GaN nanopillar LEDs

Accepted Manuscript Title: Diameter-dependent photoluminescence properties of strong phase-separated dual-wavelength InGaN/GaN nanopillar LEDs Authors: Qiang Wang, Ziwu Ji, Yufan Zhou, Xuelin Wang, Baoli Liu, Xiangang Xu, Xingguo Gao, Jiancai Leng PII: DOI: Reference:

S0169-4332(17)30752-3 http://dx.doi.org/doi:10.1016/j.apsusc.2017.03.093 APSUSC 35469

To appear in:

APSUSC

Received date: Revised date: Accepted date:

10-8-2016 6-3-2017 9-3-2017

Please cite this article as: Qiang Wang, Ziwu Ji, Yufan Zhou, Xuelin Wang, Baoli Liu, Xiangang Xu, Xingguo Gao, Jiancai Leng, Diameter-dependent photoluminescence properties of strong phase-separated dual-wavelength InGaN/GaN nanopillar LEDs, Applied Surface Science http://dx.doi.org/10.1016/j.apsusc.2017.03.093 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

Diameter-dependent photoluminescence properties of strong phase-separated dual-wavelength InGaN/GaN nanopillar LEDs

Qiang Wang a, b, *, Ziwu Ji b, **, Yufan Zhou c, Xuelin Wang c, Baoli Liu d, Xiangang Xu e, Xingguo Gao a, Jiancai Leng a

a

School of Science, Qilu University of Technology, Jinan 250353, China

b

School of Microelectronics, Shandong University, Jinan 250100, China

c

School of Physics, Shandong University, Jinan 250100, China

d

Beijing National Laboratory for Condensed Matter Physics, Institute of Physics, Chinese Academy of Sciences, Beijing 100190,

China e

Key Laboratory of Functional Crystal Materials and Device (Ministry of Education), Shandong University, Jinan 250100, China

*Corresponding author: [email protected] (Q. Wang); Tel.: +86 531 89631268; fax: +86 531 89631268. **Corresponding author: [email protected] (Z. Ji); Tel.: +86 531 88365248; fax: +86 53188564886.

The highlights of our manuscript are summarized as follows: · Nanopillar LED with smaller diameter shows a larger strain relaxation in the MQWs. · Nanopillar induced blue shift of green peak is smaller than that of blue peak. · Nanopillar induced blue shift of green/blue peak at 300 K is smaller than at 4 K. · PL intensity decreases with reducing nanopillar diameter with same pillar density.

Abstract In this paper, strong phase-separated blue/green dual-wavelength InGaN/GaN nanopillar (NP) light emitting diodes (LEDs) with the same NP density and various NP diameters were fabricated using focused ion beam etching. Micro-Raman spectroscopy was used to show the effect of NP diameter on the strain relaxation in the multi-quantum-wells (MQWs). The effect of NP diameter on optical behaviors of the strong phase-separated dual-wavelength InGaN/GaN NP LEDs was investigated for the first time by using micro-photoluminescence (PL) spectroscopy. The blue

shifts of PL peak energies of the NP LEDs showed that the NP LED with a smaller diameter exhibited a larger strain relaxation in the MQWs, as confirmed by micro-Raman results. And the blue shift of green emission was smaller than that of blue emission. The total integrated PL intensities from the NP arrays were enhanced compared to the as-grown sample due to the increased recombination rate and light extraction efficiency. The enhancement factor decreased with decreasing the NP diameter in our experiments, which indicated that the loss of active volume was gradually dominant for the luminous efficiency of NP LEDs compared to the increased recombination rate and light extraction efficiency. Keywords: Nanopillar LED, Phase-separation, Dual-wavelength, Strain relaxation, Photoluminescence 1. Introduction The InGaN-based semiconductor has been widely investigated for its applications in solid state lighting. Its direct band gap, tunable from 0.7 eV (for InN) to 3.4 eV (for GaN), covers the near infrared to near ultraviolet spectral region and makes this material system important for full-color multimedia applications [1–3]. Nevertheless, it is well-known that the luminous efficiency of the conventional planar InGaN-based light emitting diodes (LEDs) is obstructed by the following two main drawbacks. One hand, the conventional planar InGaN-based LEDs are usually grown on a c-plane sapphire substrate. InGaN/GaN multi-quantum-well (MQW) suffers a strong internal polarization fields directed along the c-axis, especially for higher indium content. The polarization fields occur due to a combination of spontaneous polarization and strain-induced piezoelectric polarization induced by the large lattice mismatch between GaN and InGaN. Polarization-induced quantum-confined Stark effect (QCSE) limits the overlap of the electron and hole wave functions leading to a reduction in the radiative recombination rate of InGaN/GaN MQW LEDs [4,5]. On the other hand, the large refractive-index (n) difference between GaN (n = 2.5) and air (n = 1) causes most photons emitted from the quantum well (QW) to be reflected within the LED and ultimately converted to heat. Only 4% of emitted photons can escape into air due to low escape cone (23.6°), resulting in poor light extraction

