InGaN Multiple Quantum Wells

ARTICLE pubs.acs.org/crystal

Growth and Characterization of GaN/InGaN Multiple Quantum Wells on Nanoscale Epitaxial Lateral Overgrown Layers W. K. Fong, K. K. Leung, and C. Surya* Department of Electronic and Information Engineering and Photonics Research Centre, The Hong Kong Polytechnic University, Hong Kong, People's Republic of China ABSTRACT: GaN/InGaN multiple quantum wells (MQWs) were fabricated on nanoscale epitaxial lateral overgrown (NELO) GaN layers (type N) by metal
organic chemical vapor deposition using a SiO2 layer with nanometer scale windows as the growth mask. Transmission electron microscopy (TEM) results clearly demonstrate coherent growth of GaN in the window regions while lateral growth is observed over the SiO2 layer. Based on the TEM and atomic force microscopy measurements, we observed substantial reduction in the threading dislocation density for the type N GaN films. Experimental results on electroluminescence (EL) measurement indicate substantial improvement in the EL intensity as well as a 15 nm blue shift in the EL peak wavelength. High resolution X-ray diffraction and reciprocal space mapping characterizations clearly indicate significant reduction in strain generation in the MQWs grown on NELO GaN layers compared to the control samples (type C). Such reduction in strain generation in the MQW gives rise to the reduction in the quantum-confined Stark effect. This is also consistent with the observed blue-shift in the EL peak. Detailed analyses of the optoelectronic properties of the devices indicate significant improvements in the internal quantum efficiency, ηi, and the extraction efficiency, ηe, by as much as 55.8 and 57.3%, respectively, compared to the type C devices. The improvement in ηi is attributed to reductions in both the nonradiative recombination centers and the quantum-confined Stark effect in the type N devices and the improvement in the extraction efficiency is attributed to the texturing of the GaN layer due to the incorporation of the SiO2 layer with nanometer scale windows.

’ INTRODUCTION III-nitrides are the materials of choice for the development of high-power light-emitting diodes (LEDs).1,2 To date, c-plane sapphire remains to be the most commonly used substrate for the manufacturing of GaN-based LEDs. Because of the large mismatches in the lattice constants and the thermal expansion coefficients between GaN and the sapphire substrate, the growth of high quality GaN/InGaN multiple quantum wells (MQWs) is still a critical issue for the development of high-brightness LEDs.3,4 The observed optoelectronic properties of the GaN/ InGaN MQWs are attributed to miscibility between GaN and InN, resulting in the so-called exciton-localization effect.5 It is generally accepted that the localization effects arising from inhomogenous indium distribution play an important role in the spontaneous emission for InGaN/GaN multiple quantum well structure.6
10 The In-rich regions in the MQWs are typically of the order to nanometers. Such regions, acting as quantum dots, facilitate the localization of excitons and are believed to be responsible for the high efficiencies in GaN-based LEDs. The efficiency of the device can be significantly affected by the presence of nonradiative recombination centers in the material. Epitaxial GaN layers on sapphire wafers typically have threading dislocation (TD) r 2011 American Chemical Society

density between 109 and 1011 cm
2.11,12 It has been demonstrated that the use of a two-step growth technology,13,14 with a lowtemperature buffer layer inserted between the sapphire substrate and the epitaxial GaN layer, can reduce the TD density to low 108 cm
2, but this is still significantly higher than the typical values of homoepitaxially grown GaAs which is between 102 and 104 cm
2.15,16 It has been shown that TDs have significant impacts on the performances of the devices;17,18 thus, it is important to investigate novel growth techniques for further reduction in the TD density. A substantial amount of work had been performed on one-step epitaxial lateral overgrown (1S-ELO) and two-step epitaxial lateral overgrown (2S-ELO)19
22 processes for the reduction of the dislocation density in GaN thin films. For the 1S-ELO process, it has been shown that the use of a SiO2 growth mask can effectively terminate the TD from propagating to the film surface thereby significantly reducing the TD density. However, the TD density in the coherent GaN in the window regions is still the Received: September 3, 2010 Revised: March 8, 2011 Published: March 29, 2011 2091

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Figure 1. SEM picture of the Ni nanoclusters on the SiO2 film.

