JOURNAL OF APPLIED PHYSICS 108, 093501 共2010兲 Generation of amber III-nitride based light emitting diodes by indium rich InGaN quantum dots with InGaN wetting layer and AlN encapsulation layer C. B. Soh,1,a兲 W. Liu,1 S. J. Chua,1,2,b兲 S. S. Ang,1 Rayson J. N. Tan,1 and S. Y. Chow1 1 Institute of Materials Research and Engineering, Agency for Science, Technology and Research (A*STAR), 3 Research Link, Singapore 117602 2 Singapore-MIT Alliance, National University of Singapore, 4 Engineering Drive 3, Singapore 117576 共Received 5 July 2010; accepted 28 August 2010; published online 2 November 2010兲 Indium rich InGaN nanostructures grown by metalorganic chemical vapor deposition were incorporated in InGaN/GaN quantum wells for long wavelength generation. These results were achieved by optimizing the growth temperature of the nanostructures, InGaN quantum well, the AlN capping layer and the GaN barrier layers. Before the growth of nanostructures, a thin InGaN wetting layer was included to reduce the lattice mismatch as well as to enhance the deposition of indium-rich InGaN nanostructures These individual quantum wells were each subsequently capped with an AlN layer which better preserved the In-rich phase in the nanostructures and prevented the indium interdiffusion between the InGaN/GaN heterojunctions. The AlN capping layer also reduces the effect of piezeoelectric field in the active layers of the light emitting diodes as seen from the reduction in the blueshift in the electroluminescence peaks with higher injection currents. The energy band profile of such a structure is discussed. © 2010 American Institute of Physics. 关doi:10.1063/1.3499618兴 I. INTRODUCTION In the past decade, there has been a considerable effort in research and development into white InGaN light emitting diodes 共LEDs兲 due to it potential saving for lighting applications. Commercial white LEDs make use of blue LEDs coated with luminescent materials such as yellow emitting phosphor Y3Al5O12 : Ce3+ 共YAG:Ce兲,1 CdSe-ZnS,2 or capped CdSe quantum dots 共QDs兲 with Sr3SiO5 : Ce3+, Li+ phosphor.3 However, it is difficult to achieve a high color rendering index 共of higher than 85%兲 due to their red spectral deficiency.4,5 There is also the drawback due to additional coating step for the light converter. Recently, there have been reports of phosphor free white LEDs using prestrained InGaN wells,6 combination of quantum wells 共QWs兲 with varying thickness7 and indium phase separated multiple quantum wells 共MQWs兲 as the active layers.8 However, the internal quantum efficiency of the long wavelength component was relatively low, leading to drastic degradation in its emission efficiency after fabrication9 with the issue of quantum confined Stark effect not addressed. To realize superior device performances, it is desirable to introduce QDs or nanostructures into the active layers of the devices.10 Self assembled QDs structures or indium rich InGaN nanostructures have been known to increase the recombination efficiency of the emitting layers.11,12 The QDs generated by the self assembly process generally have a random distribution in terms of its size12 which causes the inhomogeneous broadening of the electroluminescence 共EL兲 spectra. Chen et al.13 has reported on the growth of InGaN nanostructures with homogeneous size using electron beam lithography techa兲 Electronic mail: cb-soh@imre.a-star.edu.sg. Electronic mail: elecsj@nus.edu.sg. b兲 0021-8979/2010/108共9兲/093501/5/$30.00 nique to pattern a template. However, this technique is not feasible for large area patterning of a 2 in. LEDs wafer as it is time consuming. Furthermore, LEDs have to undergo subsequent p-GaN growth and annealing for Mg activation which is carried out at a temperature much higher than that for the growth of indium rich InGaN QDs. This will lead to unavoidable spectra broadening and blueshift in peak wavelength. Wang et al.14 has reported on the effect of increasing the annealing temperature for p-GaN activation in InGaN QDs based LEDs. With the annealing temperature at 830 ° C, degradation of the QDs takes place through atomic diffusion of indium into the