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High indium compositions InGaN films were grown on sapphires using low temperature pulse laser deposition (PLD) with a dual-compositing target. This target was used to overcome the obstacle in the InGaN growth by PLD due to the difficulty of target preparation, and provided a co-deposition reaction, where InGaN grains generated from the indium and GaN vapors deposit on sapphire surface and then act as nucleation seeds to promote further InGaN growth. The effects of co-deposition on growth mechanisms, surface morphology, and electrical properties of films were thoroughly investigated and the results clearly show promise for the development of high indium InGaN films using PLD technique with dual-compositing targets.
©2012 Optical Society of America
1. Introduction
InGaN alloys are of great fundamental and technological interest owing to their tunable bandgap energy - from the ultraviolet (GaN, 3.4eV) to the near infrared (InN, 0.7eV) [1–3]. This broad range spans nearly the entire solar spectrum, rendering the material highly suitable for use in optoelectronic devices such as light-emitting diodes and multi-junction solar cells [4–6]. However, due to a large lattice mismatch between InN and GaN, InGaN alloys with high indium content can readily undergo phase separation during the spinodal decomposition process. This leads a variable in indium concentration, which negatively impacts device performance and spectral quality [7]. Moreover, to avoid the deleterious effects of the extremely high vapor pressure of InN, some studies have suggested the use of increased growth rate and growth temperature reduction in order to increase the quantity of indium incorporated into InGaN films [8,9]. While high indium content InGaN alloys have been obtained using such methods, the alloys formed typically exhibit poor crystalline quality and an inhomogeneous distribution of indium. Therefore, preparing high quality InGaN films with sufficiently high indium content remains a significant challenge in the field. In this study, we employed low temperature pulsed laser deposition (LT PLD) to grow high indium InGaN films. PLD is a highly nonequilibrium evaporation process, where the stoichiometry of the deposited film is very closely matched to that of the target used [10]. The deposition of a number of nitride films such as GaN, AlN, and InN have been achieved by PLD [11–13]. However, the difficulty of target preparation is an obstacle to achieving high quality InGaN films by PLD. PLD targets are commonly fabricated by pressing alloy powders at high pressure, followed by sintering at high temperatures. When using an InGaN target by sintering InN and GaN powders to deposit InGaN films, a film of non-uniform composition will often be produced owing to the difference in melting points between the two powders. Recently, a 44% indium content InGaN film was deposited by employing an In0.65Ga0.35N eutectic target compounded by liquid gallium and solid indium [14]. Although such eutectic targets indeed allow for InGaN film growth, details regarding the growth mechanisms of the dual atom (Ga and In) reactions along with a discussion of the electrical properties of the deposited InGaN films are rarely provided. To this end, we have designed a target consisting of a 3-inch indium sheet drilled with periodic rectangular-holes mounted on a GaN wafer grown by Hydride Vapor Phase Epitaxy. This target allows us to directly study co-deposition behavior. By changing the ratio of hole area to total sheet area, the indium composition in the InGaN films could be modulated from 33% to 62%. The effects of co-deposition on the crystalline quality and electrical properties of the InGaN films fabricated using this target were thoroughly investigated by scanning electron microscopy (SEM), x-ray diffraction analysis (XRD), atomic force microscopy (AFM), and Hall measurements.
2. Experimental
All InGaN films were grown on c-plane sapphire substrates at 300°C in nitrogen plasma ambient by PLD, using the dual-compositing target described above. The percentage ratio of the total rectangular holes area to the indium sheet area is defined as the F factor. The thickness of four InGaN films using targets with F factors of 0.682, 0.602, 0.443, and 0.364 were about 150 nm thick (the average growth rate: ~28-31 nm/hr). To simplify the discussion, films deposited using targets with F = 0.682, 0.602, 0.443, and 0.364 are denoted as samples A, B, C and D, respectively. A KrF excimer laser (λ = 248 nm) operated with a repetition rate of 1 Hz and an energy fluence of 60 mJ/pulse was employed as the ablation source. The target was set at a fixed distance of 9 cm from the substrate, and was rotated at 30 rpm during film deposition. The working pressure was 1.13 × 10−4 Torr with the injection of N2 plasma. The input power for plasma generation was fixed at 600 W. In order to understand the InGaN film growth mechanism on sapphire, we closely monitored deposition from F = 0.682 target as a function of time.