same as the underlying GaN template. Since the typical width of the conventional dielectric growth mask and the window are 7 and 4 μm, respectively, an LED die will cover the highly defective coherent GaN regions that may significantly reduce the lifetimes of the devices. For the 2S-ELO process, the dislocation density can be further reduced and limited to the coalescence edge. However, the drawback is that relatively thick ELO-GaN, between 8 and 12 μm depending on the width of dielectric growth mask, is needed for full coalescence. This leads to significant bowing of the wafer and hence high nonuniformity in the material quality over the 2-in. wafer. Recently, the nanoscale epitaxial lateral overgrown (NELO)23
26 technique was applied directly on the sapphire substrates for the growth of GaN-based devices. In this paper, we report on the development of a novel LED structure in which a 5-period InGaN/GaN MQW was grown on top of a novel NELO GaN layer. Detailed investigations on the structural and optoelectronic properties of the devices were performed to identify the mechanisms underlying the improvements in the device properties. The NELO technique may also find applications in nonlinear optics and photovoltaics due to the improvement in the material quality of the NELO GaN layer and the reduction in the material stress generation. In particular, regarding the potential applications of III-nitrides in photovoltaics, the wide coverage of the III-nitride family in the optical spectrum, from 0.6 eV for InN to 3.4 eV for GaN, facilitates the development of all-nitride fullspectrum photovoltaic cells. This will have a significant implication on the cost of the device as the complete photovoltaic cell can be fabricated within one single metal
organic chemical vapor deposition (MOCVD) growth run. Another important advantage is the superior radiation hardness of the III-nitride materials. This will have significant implications on the lifetimes of the devices as well as the application of the device in space.

’ DEVICE FABRICATION An undoped GaN layer of thickness 2 μm was first deposited on a sapphire substrate in a Thomas Swan close-coupled showerhead MOCVD reactor. To grow the NELO GaN layer, a SiO2 layer of thickness 200 nm was deposited on the GaN template by plasma enhanced chemical vapor deposition technique. A 10-nm thick metallic Ni layer was deposited on top of the SiO2 layer by e-beam evaporation. This was followed by thermal annealing at 800 °C for 5 min in nitrogen ambient resulting in the formation of nanometer scale Ni clusters which were used as the etch mask for the etching of the nanometer scale windows in the SiO2 by

Figure 2. (a) TEM image of a Type N GaN thin film indicating nanocluster SiO2 islands, the termination of a dislocation by the SiO2 island and the propagation of dislocations in the GaN layer. (b) TEM image of a horizontal dislocation and a dislocation loop.

reactive ion etching (RIE). A scanning electron microscopy (SEM) picture of the Ni nanoclusters is shown in Figure 1. It is seen that the size and the separations of the Ni clusters are typically in the range of 100 nm to about 500 nm. CHF3 was used as the process gas in the RIE system at a flow rate of 15 sccm. Forward RF power used was 200 W and the chamber pressure was maintained at 10 mTorr during the etching process. This can produce nanometer scale SiO2 islands of average diameter and interdistance of 300 and 200 nm respectively. The Ni clusters were then removed by dipping the sample in hot HNO3 acid for 5 min. The sample was then loaded back into the MOCVD reactor to proceed with the NELO process. Undoped GaN was first deposited during the initial stage of the NELO process until full coalescence was achieved. Then, n-GaN was grown using liquid ditertiarybutyl silane (DTBSi) metal
organic dopant source with carrier concentration of the order 5  1018 cm
3. The active region of the device consists of 5-period InGaN/GaN MQWs, where the InGaN QWs were grown at 700 °C and the GaN barriers at 850 °C. Finally, a 150 nm thick p-GaN layer was grown at 1010 °C on top of the MQWs. Postgrowth annealing was done at 800 °C for 30 min under N2 ambient to activate the 2092

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Figure 4. HRXRD ω-scan of the (102) reflection of (a) type N (O) and (b) type C MQWs (0).

Figure 3. Top view AFM picture of (a) type C GaN film and (b) type N GaN thin film.

Mg dopants and hole concentration of 2  1017 cm
3 can be obtained. The GaN-based LEDs with (type N) and without (type C) the NELO structure were fabricated by conventional techniques. High transmittance Ni/Au (5 nm/7 nm) and Cr/Pd (20 nm/40 nm) layers were deposited for the formation of the p- and n-type ohmic contacts followed by the respective contact annealing processes. Around 500 nm thick Au bonding pads were fabricated on the p- and n-contacts to facilitate electrical connection.