efficiency (LEE) [6]. In order to develop high-efficiency LEDs, it is imperative to find an effective way to improve both recombination rate and LEE. Nanostructures such as nanowire, nanohole and nanopillar (NP) arrays provide an alternative way to mitigate these problems associated with conventional planar LED. Among them, NP arrays have drawn increased interest due to several advantages over planar epitaxial. Recently, several groups investigated the emission efficiency of InGaN/GaN MQW NP LEDs fabricated by bottom-up or top-down approaches [7-9]. These studies have shown enhancement in PL intensity from NP arrays compared to the planar sample. First, NP structure can relieve the strain-induced piezoelectric field across the QW. Therefore, the QCSE is alleviated and the radiative recombination rate can be improved [7]. Second, NP structure can enhance the LEE due to the increased scatter and reduced reflection at the sidewall of the NPs and the light-guiding effect along NP arrays [9]. It's worth noting that these reports have shown the influence of NP arrays on the luminous efficiency of LED with just a single blue or green wavelength arising from the InGaN matrix. Some studies have reported the blue/green dual-wavelength InGaN/GaN MQW LED, due to phase-separation in higher indium content InGaN epilayers [10,11]. Two InGaN-related emissions observed in the photoluminescence (PL) spectrum are assigned to InGaN matrix (blue emission) and In-rich quasi-quantum dots (QDs) (green emission). However, to the best of our knowledge, there is no report on the influence of NP diameter on the PL emission of strong phase-separated dual-wavelength InGaN/GaN MQW NP LED. In this work, the blue/green dual-wavelength InGaN based planar LED was etched into NPs with various diameters by focused ion beam (FIB). The influence of diameter on the micro-Raman (μ-Raman) and micro-PL (μ-PL) spectra of NP arrays were studied by comparing to the planar sample. 2. Experimental 2.1. Planar LED fabrication The GaN-based planar LED structure (defined as Planar) was grown via metal organic chemical-vapor deposition (MOCVD) on the c-plan sapphire substrate. The

epitaxial structure consisted of a 40-nm GaN nucleation layer, a 2-μm unintentionally doped GaN buffer layer, a 2-μm Si-doped n-GaN layer, ten pairs of InGaN/GaN (3 nm/8 nm) MQWs, a 65-nm low temperature Mg-doped p-GaN layer, a 28-nm p-AlGaN layer, and a 200-nm p-GaN layer. The nominal indium composition of the QWs was about 32%. 2.2. Nanopillar fabrication Figure 1 illustrates the process flow for the fabrication of InGaN/GaN NPs covering an area of 8×8 μm2 using the FIB. In brief, a ~100 nm Au layer was first deposited on the epi-surface of planar LED by e-beam evaporation in order to enhance surface conductivity and avoid surface charging. Then, the NP LEDs with different diameters were fabricated by direct etching using a FIB system (SMI-3050, SII Nanotechnology, Japan) according to the different bitmaps inserted in the SII software. As shown in Fig. 2(a), the bitmap image consists of white circle dot arrays with a black background. The area for black pixels were milled by the 30 keV Ga+ ion with a beam current of 300 pA. After the etching process, the remnants of the Au were removed with a KI/I2 aqueous solution. Subsequently, surface damage generated in the NP sidewalls due to the high-energetic ion bombardment, were further etched with KOH of 20 wt.% dissolved in ethylene glycol, heated to 60 °C, with an etching time of less than 4 min. Finally, these samples were subjected to rapid thermal annealing at 800 °C for 1 min under ambient nitrogen. According to the scanning electron microscopy (SEM) images, shown in Figs. 2(b)-2(d), NP arrays with diameters of approximately 320, 250 and 180 nm (defined as NP320, NP250 and NP180, respectively) and heights of approximately 700 nm were fabricated. Three NP LEDs were fabricated on the same wafer, and the interval between two adjacent NPs was 400 nm, with a certain NP density of 7.2×108 cm -2. 2.3. Other details The strain state of each sample was measured by μ-Raman scattering spectra at room temperature in backscattering geometry using a μ-Raman system (Jobin-Yvon LabRAM HR800). The samples were excited by the 632.8 nm line of He-Ne laser. The incident and scattered light propagated parallel to the c-axis, which in turn was

normal to the growth surface. For temperature dependent μ-PL measurements, the samples were mounted in an open-cycle helium flow cryostat, and the temperature was maintained at 4 and 300 K, respectively. A continuous wave He–Cd 325 nm laser beam at an incident power of 2 mW was focused down to a spot size of 2 μm by a Mitutoyo 100× objective lens (numerical aperture of 0.5). The emissions were dispersed by a Jobin-Yven iHR550 monochromator, and detected by a liquid N2-cooled Synapse CCD detector, with a spectral resolution of approximately 0.15 nm. 3. Results and discussion 3.1. Raman characterization

