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Abstract
The optical properties of double (a blue and a green) InGaN/GaN multiple quantum wells (MQWs) on unstressed bulk GaN and a compressive stressed GaN template on patterned sapphire substrate (PSS) were investigated. Photoluminescence intensity of green MQWs on bulk GaN was about 6 times higher than that on GaN/PSS, while the intensity difference of blue MQWs was less than two times. It was found that the existing stress played an important role in the difference in luminescence. Furthermore, the V-shape pits in blue MQWs stemmed from the pre-existing dislocations while those in green MQWs were derived from the new dislocations.
© 2018 Optical Society of America under the terms of the OSA Open Access Publishing Agreement
1. Introduction
GaN based blue and green light-emitting diodes (LEDs) have been commercialized successfully [1,2]. The most commonly used active region of blue and green LEDs is the InGaN/GaN multiple-quantum-wells (MQWs) [3,4]. The structures of blue and green LEDs are very similar, however, there exists a vast difference between the luminous efficiencies of them. For instance, the external quantum efficiency (EQE) of blue LED is reported as high as 84% [5], while the EQE of green LED is only 34% [6]. Various reports have owed this to the deteriorated material quality when the indium component augments. Because of the large lattice and thermal mismatch between InxGa1-xN and GaN as well as the poor indium incorporation, many problems are introduced, such as abundant dislocations [7], local alloy fluctuations [8] and even phase separation [9]. Many optimizations of growth processes and devices structures have been implemented to ameliorate this. Using patterned sapphire substrate (PSS) is the most successful commercially used method, which is simple and useful [10]. However, there exists mass stress in InGaN/GaN MQWs structure grown on PSS [11,12]. Whereas, few work has discussed the relationships between stress, crystalline quality and optical properties of InGaN/GaN MQWs structure grown on PSS. In this paper, we grew double InGaN/GaN MQWs structures which contain two indium components on unstressed bulk GaN and GaN/PSS template to discuss the impact of sublayers’ stress states on blue and green luminescence simultaneously.
2. Growth and experiment
Both of the double InGaN/GaN MQWs structures were grown by metal organic chemical vapor deposition (MOCVD) under identical growth condition. The substrates were free-standing GaN and GaN/PSS template. The bottom diameter, interval spacing and height of the cone-shaped PSS were 2.7, 0.3, and 1.7 µm, respectively, and the thicknesses of the bulk GaN and the GaN template were 350 µm and 4.1 µm, respectively. The MQWs structure consisted of a 200-nm-thick GaN buffer layer and a 60-nm-thick AlGaN stress regulate layer. The AlGaN layer was followed by a 190-nm-thick n+-GaN layer and a 1.3-µm-thick n-GaN layer. The first MQWs were 5 periods of blue InGaN/GaN MQWs whose indium content was 13%, then, 9 periods of green InGaN/GaN MQWs with 25% indium content followed. The subsequent growth was a low temperature p-type layer and a AlGaN electron blocking layer (EBL), which was followed by a p-GaN layer and a p+-GaN contact layer. The double InGaN/GaN MQWs structures grown on bulk GaN and GaN/PSS template were named as sample A and sample B, respectively.
3. Results and discussion
A confocal Jobin Yvon LabRam HR800 micro-Raman spectrometer with a charge-coupled device (CCD) detector and an optical microscopy system was used to study the stress status of the bulk GaN and the GaN/PSS template [13]. The 514 nm laser diode was used as an excitation source and the excitation time was 5 sec. The E2 (high) mode peak of GaN is very sensitive to stress [14], and the strain-free GaN is considered to be at 567.6 cm−1, while the peak would up shift when GaN layer is under compressive stress [15]. As shown in Fig. 1, the E2 (high) peaks of bulk GaN and GaN/PSS template are located at 567.6 cm−1 and 570.6 cm−1, respectively. So the GaN/PSS template is under compressive strain while the bulk GaN is under unstressed state. In addition, GaN film grown on sapphire substrate frequently has a great quantity of dislocations, owing to the large lattice and thermal expansion coefficient mismatch [16]. Dislocation density of GaN film can be decreased to 106~107 cm−2 by homoepitaxial growth on bulk GaN substrate [17]. The intensity of the bulk GaN is about an order of magnitude higher, it reveals the bulk GaN indeed has a lower dislocation density [18,19].
Photoluminescence (PL) spectra at room temperature are given in Fig. 2. The test instrument was still the confocal Jobin Yvon LabRam HR800 while the wavelength of the excitation laser was 325 nm and the integral time was 0.1 sec. Both samples display double peaks with peak positions located separately in blue band and green band. The blue peak intensities of the two samples differ less than two times. However, the green peak intensity of sample B is much weaker than that of sample A, and their intensities vary by about 6 times. The potential determinants of these differences are discussed in the following paragraph.
D8 DISCOVER was utilized to evaluate the crystalline quality. Figure 3 shows the high-resolution X-ray diffraction (HRXRD) ω-scan rocking curves (RCs) of the (002)/(102) reflections for the two samples. The full width at half maximums (FWHMs) of (002) reflection for sample A and B are 42 arcsec and 269 arcsec, respectively, and the FWHMs of (102) reflection for sample A and B are 48 arcsec and 271 arcsec, respectively. It infers that sample A has lower edge/screw dislocation densities [20,21]. According to results of HRXRD, the total dislocation densities of 1.6 × 107 cm−2 and 5.3 × 108 cm−2 can be calculated for sample A and B, respectively [22].