underlying wetting layer, leading to the formation of a wide InGaN QWs with higher indium contents. These effect enhances the blueshift in the emission wavelength at a higher drive current.15 Due to the constraint in the materials growth properties, we have modified the growth structures for the active regions to preserve the indium rich nanostructures formation by preventing out-diffusion of indium at the high temperature thermal annealing stage. Monolayers of InGaN wetting layer is first grown to reduce the strain for the incorporation of self-assembled indium rich InGaN QDs. This is similar to the use of an InGaAs wetting layer to grow self assembled arsenide QDs.16 A thin aluminum nitride 共AlN兲 layer is then incorporated to serve as a capping layer after the growth of the indium rich InGaN QDs. AlN capping mechanism has been studied in detail for GaN QDs/AlN system.17,18 This work will focus on the study of the mechanism for self assembly of indium-rich InGaN QDs on InGaN wetting layer, followed by capping with a thin AIN layer. The band structure which leads to generation of amber emission from these LEDs structures will be proposed. 108, 093501-1 © 2010 American Institute of Physics Downloaded 09 Nov 2010 to 132.178.122.236. Redistribution subject to AIP license or copyright; see http://jap.aip.org/about/rights_and_permissions 093501-2 Soh et al. J. Appl. Phys. 108, 093501 共2010兲 II. EXPERIMENTS The sample used in this study is grown by metalorganic organic chemical vapor deposition. c-plane sapphire was used as the substrate with hydrogen as the carrier gas. After the growth of the 25 nm low temperature 共LT兲-GaN buffer layer at 525 ° C, a 1.2 m thick undoped GaN, and 1.5 m thick Si-doped GaN were grown. These layers serve as the template for the growth of MQWs. Three MQWs samples were grown on the GaN template and are labeled as A, B, and C. In sample A, three periods of InGaN/GaN QWs were grown with dimensions of 4 nm/12 nm, respectively. Each well consists of a wetting InGaN layer, 1 nm thick grown with trimethylindium 共TMIn兲 flow rate at 24 mol/ min and trimethylgaldium 共TMGa兲 at 12 mol/ min, designated as ␣-layer, followed by 2 nm of InGaN growth with TMIn flow rate kept at 70 mol/ min and TMGa at 12 mol/ min, designated as -layer. The InGaN well layer, i.e., the ␣-layer and -layer, were grown at a temperature of 725 ° C. Sample B, consists of a set of 3 QWs as well. However, after the growth of the wetting layer, i.e., the ␣-layer, the sample underwent TMIn treatment10,19,20 where TMIn was allowed to flow under NH3 rich condition for 30 s to form the indium rich InGaN QDs before the growth of the -layer. For sample C, growth condition for the QWs was similar to sample B except for the additional AlN cap layer of ⬃1.2 nm thick grown after deposition of the -InGaN layer. The three MQWs samples were then capped by a 10 nm GaN layer. These three MQWs structures were used for subsequent fabrication of LEDs by depositing a 25 nm thick p-Al0.25Ga0.75N layer and a 200 nm thick p-GaN layer at 870 ° C. An in situ chamber annealing for 20 min in nitrogen ambient at 825 ° C was carried out for the LEDs sample to activate Mg in its pGaN layer. The LEDs sample was labeled as sample X and Y using the MQWs structures of samples B and C, respectively. A. Surface morphology study by atomic force microscopy „AFM… In order to simulate the effect of possible degradation of the QDs during the pGaN activation process, in situ chamber annealing in N2 ambient was carried out for 5 min for QW sample, B and C. Figure 1 shows the tapping mode AFM image of the surface morphology of the QWs sample. Figure 1共a兲 shows that the formation of indium rich InGaN QDs on the InGaN wetting layer after TMIn treatment. The use of N2 as contrast to H2 as a carrier gas for growth of the QDs facilitate indium incorporation as it slows down the diffusion of indium adatoms. The reduction in the atomic diffusion of indium enhances the Stranski–Krastanow growth of InGaN QDs. Coupled with the higher continuous flow of NH3 at 18 slm 共slm denotes standard liters per minutes兲 as compared to conventional InGaN well growth 共12 slm兲, a higher amount of N interacts and forms bonds with impinging indium at Ga sites of the InGaN wetting layer. This explains the formation of indium rich InGaN QDs of size ⬃20 nm in diameter and ⬃1.5 nm in height with a density of ⬃1 ⫻ 1010 cm−2. The InGaN QDs were then capped with a 1.5 nm InGaN well FIG. 1. 