Thermal stability on the structural and optical properties of high indium content InGaN films grown using pulsed laser deposition (PLD) was investigated through long-duration and high-temperature annealing. X-ray diffraction and cathode- luminescence measurements of the 33% indium InGaN revealed no differences in the line-shape and peak position even after annealing at 800°C for 95 min; similar structural stability was found for the 60% samples after annealing for 75 min. The higher thermal stability is attributed to nanoscale InN domains with different orientations create mixed-polarity InGaN/InN interfaces, resulting in higher activation energies at interfaces and increasing the thermal stability of the material. Furthermore, the InGaN films were subjected to metalorganic chemical vapor deposition treatment to regrow a GaN layer; results are promising for the development of high thermal stability InGaN films using the PLD technique.
©2012 Optical Society of America
InGaN alloy has attracted significant interest of late due to its tunable bandgap energy ranging from the ultraviolet region (GaN: 3.4eV) to the near-infrared region (InN: 0.7eV) [1,2]. This broad range of possible bandgap covers almost the entire solar spectrum, giving the alloy significant potential for use in optoelectronic devices, especially light-emitting diodes (LEDs) and full-spectrum solar cells [3–5]. However, the phase separation of an InGaN alloy occurs readily during the spinodal decomposition process due to the large lattice mismatch between InN and GaN. This leads to a variability in the indium concentration and negatively impacts LED spectral quality [6,7]. Another issue is that the extremely high vapor pressure of InN prevents the incorporation of indium into the InGaN layer. Studies aimed at enhancing the indium incorporation using metalorganic chemical vapor deposition (MOCVD) and molecular beam epitaxy have highlighted several possible approaches, such as increasing the layer growth rate or the indium vapor flow, and decreasing the growth temperature [8–10]. While the optical and electrical characteristics of high indium InGaN materials have been discussed in these studies, the thermal properties of the films have received less attention. A specific issue is that in a typical LED structure, the InGaN layers comprising the active layer are subjected to the temperatures exceeding the growth temperature (600-670°C). One such example is in p-GaN layer growth (>800°C). This results in thermal damage to the InGaN material, causing structural and optical degradation of the LED. Therefore, the development of high indium InGaN alloys with high thermal stability is an important topic in LED research. To this end, we have applied low temperature pulsed laser deposition (PLD) to the growth of high indium InGaN films . PLD is a highly nonequilibrium evaporation process and can be operated at low temperature to increase indium incorporation [12,13]. In this study, a u-GaN layer was grown using MOCVD with sapphire as the growth substrate in order to improve the InGaN film quality for LED applications. The effect of heating on the structural and optical properties of the films was investigated using x-ray diffraction (XRD), atomic force microscopy (AFM), auger electron spectroscopy (AES), transmission electron microcopy (TEM), photoluminescence (PL), and cathodeluminescence (CL) measurement.
All InGaN films were grown on a 2-µm u-GaN template at 300°C in nitrogen plasma ambient atmosphere with PLD. A dual-compositing target was used which consisted of an indium sheet drilled with periodic rectangular holes mounted on a 3-inch GaN wafer. The percentage ratio of the total rectangle-hole area to the indium sheet area is defined as the F factor. Two targets with F factors of 0.682 and 0.364 were used in this study. A KrF excimer laser (λ = 248 nm) was employed as the ablation source, and operated with a repetition rate of 1 Hz and a pulse energy of 60 mJ. During the deposition process, the target was rotated at 25 rpm. Directing the ablation laser onto the two-sectioned target described above provided a co-deposition reaction, wherein InGaN grains formed on the u-GaN surface from indium and GaN vapors and acted as nucleation seeds to further promote InGaN growth. The as-deposited InGaN films were 140 and 64 nm thick using the F = 0.682 and 0.364 targets; these are denoted sample A0 and B0, respectively. Although the use of dual-compositing target allows the fabrication of high indium content InGaN films, some InN alloy was also created during the co-deposition process. To remove this impurity, the samples were post-annealed at 800°C for 15 min, which decomposed the InN alloy.
