Firstly, GaN thin films were grown on several substrates, that is, a synthetic quartz glass, an amorphous-carbon- coated graphite, and a c-face sapphire single-crystalline wafer as a reference. The films were directly grown on the each substrates at the growth temperature of 900˚C, and their crystallinities and PL properties were compared. XRD patterns of the films are shown in Figure 1, respectively. In the case of the GaN thin film grown on the quartz glass substrate, only (0002) and (0004) GaN diffraction peaks relating the wurtzite-type crystal structure are observed in the XRD pattern, and it is similarly to the case of the film grown on the c-face sapphire substrate as shown in Figure 1(a) and Figure 1(b), respectively. Therefore, the thin film grown on the quartz glass substrate is at least highly c-axis preferential orientated GaN, though there was a clear difference in their crystallinities. That is, full-widths at the half-maximum (FWHMs) of the X-ray rocking curves (XRCs) of the (0002) GaN diffractions of the films were largely different as the cases on the quartz glass was 7.6˚ and on the sapphire was 0.7˚.

The XRD pattern of the GaN thin film grown on the amorphous-carbon-coated graphite substrate is shown in Figure 1(c). The base material of the substrate is a polycrystalline graphite, and an amorphous-carbon layer of several μm thickness is coated on the graphite base. In the XRD pattern, the diffraction peaks of the substrate are indicated by asterisks. Relatively narrow and broad peaks are simultaneously observed, and they respectively correspond to the peaks of the polycrystalline graphite and amorphous-carbon layer. From the consideration,

it can be said that the film grown on the amorphous-carbon-coated graphite substrate is also highly c-axis oriented GaN since the only (0002) and (0004) GaN diffraction peaks are observed except the peaks caused from the substrate material in the XRD pattern.

Nitride semiconductor thin films often become poor polycrystalline when they are grown on polycrystalline substrates because of random crystal orientations in the polycrystalline substrates [21] . On the other hand, the poor polycrystalline state is well improved and highly c-axis oriented nitride thin films are obtained when an amorphous layer having around 1 μm thickness is inserted between the nitride thin films and such polycrystalline substrates [22] [25] . Therefore, the similar effect on the crystallinity is considered to be occurred in the case of the GaN thin film grown on the amorphous-carbon-coated graphite substrate.

Figure 2 shows PL spectra of the GaN thin films grown on the same substrates described above. Band-edge emission at the higher photon energy and deep-level emission at the lower photon energy are respectively observed in the PL spectrum of the GaN thin film grown on the quartz glass substrate as shown in Figure 2(b). This is similar to the case of the film grown on the sapphire substrate shown in Figure 2(a) though the emission intensity of the band-edge emission at the higher energy side is weaker than that of the film grown on the sapphire substrate. On the other hand, broad emission peak at around 3 eV is observed in addition to the band-edge emission at the higher energy side in the PL spectrum of the GaN thin film grown on the amorphous-carbon- coated graphite substrate. It is different from the other PL spectra of the GaN films grown on the quartz glass or sapphire substrates.

Next, impurity doping was attempted to GaN thin films grown on the non-single-crystalline substrates and their PL properties were investigated. Doping of impurities such as Si, Ge, and Mg to the epitaxial nitride semiconductor thin films is known as an effective way to enhance the luminescence intensities [28] [29] . Ge is an n-type dopant for the nitride semiconductors [29] , and its vapor pressure is higher than Si which is the most popular n-type dopant [28] for the nitride semiconductors. Therefore, the doping of Ge in the MBE apparatus, in which materials are thermally evaporated and supplied to synthesis nitride semiconductor thin films, is much easier than Si, and heavy doping of Ge to nitride thin films becomes possible [30] . From the reasons, doping of Ge was carried out on the GaN thin films grown on the non-single-crystalline substrates.

Figure 3 shows the PL spectra of the Ge-doped GaN thin films grown on the same substrates. The growth temperature of the each films was 900˚C, and Ge was supplied at the cell temperature of 1200˚C. Similarly

to the case in which the film was grown on the sapphire substrate as shown in Figure 3(a), the emission peaks which correspond to the near band-edge emission were drastically increased compared with those of the undoped GaN thin films as shown in Figure 2. The increases of the near band-edge emission intensities by the Ge doping are considered that the impurity levels have become effective luminescence centers replacing band-to- band transitions. In the case of the Ge-doped GaN thin film grown on the amorphous-carbon-coated graphite substrate, the near band-edge emission intensity also increased though the emission peak at the slightly lower energy side remained.

Highly c-axis oriented GaN thin films were obtained on the non-single-crystalline substrates and near band- edge emission intensities in their PL spectra were enhanced by Ge doping. Next, In_xGa_1−xN thin films of several compositions were grown on the substrates and their properties were investigated. Firstly, the In_xGa_1−xN thin films were grown on the quartz glass substrate. The growth temperature and the Ga cell temperature were fixed to 800˚C and 970˚C, respectively. On the other hand, the In cell temperature was varied from 760˚C to 780˚C, and as a result, In_0.1Ga_0.9N and In_0.16Ga_0.84N thin films were respectively obtained. The XRD patterns of the In_xGa_1−xN thin films are shown in Figure 4. Phase separation, which is often appeared in In_xGa_1−xN thin film, was not observed in the obtained thin films and each of them exhibited highly c-axis oriented crystal structures.