3. Results and discussion
Figure 1 shows the surface morphologies of the InGaN film deposited from F = 0.682 target after 2, 10, 30, 60, and 300 min of deposition times. In Fig. 1(a), it is clear that initially only a few InN nanoparticles are deposited on the sapphire surface. The existence of InN alloy was confirmed by x-ray photoelectron spectrometer (XPS) shown in Fig. 2 . It can see that the In3d split to the 3d5/2 at 443.4 eV and 3d3/2 at 451.3 eV and a 397.2 eV peak corresponds to N1s for InN alloy was observed. These binding energy are close to the values of the reported InN films [15,16]. As the deposition proceeds, the InN particles progressively increase in size, while InGaN particles begin to fill the unoccupied spaces on the substrate surface. The introduction of InGaN particles accompanied a reduction in surface roughness, as shown in Fig. 1(b) and 1(c). The root-mean-squared (RMS) roughness value measured by AFM decreased from 5.89 to 2.01 nm as the deposition time increased from 10 to 30 min. It is clear that the film was composed at this time of two distinct grains with few vacancies. Following 60 min of deposition (Fig. 1(d)), RMS roughness was reduced to 1.42 nm. At 300 min, the grains of the film have merged with each other and the RMS roughness reached a minimum value of 1.28 nm, as shown in Fig. 1(e). The growth mode of the film is similar to the nucleation process characteristic of epitaxy [17]. When InN and InGaN grains are distributed on the sapphire surface, some grains act as nucleation sites, facilitating further thin film growth. Details pertaining to the film of growth mechanism are discussed below with reference to XRD data (in Fig. 3(a) ).
Fig. 1 Surface morphologies of the InGaN film deposited from F = 0.682 target after (a) 2, (b) 10, (c) 30, (d) 60, and (e) 300 min of deposition times.
Fig. 3 (a) XRD patterns of the InGaN film following 2, 10, 30, 60, and 300 min of deposition. (b) The ratio of the integrated intensities of the InGaN peak to the InN peak and the RMS roughness as a function of deposition time.
The XRD patterns of the InGaN film following 2, 10, 30, 60, and 300 min of deposition are shown in Fig. 3(a). A 33% indium composition of the InGaN films deposited from F = 0.682 target was determined by measuring the shift of the (0002) InGaN diffraction peak relative to the GaN peak, and applying Vegard’s law. It is notable that an InN peak appeared at 31.33° in the XRD patterns of all samples. In other studies of the InGaN films growth, the presence of the InN peak is mainly attributed to thermal decomposition and re-evaporation of In-N bond during the epitaxial process [18]. In our case, due to the low melting point of indium, when a pulse from the ablation laser impacted the PLD target, indium vapor was generated and allowed to react with the nitrogen plasma, resulting in the formation of InN on the sapphire substrate. This explains why only InN appears after 2 min of deposition. Subsequently, sufficient GaN vapor was provided to yield the desired InGaN alloy which forms as a result of the reaction between indium vapor and GaN vapor. The InGaN peak appears in the XRD pattern at deposition time exceeding 10 min. The reaction yielding InGaN eventually comes to dominate the InN reaction, which is evident from the evolution of the enhanced intensity of the InGaN and InN diffraction peaks with increasing deposition time. The ratio of the integrated intensities of the InGaN peak to the InN peak and the RMS roughness as a function of deposition time are displayed in Fig. 3(b). The ratio trend confirms the occurrence of co-deposition behavior. Despite the initial InN alloy deposition, the later formation of InGaN alloy provides sufficient nucleation sites to promote InGaN growth. The change of growth mode was observed once the ratio of the aforementioned XRD peak intensities reversed within the range of 0.35 to 3.74 and once the RMS roughness was reduced to 1.42 nm. Following 300 min deposition time, the sample roughness decreased and the ratio of the XRD peak intensities achieved a 112% enhancement over sample measured at 10 min, which indicates that the growth mode had completely transferred to a layer-by-layer process. Based on both SEM and XRD results, a schematic of the InGaN deposition on sapphire was displayed in Fig. 4 that illustrates the co-deposition behavior and the growth mode gradually transfers from island growth to layer growth with increasing deposition time.