3. Results and discussion
The XRD curves of the as-deposited (A0 and B0) and annealed (A15 and B15; with subscript 15 denoting annealing for 15 min) samples are presented in Fig. 1 . The 33% indium composition for the InGaN sample shown in Fig. 1(a) was calculated by determining the shift of the (0002) InGaN diffraction peak relative to the GaN peak, and applying Vegard’s law. No changes in the XRD spectra were observed for the InGaN line-shapes and positions. The full width at half maximum (FWHM) of the InGaN peak was measured to be 0.29-0.30°. It is notable that a weak InN peak appeared at 31.33° in the as-deposited state and disappeared after annealing. An InGaN peak corresponding to the indium composition of 60% and the evolution of InN peak are similarly shown in Fig. 1(b). The base of the InGaN peak was broadened slightly with the disappearance of the InN alloy. This was attributed to the indium produced during the decomposition of InN reacted with GaN to form InGaN with a lower indium content . Figure 2 shows the XRD InGaN peak position and the root-mean-squared (RMS) surface roughness of series-A and -B as a function of the annealing time at an isothermal annealing temperature of 800°C. The InGaN peaks of series-A maintained an unchanged position, indicating good thermal stability even after long-duration and high-temperature annealing. The RMS roughness of series-A increased from 1.35 nm for as-deposited state, to 2.38, 3.42 and 5.39 nm for 800°C annealing for 15, 75 and 95 min, respectively. The surface morphology of A0 and A95 were presented in Fig. 3 . It clearly indicated that no significant indium segregation occurred in sample A95 even after high-temperature annealing. Contrast, the InGaN peak for series-B shifted slightly to a higher angle and the RMS roughness rose more appreciably from 7.24 to 14.48 nm as the annealing time was extended from 75 to 95 min. Moreover, as shown in the inset of Fig. 2, the InGaN intensity for B95 over B75 was attenuated 2.3 times more strongly than that of B75 over B15. This implies that when the annealing time exceeded 75 min, the 60% indium content InGaN film decomposed rapidly. The depth of the valley between InGaN and GaN peaks also began to increase due to the formation of InGaN with varying indium content. Additionally, the decomposition of InGaN was accompanied by an increase in the surface roughness.
AES depth profiles were employed to better understand the thermal behavior of InGaN, as illustrated in Fig. 4 . In Figs. 4(a)-4(c), the distribution of Ga, O and N in sample A0, A15, and A95 are seen to be stable, and appear unaffected by the increased annealing time. The relatively high indium buildup above the InGaN/GaN interface is of interest. Due to the low melting point of indium, the indium vapor generated in the ablation process was available to react with the background nitrogen plasma, forming InN alloy on the u-GaN surface. As the annealing time increased in excess of 15 min, the InN dissociation caused a small variation in the indium concentration near the InGaN/GaN interface. After annealing for 95 min, all elemental distributions returned to their steady state. In contrast to series-A, some differences in indium concentration of series-B were found, as shown in Figs. 4(d)-4(f). The indium concentration remained almost unchanged for annealing times up to 75 min. However, the change between 75 and 95 min exhibits diffusion into the u-GaN layer. On the other hand, the 60% indium InGaN samples were post-annealed and the disappearance of InN was confirmed using XRD; however, XRD is not sensitive enough to determine the composition of nanoscale structures. Several nanoscale InN grains were discovered in sample B75 even after the annealing for 75 min. These nanoscale InN grains may play an important role in maintaining structural stability because of the larger activation energy for the intermixing reactions between InN and GaN . The phenomenon suppressed indium diffusion into GaN, which was evidenced by the distinguished InGaN/GaN interface shown in Fig. 4(e). The thermal stability of InGaN therefore depends on the amount of nanoscale InN. When the annealing time increased, InN increasingly decomposed into indium and nitrogen, resulting in indium inter-diffusion and nitrogen out-gassing (Fig. 4(f)). This explains the change in the RMS roughness for sample B95. The detailed effects of nanoscale InN are discussed below with reference to TEM images.