共Color online兲 AFM image of 共a兲 indium rich InGaN QDs deposited on the InGaN wetting layer, 共b兲 single QW sample B with clusters of InGaN QDs formed by layer growth of the InGaN well layer which wets the QDs, 共c兲 single QW sample C with stripelike AlN islands on the InGaN well formed with addition of AlN capped layer, 共d兲 second QDs layer with higher density grown on single QWs sample C, 共e兲 sample B with annealing at 820 ° C for 10 min, and 共f兲 sample C with annealing at 820 ° C for 10 min. layer with the NH3 flow rate reduced to 12 slm and the temperature increased to 725 ° C to promote layer by layer growth. Figure 1共b兲 shows the AFM image of sample B after the growth of the InGaN well layer,  layer, with the embedded InGaN QDs. The InGaN QDs clustered and formed nanostructures due to layer growth of the InGaN over the InGaN QDs. In order to mitigate the effect of the degradation of these indium rich QDs with further deposition and annealing for the pGaN layer at a higher temperature, a 1.2 nm AlN was deposited at 780 ° C to serve as a capping layer on the -InGaN. Figure 1共c兲 shows that the AlN cap leads to the formation of stripelike AlN islands on the InGaN well 共-layer兲 which consisted of the embedded InGaN QDs. GaN QDs capped by AlN displayed the unique characteristic of being insensitive to In out-diffusion.17,18 This is attributed to the formation of stable Al–N bonds and also to the relatively lower mobility of Al adatoms on the film surface at a lower chamber temperature of 780 ° C. When AlN is first deposited, it will first wet the surface of the InGaN nanostructures with its first few monolayers and subsequent AlN grows preferentially to enlarge the nanostructures. When AlN is grown to a thickness of 1.2 nm, it formed elongated Downloaded 09 Nov 2010 to 132.178.122.236. Redistribution subject to AIP license or copyright; see http://jap.aip.org/about/rights_and_permissions 093501-3 Soh et al. J. Appl. Phys. 108, 093501 共2010兲 FIG. 2. 共Color online兲 HR-STEM image taken along 共21̄1̄0兲 direction for 共a兲 sample B with embedded InGaN QDs in the InGaN well layers, 共b兲 sample B after thermal annealing which illustrated the out-diffusion of QDs to the surrounding matrix and interdiffusion into the GaN barrier causing the shrinkage of the QDs dimension, and 共c兲 sample C after thermal annealing with the QDs better preserved but with some out-diffusion into the InGaN active well region. islands to coalesce the InGaN nanostructures. Aluminum also has a lower vapor pressure as compared to indium and this can reduce the diffusion length of other group III atoms. This implies that the growth of the subsequent GaN barrier and InGaN wetting layer can be affected by the base AlN layer. There will be a reduction in the diffusion length of In atoms on the first QW layer leading to the increase in the QD density as shown in Fig. 1共d兲 compared with Fig. 1共a兲. The QDs density for the second repeat layer is estimated to be 3 ⫻ 1010 cm−2 as compared to a QDs density of ⬃5 ⫻ 109 cm−2 for a single well layer. The results suggest that there is substantial reduction in strain for the embedded InGaN in the second repeat layer. After annealing for 10 min in N2 ambient at 825 ° C, it was observed that there is a change in the surface morphology for sample B and C. In sample B, the nanostructures of the QW formed through clustering of the QDs experienced further agglomeration to give rounded islands as shown in Fig. 1共e兲. As for sample C with AlN cap layer, the elongated structure further bunched together to form connected islands with dimension of ⬃120– 220 nm. The observed pits separating the structures in the layer are also significantly wider and deeper. This enables the formation of stable QW elongated structures as observed in Fig. 1共f兲. B. Transmission Electron Microscopy „TEM… analysis In order to explore the impact of annealing on Mg activation in the p-GaN layer on the optical properties of the QDs embedded LEDs, a detailed investigation by TEM was carried out as illustrated in Fig. 2. In situ chamber annealing was carried out for LEDs sample, X and Y, at 820 ° C for FIG. 3. 