Surface morphologies of sample A and B were measured by a Dimension Icon atomic force microscope (AFM) in tapping mode, the samples were etched by H2 for 10 min under 1000°C to evaluate the dislocation density. 20 × 20 µm2 AFM images are presented in Fig. 4. It can be seen that there are much more dark points on surface of sample B, undoubtedly, sample B has a higher dislocation density. As a rough estimate, the dislocation densities of sample A and B are 1.2 × 107 cm−2 and 2.6 × 108 cm−2, respectively. The results of AFM are in line with the HRXRD results.
Fig. 4 AFM images of (a) sample A and (b) sample B. Samples were etched by H2 in 1000°C environment for 10 min to show up dislocations.
In order to further investigate the strain induced microstructures and dislocations, cross-sectional transmission electron microscope (TEM) images of MQWs were acquired on a Tecnai G2 F20 S-Twin operating at 200 kV. Figures 5(a), 5(c) and 5(e) are the TEM images of different sizes of sample A, while 5(b), 5(d) and 5(f) are images of sample B. It’s very striking that there are numerous new dislocations generated in green MQWs of sample B, by contrast, no dislocation is generated in green MQWs of sample A at all. We assume that the intrinsic stress of GaN sublayer affected the subsequent growth. It is reported that tensile stress will be introduced when growing InGaN on GaN due to the large size of indium atom [23,24]. Although the interlaminar stresses between InGaN and GaN layers are large, the layers of MQWs are thin so they are strained completely. Whence if the underlying stress is zero, ideally, there is no dislocation generates when growing green MQWs, just like sample A. Nevertheless, the compressive stress of the GaN/PSS template beneath green MQWs of sample B is too large. The consequence of this compressive stress entangled with tensile stress of InGaN is new dislocations being produced in the green MQWs of sample B. In contrast, the indium content of blue MQWs is low so the InGaN can endure some compressive stress and scarcely any new dislocation forms in blue MQWs of both samples. In addition, some V-shape pits are found to generate at the interfaces of InGaN/GaN in sample B. Each V-shape pit is connected to the threading dislocation at their bottom, indicating that the pits are initiated at dislocation cores in the presence of indium [25,26]. The V-shape pits in blue MQWs stem from the pre-existing dislocations while those in green MQWs are derived from the new dislocations caused by large stress. It should also be attached attention that the interfaces of InGaN/GaN of green MQWs in sample A are very abrupt and smooth as shown in Figs. 5(a) and 5(b). While in sample B the interfaces are not abrupt and smooth, even in some parts of the InGaN layers there appear fractures or uneven thickness (as shown in Figs. 5(b) and 5(b)). According to previous reports, the indium-rich InGaN/GaN MQWs have suffered from various effects such as diffusion/cluster formation and desorption compete with indium incorporation [27], most often, indium has trended to cluster around dislocations [28]. This phenomena occur in sample B rather than in sample A, indicating that reducing the underlying stress can suppress the separation and precipitation of indium, and improve the material quality of indium-rich InGaN layer. Because of the luminescence of indium clusters with different components in green MQWs of sample B, the green PL peak of sample B is much wider than that of sample A and the peak red shift. The blue MQWs of samples are not very clear but still can see them grow well as shown in Figs. 5(e) and 5(f), which attributes to the more efficient indium incorporation of low indium MQWs. So the stress of sublayer has little effect on the quality of the blue MQWs.
In order to research the influence of stress on luminescence, wavelength-dependent cathodoluminescence (CL) mapping images of samples are shown in Fig. 6. The green light emitted from green MQWs is demonstrated by Figs. 6(a) and 6(b). Sample B comprises some bright luminescent spots which represent the indium clusters, the dot-like clusters play a key role in emission of indium-rich InGaN/GaN MQWs [29]. However, the newly generated dislocations, which act as nonradiative recombination centers [30] in the green MQWs of sample B are too many, so the emissions of indium clusters around dislocations are drastically reduced just like the black holes engulfing photons. While in sample A the green luminescence is uniformer and brighter than sample B, which means stronger luminescence. These phenomena are due to the well-grown green MQWs of sample A and few new dislocations in it. This also explains the fact that the green light intensity of PL spectrum of sample A far outpace that of sample B. When involved in the emission of blue InGaN/GaN MQWs, CL images look kind of similar with each other as depicted by Figs. 6(c) and 6(d). The absence of bright spots may on account of the high distribution homogeneity of indium component. The dislocation densities of blue MQWs of two samples do not differ much so that the blue luminescence difference is little. Overall, the influence of the initial stress on blue band emission won’t be as remarkable as on green band emission.
Fig. 6 20 × 20 µm2 plan-view CL mapping images for (a) (c) sample A and (b) (d) sample B. CL mapping images are taken at wavelengths of (a) 540 nm and (c) 435 nm for sample A and (b) 566 nm and (d) 426 nm for sample B.
4. Conclusions
To investigate the correlation between stress and luminescence, double InGaN/GaN MQWs which contain two indium components were grown on bulk GaN and GaN/PSS template under identical growth condition. The bulk GaN was free of stress while the GaN/PSS template was under compressive strain. It was found that PL intensity of green peak of sample A was about 6 times higher than that of sample B, while the blue peak intensities of two samples were differ less than a factor of two. There were many new dislocations formed in green MQWs of sample B, while that of sample A had few. Therefore the weaker green luminescence and clustered indium appeared in the green CL image of sample B. For blue MQWs, they had low indium content so few new dislocations generated in them. CL images exhibited the uniformly distributed and similar blue luminescence of two samples. On the whole, the initial stress of GaN had a great influence on quality and optical properties of green MQWs grown on them while had less on blue MQWs.
Funding
National Key R&D Program of China (2016YFB0400800, 2016YFB0400105); National Natural Science Foundation of China (NSFC) (61574108, 61334002, 61474086, 51302306); Fundamental Research Funds for the Central Universities (JBZ171101).
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