’ EXPERIMENTAL RESULTS AND DISCUSSION The device structure above is designed specifically for the reduction of TDs in the overgrown GaN layer. Several studies27,28 demonstrated that high concentrations of nonradiative recombination centers are associated with the TDs. Furthermore, it has been shown that the lifetimes of the LEDs are strongly affected by the presence of TDs.29 Thus, reduction in the TD density will potentially result in significant improvements in the internal quantum efficiency and reliability of the device. The TDs in the GaN films are investigated by transmission electron microscopy (TEM) technique. Thin foil specimens for cross-sectional TEM measurement were prepared by standard procedure including polishing and ion milling. The TEM images were observed by a Jole-2010 high resolution transmission electron microscopy (HRTEM) operated at 200 kV. In Figure 2a, one observes the termination of a TD by the SiO2 growth mask. Also, it is interesting to note that the TEM pictures indicate that some of the TDs are located directly above the SiO2 growth mask. It has been shown that the wing regions of the

ELO-GaN coalesce with each other above the SiO2 growth mask leaving a TD at the coalescence edge.30 It is also shown that stress relaxation in the GaN layer may take place when this happens. In the middle of Figure 2b, a horizontal dislocation (HD) is seen to associate with a dislocation loop across two adjacent nano-SiO2 islands. It has been pointed out that such bending of the TD resulting in the formation of a dislocation loop may arise from a change from the 3D to 2D growth mode.31 The three-dimensional (3D) growth of GaN usually generate a lot of HDs because of the introduction of internal stresses during the growth of the vertical side-wall facets,32,33 which is consistent with our observation that the HDs are limited to the lower part of the overgrown layer. The TEM results clearly demonstrate the mechanisms that lead to the reduction in the TD density in GaN films utilizing the NELO technique. The estimated TD density based on the TEM results is ∼7.5  107 cm
2. The surface morphology of the GaN layer is characterized by atomic force microscopy (AFM) measurements over an area of 5 μm  5 μm on both type N and type C films. The results are indicated in Figure 3a and b, in which the TD densities are found to be about 6  107 cm
2 and 3.1  108 cm
2 respectively. The rms roughness of the type N and type C GaN films are found to be 0.138 and 0.238 nm respectively over a 2 μm  2 μm area. The results demonstrate excellent morphology for the type N films. The AFM figures may not represent the actual dislocation densities in the samples. However, the estimated TD densities from the AFM pictures serve to provide a comparative indication of the material quality between two different types of samples, which has been widely adopted for nitride material.34,35 Moreover, the AFM results are compared to the TEM and X-ray diffraction (XRD) results, which can provide detail information of the dislocation density. The ω-scan full width at halfmaximum (FWHM) of the symmetric 102 reflections can be used to provide structural properties of the samples, which is sensitive to (aþc)- and a-type dislocations.36 Our high resolution X-ray diffraction (HRXRD) results in Figure 4 show that the FWHMs of the 102 reflection for type N and C samples were 335 and 389 arcsecond respectively, which indicated that the crystallinity of the type N samples is better than the type C samples. Detailed characterizations of the optoelectronic properties of the devices were performed. Photoluminescence (PL) measurements were conducted over a wide range of temperatures to examine the internal quantum efficiency of the devices. A 15 mW 2093

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Figure 5. PL spectra for: a typical type N LED at room temperature (I) and 15 K (III); a typical type C MQW at room temperature (II) and 15 K (IV).

HeCd cw laser (λ = 325 nm) was used as the excitation source. The samples were placed inside a closed-cycle refrigerator and the temperature was controlled by a Lakeshore temperature controller. The PL emission was collected by fused silica lens and focused on an Oriel double monochromator. The PL signal was detected by a Hamamatsu photomultiplier tube and measured using a lock-in amplifier. The PL spectra of the devices are shown in Figure 5. Significant improvement in the PL intensity is observed in the type N devices compared to the control. The room temperature data demonstrate an improvement in the quantum efficiency, ηq, by 169% for the type N device. It is noted that while the PL spectrum for the type N device measured at 15 K is still higher than that of the type C device the difference is significantly less than the room temperature results. The experimental data strongly indicate significant effects of nonradiative recombination centers on the PL spectra. The emission peak around 390 nm at 15 K from both samples comes from the conduction band-to-acceptor transition, which is associated with shallow Mg impurities.37
39 To investigate the internal quantum efficiency, ηi, of the devices we have conducted a systematic study on the temperature dependencies of the PL. On the basis of the model proposed by A. Sasaki et al.,40 ηi is given by ηi ðTÞ ¼