共Color online兲 PL spectra at 10 K for 共a兲 sample C and 共b兲 sample B before and after thermal annealing for 10 min in N2 ambient. activation of Mg in the p-GaN layer of the LEDs. As shown in Fig. 2共b兲, for LEDs sample X after chamber annealing, there is a shrinkage in the QDs size with diffusion of indium atoms out of the QDs to the underlying InGaN wetting layer. This caused the InGaN well 共 layer兲 to form InGaN clusterlike structure as compared to the more uniform well layer of LEDs sample X before in situ chamber annealing as shown in Fig. 2共a兲. The QDs appears to form small dotlike structures and this shrinkage in size can contribute to the blueshift in emission wavelength for the LEDs. Chuo et al.21 has reported that the density of indium rich dotlike structures tends to decrease after thermal annealing due to out-diffusion of indium into the surrounding matrix including the GaN barrier and the InGaN QWs. With the addition an AlN cap layer, we observed from TEM image in Fig. 2共c兲 that the InGaN QDs in the InGaN well are better preserved as there is minimal out-diffusion of the InGaN nanostructures. The capping mechanism of GaN QDs by AlN capping layer has been demonstrated by Coraux et al.17 It is observed that the capping of InGaN nanostructures in the QWs by AlN carried out in this work shows a similar characteristic in that its involves the wetting of InGaN nanostructures by AlN followed by intervalley filling where preferential AlN growth occurs. Downloaded 09 Nov 2010 to 132.178.122.236. Redistribution subject to AIP license or copyright; see http://jap.aip.org/about/rights_and_permissions 093501-4 Soh et al. J. Appl. Phys. 108, 093501 共2010兲 FIG. 4. 共Color online兲 Current dependent EL spectra for 共a兲 LEDs X and 共b兲 LEDs Y with chamber annealing for 20 min in N2 ambient at 820 ° C for Mg activation. 共c兲 Plot of the light-output power and peak emission wavelength of LEDs as a function of injection current. C. Photoluminescence „PL… from QDs/InGaN MQWs TEM images show that with annealing, there is shrinkage in the size of the QDs with out-diffusion of indium from the In-rich InGaN QDs into the InGaN well and interdiffusion into GaN barrier. With the addition of the AlN capping layer, the InGaN QDs are better preserved. Rapid thermal annealing for 10 min at 820 ° C in nitrogen ambient was carried out for QWs, sample B and C to simulate the effect of chamber annealing during Mg activation of LEDs. Figure 3 shows a typical low temperature PL at 10 K for samples B and C before and after thermal annealing process. There is a blueshift in the PL emission for sample B whereas for sample C, a redshift in its PL emission is observed. The shift in the PL emission can be attributed to following reasons. First, there exists a strong piezoelectric field in InGaN/ GaN QW which is further enhanced in InGaN/GaN QDs filled QWs 共Ref. 22兲. The strain in the mismatched layers, reduces the diffusion length of the QDs and enhanced their incorporation as well as its density. During thermal annealing at 820 ° C, there is a reduction in the strain of InGaN 共well兲 active region and hence the indium in the QDs starts to diffuse with greater ease. The out-diffusion of indium enhances the indium content in the QW region 共i.e., the InGaN wetting and well layer兲. As shown in Fig. 3共a兲, after thermal annealing of sample C 共the QW sample with the AlN cap layer兲, the PL spectra of the dominant peak associated with the InGaN wetting and the InGaN well layer are redshifted by 15 nm due to out-diffusion of indium from the indium-rich QDs. For sample B without the AlN cap layer, out-diffusion of the In-rich phase into the surrounding matrix and the interdiffusion between the InGaN/GaN heterojunctions would occur simultaneously. With longer duration of annealing, indium diffusion into the GaN barrier, leading to a