T. H. Myers, Appl. Phys. Lett. 77, 2479 (2000).
[76] N. Grandjean, A. Dussaigne, S. Pezzagna, P. Vennegues, J. Cryst. Growth 251, 460 (2003).
[77] J. E. Northrup, Appl. Phys. Lett. 82, 2278 (2003).
[78] EJ Tarsa, B Heying, XH Wu, P Fini, SP DenBaars, JS Speck, J. Appl. Phys. 82, 5472-5479 (1997).
[79] R. Held, D. E. Crawford, A. M. Johnston, A. M. Dabiran, P. I. Cohen, Surf. Rev. Lett. 5, 913-934 (1998).
[80] T. H. Myers, L. S. Hirsch, L. T. Romano, M. R. Richards-Babb, J. Vac. Sci. Technol. B 16, 2261 (1998).
[81] I. Ohkubo, A. Ohtomo, T. Ohnishi, Y. Matsumoto, H. Koinuma, M. Kawasaki, Surf. Sci. Lett 443, L1043 (1999).
[82] O. Takahashi, T. Nakayama, R. Souda, F. Hasegawa, Phys. Stat. Sol. B 228, 529 (2001).
[83] J.H. Edgar, S.Strite, I.Akasaki, H.Amano, C.Wetzel, Properties, processing and applications of GaN and related semiconductors (INSPEC, Institution of Electrical Engineers, London, UK, 1999) .
[84] J. E. Northrup, R. Di Felice, J. Neugebauer, Phys. Rev. B 55, 13878 (1997).
[85] R. D. Dispute, H. Wu, K. Jagannadham, J. Narayan, Mater. Res. Soc. Symp. Proc. 395, 325 (1995).
[86] J. Ohta, H. Fujioka, M. Ohsima, Appl. Phys. Lett. 83, 3060 (2003).
[87] J. Ohta, H. Fujioka, M. Furusawa, A. Sasaki, M. Yoshimoto, H. Koinuma, M. Sumiya, M. Ohsima, J. Cryst. Growth 237-239, 1153 (2002).
[88] H. Zhou, T. Rupp, F. Phillipp, G. Henn, M. Gross, A. Ruhm, H. Schroder, J. Appl. Phys. 93, 1933 (2003).
[89] J. Ohta, H. Fujioka, M. Oshima, K. Fujiwara, A. Ishii, Appl. Phys. Lett. 83, 3075 (2003).
[90] S. Oktybrsky, K. Dovidenko, A. K. Sharma, V. Joshkin, J. Narayan, Mater. Res. Soc. Symp. Proc. 537, G6.43 (1999).
[91] A Usui, H Sunakawa, A Sakai, AA Yamaguchi, Jpn. J. Appl. Phys. 36, L899 (1997).
[92] Y. Kawaguchi, S. Nambu, H. Sone, M. Yamaguchi, H. Miyake, K. Hiramatsu, N. Sawaki, Y. Iyechika, T. Maeda, Mater. Res. Soc. Symp. Proc. 537, G4.1 (1999).
[93] H. P. Maruska, J. J. Tietjen, Appl. Phys. Lett. 15, 327 (1969).
[94] B. Monemar, O. Lagerstedt, H. P. Gislason, J. Appl. Phys. 51, 625 (1980).
[95] K. Naniwae, S. Itoh, H. Amano, K. Itoh, K. Hiramatsu, I. Akasaki, J. Cryst. Growth 99, 381 (1990).
[96] T. Detchprochm, K. Hiramatsu, H. Amano, I. Akasaki, J. Appl. Phys. 61, 2688 (1992).
[97] R.J. Molnar, W. Götz, L.T. Romano, N.M. Johnson, J. Cryst. Growth 178, 147-156 (1997).
[98] H Lee, M Yuri, T Ueda, JS Harris, K Sin, J. Electron. Mater. 26, 898-902 (1997).
[99] T. Paskova, J. Birch, S. Tungasmita, R. Becard, M. Heuken, E. B. Svedberg, P. Runesson, E. M. Goldys, B. Monemar, Phys. Stat. Sol. A 176, 415 (1999).
[100] V. Wagner, O. Parillaud, H. J. Buhlmann, M. Ilegems, Phys. Stat. Sol. A 176, 429 (1999).
[101] P. R. Tavernier, E. V. Etzkorn, Y. Wang, D. R. Clarke, Appl. Phys. Lett. 77, 1804 (2000).
[102] S. Gu, R. Zhang, Y. Shi, Y. Zheng, L. Zhang, F. Dwikusuma, T. F. Kuech, J. Cryst. Growth 231, 342 (2001).
[103] M. Namerikawa, T. Sato, O. Takahashi, T. Suemasu, F. Hasegawa, J. Cryst. Growth 237-239, 1089 (2002).
[104] H. Murakami, Y. Kumagai, H. Seki, A. Koukitsu, J. Cryst. Growth 247, 245 (2003).
[105] Y. Golan, X. H. Wu, J. S. Speck, R. P. Vaudo, V. M. Phanse, Appl. Phys. Lett. 73, 3090 (1998).
[106] W. Zhang, T. Riemann, H. R. Alves, M. Heuken, D. Meister, W. Kriegseis, D. M. Hofmann, J. Christen, A. Krost, B. K. Meyer, J. Cryst. Growth 234, 616 (2002).
[107] T. Kawabata, T. Matsuda, S. Koide, J. Appl. Phys. 56, 2367 (1984).
[108] M. Hashimoto, H. Amano, N. Sawaki, I. Akasaki, J. Cryst. Growth 68, 163 (1984).
[109] C.-Y. Hwang, M. J. Schurman, W. E. Mayo , Y. Li, Y. Lu , H. Liu, T. Salagaj, R. A. Stall , J. Vac. Sci. Te