IR ðTÞ 1=τR ðTÞ ¼ I0 1=τðTÞ

ð1Þ

where IR is the integrated radiative intensity and I0 is the excitation power, τR is the radiative recombination lifetime and τ is the overall carrier recombination lifetime which is related to τR and the nonradiative recombination lifetime, τNR, by 1/τ(T) = 1/τR(T) þ 1/τNR(T). To determine τ(T), the radiative recombination process dominates at low temperature while the nonradiative recombination processes may be significant at room temperature. For simplicity one can assume
 T and τR ðTÞ ¼ τR ð0Þexp
TR

ð2Þ


 T τNR ðTÞ ¼ τNR ð0Þexp
TNR

ð3Þ

Figure 6. Arrhenius plot of the integrated PL spectra for a typical type N MQW (b) and type C MQW (9). The solid lines represent the fitting of the experimental data with eq 4.

Table 1. Low-Temperature PL Fitting Result Using the Levenberg
Marquardt Approximation Technique normalized

normalized

ηi (T = 0 K)

ηi (T = 300 K)

sample

η0

T0

Type N

3.71

89.24

1

1

Type C

4.72

85.44

0.82

0.585

giving ηi ðTÞ ¼

IR ðTÞ 1 ¼ I0 1 þ η0 expðT=T0 Þ

ð4Þ

Here η0 t τR(0)/τNR(0) and T0 is the characteristic temperature. The values of η0 and T0 can be obtained by fitting the experimental data to eq 4 using the Levenberg
Marquardt approximation technique as shown in Figure 6 and the fitted results are summarized in Table (1). The experimental data in Figure 6 were obtained from the integrated intensity of the MQW emission peak. It is noteworthy that the values of η0 and T0 obtained using this technique are relatively insensitive to the excitation power used. Hence, it would be more meaningful to compare ηi(0) for different samples. Equation 4 had been used in the characterization of the temperature dependences of the PL spectra in various disordered superlattices where η0 is a proportionality constant which is material dependent and T0 is used as a thermal quenching parameter for the PL intensity and is speculated to correlate to the energy depth of the localized states. A large T0 corresponds to a small temperature dependence for ηi(T). A large η0 indicates the domination of nonradiative recombination in a sample with high defect density. Thus, comparing the values of η0 for different materials facilitates the evaluation of material quality of the samples. Substituting the values for η0 and T0 from Table 1 into eq 4 we obtained a 71% increase in the internal quantum efficiency, Δηi(T), for type N device compared to the conventional type C device at room temperature. From the room temperature PL spectra the total improvement in the quantum efficiency, Δηq, for type N device is found to be 169%. Utilizing the relationship, ηq = ηi  ηe, where ηe is the extraction efficiency, we obtain Δηe = 57.3%. The improvement in ηe is believed to arise from the introduction of nanometer scale SiO2 islands resulting in the texturing of the NELO GaN layer. 2094

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Figure 7. Experimental data on EL spectra measured from typical type N (b) and C (9) devices at 210 mA.

Figure 9. HRRSM of the (105) diffraction of (a) type N and (b) type C MQWs.

Figure 8. HRXRD omega scans for the (002) reflection of (a) experimental data (b) and simulated results (O) for a typical type N device and (b) experimental data (9) and simulated results (0) for a typical type C device.

It is important to note that ηi(T) evaluated based on the low temperature PL results may differ substantially from that evaluated based on carrier injection experiments. This is because ηi(T) is found to be bias dependent. Under large bias conditions the localized states in the MQWs will be saturated resulting in the

increase in ηi(T). Furthermore, carrier motion in the MQW under a large voltage bias will be strongly influenced by the electric field imposed on the device. Hence, ηi(T) evaluated by fitting eq 4 to the low temperature PL data, cannot be compared directly to the EL results. To determine the improvement in ηi(T,VA) for type N devices under voltage bias, VA, we assume that ηe(T) is independent of bias, thus the value of ηe(T) obtained from the PL data can be applied directly to the EL spectra for the evaluation of Δηi(T,VA). The experimental results for the room temperature EL under a bias VA are shown in Figure 7 from which Δηq(T,VA) for the type N device is found to be ∼145% and Δηi(T,VA) = 55.8%. For the electroluminescence (EL) measurements, the EL spectra were measured using a Labsphere LMS-100 integrating sphere. It is interesting to note that a 15 nm blue-shift in the EL peak is observed for type N devices. It is not likely that this originates from random device-to-device variation of the EL peak because both types of devices were grown within the same MOCVD run and the random variation in the EL is typically within a range of 2
3 nm. It has been suggested that such significant blue shifts in the EL peak may originate from reduction in strain generation in the MQWs leading to a substantial reduction in the quantum-confined Stark effect.41,42 The amount of blue shifts in the peak positions between the type N and type C samples differ substantially as observed in the EL and PL data. This is attributed to the run-to-run variation as the samples used in the two experiments were deposited in different MOCVD growth runs that are well separated apart which may result in large variations in the peak positions of the devices. HRXRD and 2D-reciprocal space mapping (2D-RSM) were used for the characterization of strain in the MQWs. Detailed experimental procedures can be found in a previous publication.43 2095