lower indium FIG. 5. Schematic representation of the energy band profile of strained AlN/InGaN QDs/InGaN/GaN QW system. Downloaded 09 Nov 2010 to 132.178.122.236. Redistribution subject to AIP license or copyright; see http://jap.aip.org/about/rights_and_permissions 093501-5 J. Appl. Phys. 108, 093501 共2010兲 Soh et al. content in the well. This accounts for the blueshift in the dominant peak position for the well and the wetting layers for sample B as illustrated in Fig. 3共b兲. D. EL measurement Figures 4共a兲 and 4共c兲 shows the injection current dependence of the EL spectra for LEDs X and Y with MQWs of sample B and C, respectively. The EL spectra for LED X in Fig. 4共a兲 shows a yellow emitting LED with a emission band ranging from 500 to 680 nm and a dominant peak at ⬃562 nm at 70 mA. The peak intensity saturated with an injection current of 70 mA. The dominant peak wavelength shifts from 582 to 562 nm with the increase in injection current from 30 to 70 mA as illustrated in Fig. 4共b兲. On the otherhand, for LEDs Y, its EL spectra shown in Fig. 4共c兲 have an emission band ranging from 550 to 750 nm with a shoulder emission 共at 500 nm兲 attributed to the InGaN wetting layer. Unlike LED X, there is a no prominent shoulder peak is observed other than the slight broadening at its lower wavelength spectrum. The minimal out-diffusion of the indium rich QDs enables the composition and thus the emission wavelength of InGaN wetting layer to be retained during the p-Mg activation step. The EL spectrum with injection current of 80 mA onwards shows a splitting of the dominant peak into peak A and B. Peak A is likely due to InGaN well while Peak B is due to the high indium QDs content nanostructures leading to its higher emission wavelength. As shown in Fig. 4共d兲, saturation of its peak emission was not encountered. However, there is slight shift in its dominant peak from 668 to 658 nm as injection current increased from 30 to 70 mA. The minimal shift also suggests that the effect of piezeoelectric field in the QWs layer is smaller in LED Y. Thus the AlN capping layer has not only prevented outdiffusion of indium in the QDs but also reduces the strain between the InGaN well and GaN barrier layer for the second and the third periods of the QW growth. E. Energy band profile Figure 5 gives the schematic representation for the polarizations and the energy band diagram of the strained InGaN wetting layer/InGaN QDs/InGaN well/AlN/GaN QWs structure. In a polar Ga-face 共0001兲 material, negative and positive spontaneous polarization 共Psp兲 charges are induced at the top and bottom interfaces of the wurtzite material system. 共Al, InGaN, Ga兲N is a polar material and the polar 共0001兲 axis is parallel to the QW growth direction. As the InGaN layer is compressively strained, piezoelectric polarization charges are induced in opposite direction to the Psp charges.23 On the otherhand, for the sample with AlN capping layer, it is under tensile strain as the Al atomic size is much smaller than Ga and In. This will assist in lowering the piezoelectric polarization charge accumulation at the interface of the InGaN well. This accounts for the much smaller blueshift in the EL emission 共by 10 nm兲 for the amber LED, Y, with the AlN cap layer as compared to yellow LED, X, 共by 20 nm兲 as the injection current increases from 30 to 70 mA. III. CONCLUSIONS In summary, we report on the generation of amber emitting LEDs with minimal shift in its EL spectrum under higher injection current. This is attributed to implementation of an AlN capping layer on top of the InGaN well with embedded indium rich QDs. The AlN cap layer enables the LEDs to undergo an annealing process with a longer duration and at a higher temperature. 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Onuma, B. A. Haskell, A. Chakraborty, T. Koyama, P. T. Fini, S. Keller, S. DenBaars, J. Spexk, U. K. Mishra, S. Nakamura, S. Yamaguchi, S. Kamiyama, H. Amano, I. Akasaki, J. Han, and T. Sota, Nature Mater. 5, 810 共2006兲. Downloaded 09 Nov 2010 to 132.178.122.236. Redistribution subject to AIP license or copyright; see http://jap.aip.org/about/rights_and_permissions