The XRD patterns of as-deposited and annealed (800°C for 15min) samples labeled A, B, C, and D are presented in Fig. 5 . As shown in Fig. 5(a), the InGaN position shifts from 33.43° to 32.51° (corresponding to indium composition increasing from 33% to 62%) when the F factor was varied from 0.682 to 0.364. The lower F target therefore provides a larger quantity of indium vapor upon ablation, which allows for the fabrication of higher indium content InGaN. In addition, the indium content of the film could be modulated through controlling the concentration of both the indium and GaN vapor, which in turn is controlled by the composition of the dual-compositing target. The sample RMS roughness values of the samples A-D were determined by AFM to be 1.28, 1.56, 1.38, and 3.25 nm, respectively. Although the dual-compositing target allows the fabrication of high indium content InGaN films, the presence of InN inevitably affects the structural and electrical characteristics of the films. Fortunately, InN can be fully decomposed by annealing the film at 800°C. In Fig. 5(b), we observed that all annealed samples retain the same indium composition as the as-deposited films with their peak intensities being slightly larger. The indium produced from the decomposition of InN within the InGaN films tends to react with surrounding InGaN grains to yield InGaN film. It is for this reason that the intensities of the InGaN XRD peak increases upon annealing. However, the magnitude of the intensity increase still depends on the original film quality. For the poorer InGaN films, such as sample D, high temperature annealing led to both the decomposition of InN as well as InGaN degradation.
Fig. 5 XRD patterns of (a) as-deposited and (b) annealed (800°C for 15min) samples labeled A, B, C, and D, respectively.
The relationship between the integrated XRD intensities of InGaN and InN peak (from Fig. 5(a)) and F factor is shown in Fig. 6 . It was found that the quantity of InN almost remains the same even using smaller F target, which indicates the limitation of InN formation in the same PLD conditions and the major indium vapor prefers InGaN reaction to InN reaction. However, the InGaN intensities decrease with F decreases from 0.682 to 0.364. The reason for the InGaN intensities reduction could be the large miscibility gap between GaN and InN [19]. On the other hand, using smaller F target implies decreasing the GaN area exposed by pulse laser, resulting in a lower GaN vapor concentration and thus causing InGaN intensity reduction.
The Hall mobility (μ) and electron concentration (n) of the four annealed samples (InN-free) are plotted in Fig. 7 and listed in Table 1 . It was found that μ increases from 25.5 to 73.8 cm2/V·s then decreases to 22.4 cm2/V·s as the indium composition increases from 33% to 62%. The trend in μ was attributed to the low effective electron mass of the InGaN films and the larger lattice mismatch between the film and the sapphire substrate. Moreover, the high electron concentrations (1.9-3.2 × 1019 cm−3) are ascribed to the presence of nitrogen vacancies generated from fabrication imperfection [20]. The μ trend is similar for the as-deposited and annealed cases, but μ of the as-deposited samples was higher than that of the annealed samples due to the presence of InN in the InGaN film, which results in a lower effective mass. Mixing of InN in the as-deposited samples also reduced the electron concentration compared with the annealed samples.
Fig. 7 Hall mobility and electron concentration of the four annealed samples labeled A, B, C, and D, respectively.
Table 1. Crystal qualities and electrical properties of as-deposited and annealed samples.
4. Conclusion
In conclusion, we have demonstrated the fabrication of InGaN films with indium concentration of 33, 39, 49 and 62% by LT PLD using a controllable InGaN target. This target facilitates a co-deposition reaction, where InGaN grains generated from the reaction of indium and GaN vapor deposits on a sapphire surface and then acts as nucleation seeds to promote further InGaN growth. The growth mode of the InGaN films gradually changes from island growth to layer-by-layer growth. The surface morphology along with the structural and electrical properties of the resulting films indicate that the PLD method used in this study is indeed promising for the development of high indium content InGaN films.
Acknowledgment
This research was supported by National Science Council, The Republic of China, under the Contract No. 98-2221-E-005-006-MY3 and 101-3113-E-005-002-CC2.
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