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Crystal Growth & Design Figure 8a and b illustrate the XRD patterns for (002) reflection of a 5-period InxGa1-xN/GaN QW together with a simulated curve for both types N and C samples. The simulations were performed using the LEPTOS software based on the Recursive Matrix Formalism. From the simulation results, the In concentration in the QWs is estimated to be 15.27%. From the asymmetric (105) HRRSM shown in Figure 9a and b, it is observed that the type C sample exhibits a wider spread in the diffraction intensity compared to the type N sample indicating that the NELO sample has a much better orientation distribution and crystallinity.44 Considering the reciprocal lattice point (RLP) system, a parameter θ is used for the characterization of lattice relaxation, which can be expressed as a function of the elastic parameters45 " # c33 ðxÞ c 4ðh2 þ k2 Þ   θðxÞ ¼ arctan ð5Þ 2c13 ðxÞ a l2 where c(x)13 and c(x)33 are the elastic constants in which c(x)13 = 103 and 92 GPa and c(x)33 = 405 and 224 GPa for GaN and InN respectively. The values for lattice constants a and c for GaN and InN are aGaN = 3.1892 Å, aInN = 3.5378 Å, cGaN = 5.1850 Å and cInN = 5.7033 Å respectively.46
48 The values of θ, obtained using eq 5 above and based on x = 15.27% for (105) reflections, are found to be 19.2° and 0° for the fully relaxed and fully strained InxGa1-xN layer respectively. Figure 9 shows that the experimental values for θ are 3.8° and 0° for the type N and type C samples. The HRRSM results are consistent with the EL results indicating substantial reduction in strain generation in the type N sample. Reduction in strain generation in the MQW stipulates corresponding reduction in the quantum-confined Stark effect since III-nitrides are piezoelectric material. This is consistent with the observed 15 nm blue shift among the type N samples. An important consequence in the reduction of the quantum-confined Stark effect is the increase in the overlap in the electron and hole wave functions resulting in an increase in the internal quantum efficiency. Thus, the experimental data show that the observed improvements in ηi(T) for the type N devices arise from (i) reduction in the density of recombination centers in the MQWs; and (ii) reduction of strain in the MQWs resulting in the reduction in the quantum-confined Stark effect in the MQWs.

’ CONCLUSIONS We have investigated the incorporation of a NELO layer in the GaN LED structure. AFM characterizations of the surface morphology of the sample indicate significant reduction in the dislocation density. From the TEM pictures we observed termination of the TDs by the SiO2 growth mask as well as the formation of dislocation loops which are being attributed as the mechanisms responsible for the reduction of the TD density. Detailed investigations of the optoelectronic properties of the devices demonstrate that the utilization of a NELO layer leads to significant improvements in both the extraction and internal quantum efficiencies of the LEDs. High resolution reciprocal space mapping indicate substantial reduction in strain generation in the MQWs when a NELO layer is incorporated in the device structure. Based on the experimental results, the improvements in the internal quantum efficiency are attributed to two different mechanisms: the reduction in the defect density in the MQW resulting in the corresponding reduction in both the nonradiative recombination process and reduction in strain generation in the

ARTICLE

MQWs leading to the reduction in the quantum-confined Stark effect.

’ AUTHOR INFORMATION Corresponding Author

*E-mail: [email protected].