PL properties of the In_xGa_1−xN thin films are shown in Figure 5. The value of the energy band gap of In_xGa_1−xN varies as a function of the molar fraction x with a certain bowing. The values of the energy band gaps of the obtained In_xGa_1−xN thin films were respectively estimated as about 3.0 and 2.8 eV under a consideration of the bowing parameter. Near band-edge emissions at the higher energy sides were observed in the PL spectra, though other emission peaks, which were considered to correspond to deep-level emissions, were also observed at the lower energy sides. Similar tendencies were observed in the properties of the In_xGa_1−xN thin film directly grown on the amorphous-carbon-coated graphite substrate. The XRD pattern and PL spectrum of the In_0.1Ga_0.9N thin film grown on the amorphous-carbon-coated graphite substrate are shown in Figure 6, respectively. Only the diffraction peaks relating to the c-face of In_xGa_1−xN were observed in the XRD pattern except the diffraction peaks relating to the substrate used. That is, highly c-axis oriented thin film was obtained, and in addition, phase separation was not observed in this case. Near band-edge emission peak and an emission peak which was

considered to be caused from deep-level defects were observed in the PL spectrum. The PL peak intensity in the higher energy side is relatively higher than that of the In_0.1Ga_0.9N thin film grown on the quartz glass substrate. There might be a different mechanism connecting to the enhancement of the near band-edge emission in this case.

Ge doping was also attempted to the In_xGa_1−xN thin film grown on the quartz glass substrate. It had been confirmed that the Ge doping to the GaN thin films directly grown on the non-single-crystalline substrates brought increases of the near band-edge emission intensities in their PL spectra as described before. A similar result was obtained in the case of Ge-doping to the In_xGa_1−xN thin film as shown in Figure 7. The Ge cell temperature was 1100˚C in this case. Effective luminescence centers might be occur in the In_xGa_1−xN thin film by the doping similarly to the cases of the GaN thin films grown on the same substrate.

Highly c-axis oriented films were obtained and increases of the near band-edge emission intensities by Ge doping were confirmed about the thin films directly grown on the non-single-crystalline substrates as described above section. Some effects of Ge or Mg-doping to GaN thin films grown on non-single-crystalline substrates are described in this section. Firstly, effects of n-type Ge-doping were investigated on the resistivity of the GaN thin films. The Ge-doped GaN thin films were directly grown on the quartz glass substrate. The growth temperature was fixed to 900˚C, and the Ge cell temperature was varied from 1200˚C to 1300˚C. The variation of the resistivity is shown in Figure 8. The resistivity gradually decreased with the increase of the Ge cell temperature, that is, free electron density increased by the increase of the amount of the n-type impurity. The carrier density in the thin film was confirmed by a Hall effect measurement, and it reached to 3 × 10²⁰ cm⁻³ when Ge was doped to it at the cell temperature of 1300˚C. Thus, Ge doping is quite effective to obtain GaN thin films of high carrier density when the film is grown on the non-single-crystalline substrate similarly to the epitaxial nitride thin films.

Next, effects of Mg doping were investigated on controlling conduction types of GaN thin films grown on the

non-single-crystalline substrates. Effects of Mg-doping to GaN thin films grown on such non-single-crystalline substrates have not been well described, though Mg is the well-established p-type dopant on GaN epitaxial thin films [1] . The Mg-doped GaN thin films were respectively grown on an amorphous-carbon substrate of 0.4 mm thickness, and on a quartz glass substrate on which a heavily Ge-doped conductive GaN underlayer was grown. The growth temperature of the films was 900˚C, and Mg was doped to the films at the Mg cell temperature of 370˚C.

The conducting types of the obtained thin films were evaluated by the ECV measurements. KOH aquatic solution was used for fabrications of Schottky contacts to the Mg-doped GaN thin films. On the other hand, the amorphous-carbon substrate has relatively high conductivity and an ohmic relationship to the Mg doped GaN thin film. The heavily Ge-doped GaN thin film grown as the conductive underlayer on the quartz glass substrate also has a similar ohmic relationship to the Mg-doped GaN thin film. Therefore, the amorphous-carbon substrate and the heavily Ge-doped GaN underlayer were respectively used as ohmic contacts in the ECV measurements.

Figure 9 shows the Mott-Schottky plots about the Mg-doped GaN thin films obtained from the ECV measurements. The directions of the slopes in these plots are clearly different between the cases of the Mg-doped GaN thin films and the undoped n-type GaN thin film. Usually, undoped GaN shows n-type conduction because of naturally occurred nitrogen vacancies or excess metallic Ga. It can be determined that the undoped GaN thin film has n-type conduction property from the result of the ECV measurement as shown in the Figure 9(c). On the other hand, the measurement results on the Mg-doped thin films show clearly opposite direction of the slopes, that is, the slopes declined with the increase of the applied voltage. Therefore, it can be determined that the conduction type of the each Mg-doped film is p-type. The ECV measurements were carried out under applying alternating voltage of 69 Hz to avoid effects of the relatively high impedance of the KOH aquatic solution on the decrease of applied voltage to the GaN thin films. Therefore, estimated values of acceptor concentrations in the films might include the number of non-activated acceptors. Nevertheless, it was shown that p-type GaN thin films could be obtained on such non-single-crystalline substrates.