High-resolution cross-sectional TEM images of series-B are shown in Fig. 5 . The interface between InGaN and GaN is identified clearly and the d-spacing values of series-B were calculated to be 2.75 Å, corresponding to an indium content of 60%. Several regions with a different d-spacing value of 2.85 Å were found which are indexed to the InN (0002) plane. These are indicated by a red circle in Figs. 5(b) and 5(c). The existence of nanoscale InN alloy indicates that the post-annealing time was too short to fully decompose InN. The nanoscale InN gradually diminished in amount and size as the annealing time increased. Furthermore, Different orientations of nanoscale InN was observed in sample B15 (Fig. 5(b)). This was considered to be evidence for InN growth trending towards a “deposition” rather than a “growth” mechanism . The varying InN orientations embedded within InGaN creates mixture polarity interfaces between InGaN and InN and causes different decomposition rates for each InN domain . It enhances the activation energy of interface and maintains the structural stability. Particularly, a number of stacking faults were found near the interface between InN domain and InGaN in Fig. 5. It may be attributed to the heat energy from thermal annealing was absorbed by InGaN to form stacking faults near the InN/InGaN interface before the occurrence of InN decomposition due to the higher stacking fault energy of InN [17–19]. The formation of stacking faults contributed indirectly to the increase in the stability ability of thermal treatment.
Figure 6 shows the PL spectra of samples B15 and B95 using a 405 nm laser excitation source. Two closely spaced peaks are visible at 855 and 862 nm in sample B15. The indium content can be determined from the PL peak position using the following equation:20] and is consistent with the results above. After long-duration and high-temperature annealing, a slight variation in the indium composition caused the two peaks to be shifted to 852 and 865 nm. However, the use of the PL signal to confirm the indium composition of samples A15 and A95 is difficult as their PL peak is located in the yellow band. This may be confused with the commonly observed yellow luminescence. This yellow luminescence is considered to be a result of defects in the GaN layer. CL spectra at 3 kV were recorded in order to verify the indium composition of both samples as shown in the inset of Fig. 6. The 88 nm electron penetration depth corresponded to 3 kV  and confirmed that the CL signals originated from InGaN layer rather than u-GaN. To summarize, the same CL InGaN peak with indium composition of 33% was visible at 554 nm for samples A15 and A95. As expected, these consistently exhibited high thermal stability. Despite the slight composition variation observed in samples B15 and B95, the observed thermal stability is still superior to that demonstrated in other studies of low indium composition InGaN alloys. Based on the earlier discussion in this paper, the 60% indium InGaN film was subjected to MOCVD to regrow a GaN layer and demonstrate the thermal stability. To avoid etching of the InGaN film by H2 introduced during the regrowth process , a 30 nm-thick GaN capping layer was grown on the InGaN film by PLD. Figure 7 shows the PL spectra of the GaN/InGaN/GaN structure before and after GaN regrowth at 880°C for 20 min. The peak position of 860 nm remained unchanged after MOCVD regrowth. The flat and uniform nature of the regrown sample shown in the inset of Fig. 6 indicates that the PLD method used in this study is indeed promising for the development of high indium content InGaN films with a high thermal stability.
In conclusion, the high thermal stability of InGaN films with indium contents of 33 and 60% was demonstrated through the measurement of structural and optical characteristics following long-duration high-temperature annealing. The existence of nanoscale InN causes mixed polarity interfaces, increasing the activation energy of the interface and further increasing the thermal stability. As a proof of concept, an InGaN film was treated with MOCVD to regrow GaN layer. The results clearly show promise for the development of high thermal stability InGaN films using the PLD technique.
This research was supported by National Science Council, The Republic of China, under the Contract No. 98-2221-E-005-006-MY3 and 101-2221-E-005-023-MY3.
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