’ ACKNOWLEDGMENT This work is supported by a grant from the Innovative Technology Commission under the Guangdong/Hong Kong Scheme (Project no. GHP/031/07GD) and a GRF Grant (PolyU 5269/07E). Further support is provided by a Niche area grant of the Hong Kong Polytechnic University. ’ REFERENCES (1) Cao, X. A.; Arthur, S. D. App. Phys. Lett. 2004, 85, 3971. (2) Wetzel, C.; Salagai, T.; Detchprohm, T.; Li, P.; Nelson, J. S. App. Phys. Lett. 2004, 85, 866. (3) Yamaquchi, S.; Kariya, M.; Nitta, S.; Takeuchi, T.; Wetzel, C.; Amano, H.; Akasaki, I. J. Appl. Phys. 1999, 85, 7682. (4) Song, T. L.; Chua, S. J.; Fitzgerald, E. A.; Chen, P.; Tripathy, S. Appl. Phys. Lett. 2003, 83, 1545. (5) Bai, J.; Wang, T.; Sakai, S. J. Appl. Phys. 2001, 90, 1740. (6) Narukawa, Y.; Kawakami, Y.; Funato, M.; Fujita, S.; Nakamura, S. Phys. Rev. B 1997, 55, R1938. (7) Statake, A.; Masmoto, Y.; Miyajima, T.; Asatsuma, T.; Nakamura, F.; Ikeda, M. Phys. Rev. B 1998, 57, R2041. (8) Sing, R.; Doppalapudi, D.; Moustakas, T.; Romano, L. Appl. Phys. Lett. 1997, 70, 1089. (9) El-Masry, N. A.; Piner, E. L.; Liu, S. X.; Bedair, S. M. Appl. Phys. Lett. 1998, 72, 40. (10) Sun, C.; Zubair Anwar, M.; Chen, Q.; Yang, J.; Asif Khan, M.; Shur, M.; Bykhovski, A.; Liliental-weber, Z.; Kisielowski, C.; Smith, M.; Lin, J.; Xiang, H. Appl. Phys. Lett. 1997, 70, 2978. (11) Kapolnek, D.; Wu, X. H.; Heying, B.; Keller, S.; Keller, B. P.; Mishra, U. K.; DenBaars, S. P.; Speck, J. S. Appl. Phys. Lett. 1995, 67, 1541. (12) Lester, S. D.; Ponce, F. A.; Craford, M. G.; Steigerwald, D. A. Appl. Phys. Lett. 1995, 66, 1249. (13) Beaumont, B.; Bousquet, V.; Vennegues, P.; Vaille, M.; Bouille, A.; Gibart, P.; Dassonneville, S.; Amokrane, A.; Sieber, B. Phys. Status Solidi A 1999, 176, 567. (14) Vennegues, P.; Beaumont, B.; Bousquet, V.; Vaille, M.; Gibart, P. J. Appl. Phys. 2000, 87, 4175. (15) Fischer, R.; Neuman, D.; Zabel, H.; Morkoc, H.; Choi, C.; Otsuka, N. Appl. Phys. Lett. 1986, 48, 1223. (16) Bedair, S. M.; Humphreys, T. P.; Ei-Masry, N. A.; Lo, Y.; Hamaguchi, N.; Lamp, C. D.; Tuttle, A. A.; Dreifus, D. L.; Russell, P. Appl. Phys. Lett. 1986, 49, 942. (17) Detchprohm, T.; Xia, Y.; Xi, Y.; Zhu, M.; Zhao, W.; Li, Y.; Schubert, E. F.; Liu, L.; Tsvetkov, D.; Hanser, D.; Wetzel, C. J. Cryst. Growth 2007, 298, 272. (18) Ling, S. C.; Wang, T. C.; Chen, J. R.; Liu, P. C.; Ko, T. S.; Lu, T. C.; Kuo, H. C.; Wang, S. C.; Tsay, J. D. Jpn. J. Appl. Phys. 2009, 48, 04C136. (19) Gibart, P. Rep. Prog. Phys. 2004, 67, 667. (20) Ko, C. H.; Su, Y. K.; Chang, S. J.; Tsai, T. Y.; Kuan, T. M.; Lan, W. H.; Lin, J. C.; Lin, W. J.; Cherng, Y. T.; Webb, J. B. Mater. Chem. Phys. 2003, 82, 55. (21) Ni, X.; Ozgur, U.; Morkoc, H.; Liliental-Weber, Z.; Everitt, H. O. J. Appl. Phys. 2007, 102, 053506. (22) Yu, Z.; Johnson, M. A. L.; Brown, J. D.; Ei-Masry, N. A.; Cook, J. W.; Schetzina, J. F. J. Cryst. Growth 1998, 195, 333. 2096

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