In order to study the state of strain in the MQW regions after NP formation, μ-Raman scattering spectra were conducted on the planar and NP samples. As shown in Fig.3, the 569 cm−1 peaks are the E2(H) mode of GaN, and the shoulder lines around 560 cm−1 are the E2(H) mode of InGaN [12,13]. The InGaN E2(H) phonon mode can be used indirectly to evaluate the strain states of InGaN/GaN MQW regions. In addition, the reduced compressive strain in the InGaN/GaN MQWs regions can lead to low-frequency shifts of InGaN E2(H) phonon mode [14–16]. In Fig.3, the μ-Raman spectra were fitted with two Voigt line-shape functions, and the peak position of InGaN E2(H) mode was 563.77, 562.82, 562.05 and 560.91 cm−1 for planar, NP320, NP250 and NP180, respectively. That Raman peaks shift to a lower wavenumber indicates that: 1) the compressive strain in the MQWs regions, due to the lattice mismatch and thermal mismatch between high In-content InGaN active layer and GaN barrier layer, partially relax as a consequence of NP formation with large free surface areas and high surface/volume ratios; 2) the degree of strain relaxation of MQWs embedded in the NP LEDs increases with decreasing NP diameter, due to the thinner NP possessing the higher surface/volume ratio. 3.2. PL characterization Figures 4(a) and (b) show the PL spectra of NP320, NP250, NP180 and planar LED at 4 and 300 K, respectively. Two separated InGaN-related peaks due to strong

phase separation were clearly found in all the PL spectra. The shorter (blue) and longer (green) wavelength are assigned to the InGaN matrix-related emission and the In-rich quasi-QDs-related emission, respectively [10,11]. Figure 5 shows the blue and green PL peak energies, as obtained by fitting the corresponding PL spectrum to Gaussian functions, versus the NP diameters. As the NP diameter becomes smaller, both blue and green peaks energies show a clear blue shift at 4 and 300 K, respectively. There are two possible reasons for the blue shift: 1) the quantum size effect; 2) the strain relaxation of the QW. However, the blue shift caused by quantum size effect can be neglected for the QW-disks in the NPs with large diameters of above 15 nm [17]. Thus, the obvious blue shift can be attributed to the strain relaxation as confirmed by μ-Raman spectra mentioned above. Benefiting from the NP structures, the surface/volume ratio is increased and the internal strain of MQWs is partially released, due to the weaker lateral confinement in QWs embedded in the NP LEDs [18]. The relaxed strain reduces the piezoelectric field inside the QWs, so that the PL peak energies of the NP LEDs show blue shift. As the NP diameter becomes smaller, the weight of strain-relaxed region becomes larger, due to the fact that the major strain relaxation region is limited to 20 nm from the edge of the NP [9]. Thus, the smaller diameter NP LED should undergo greater strain relaxation in the MQW region, and its PL peak energies show more significant blue shift. Furthermore, PL peak energies for both blue and green emission at 300 K show a pronounced red shift relative to the corresponding PL peak energies at 4 K, as shown in Fig. 5, mainly due to the band gap narrowing with temperature [19]. Figures 6(a) and (b) show the PL peak energy shift versus NP diameters at 4 and 300 K, respectively. The blue shift of blue and geen emission of the NP LEDs, as compared with the planar sample, are defined to △EM and △ED, respectively. It is noticed that △ED is smaller than △EM for all the NP samples at both 4 and 300 K. This could be explained by the fact that the QCSE in the quasi-QDs is weaker than that in the InGaN matrix, due to the smaller sizes of the strong localized quasi-QDs [20,21]. Moreover, the difference between △EM and △ED shows a slight increase as the NP diameter decreases at both 4 and 300 K. This is because that thinner NP sample has a

larger strain relief, and the blue emission is easier affected by the strain relaxation as mentioned above. In addition, △EM and △ED at 300 K for all the NP samples are smaller than the corresponding △EM and △ED at 4 K. The main reason is because the thermal-expansion mismatch between the GaN and InGaN layers decreases as the temperature increases, and the corresponding strain relaxation-induced blue shift of PL peak energy decreases. Figure 7 shows the enhancement factors for total integrated PL intensities of NP samples compared to the planar sample as a function of the NP diameter at 4 and 300 K. At 4 K, the enhancement factors, decreasing with decreasing NP diameter, were about 2.5, 1.9 and 1.2 for NP320, NP250 and NP180, respectively. Under the same laser excitation density, the changes in PL intensity of NP LEDs could be due to the active area, the recombination rate and LEE. First, the filling factors (total emission area of NPs/total emission area of the planar LED) are 69%, 35% and 18% for NP320, NP250 and NP180, respectively. Second, as mentioned above, the weight of the strain relaxation area of MQW embedded in the NP LEDs increases with decreasing NP diameter. The strain relaxation weakens the piezoelectric field in QWs verified by μ-PL and μ-Raman, and enhances electron-hole wave function overlap. Therefore, thinner NP LED possesses a higher recombination rate. Third, the surface/volume ratio becomes larger as the diameter decreases, and the LEE will increase. As we know, NP LED with smaller diameter has a larger surface/volume ratio, so the light form smaller diameter NP can be emitted more easily through the sidewall than that form the larger ones. Although the NP samples sacrifice partial active area, the integrated PL intensities of NP samples are larger than that of the planar sample. It is because that the recombination rate is enlarged due to strain relaxation, and the LEE is also enlarged due to the higher surface/volume ratio of NP samples. Moreover, the enhancement factor decreased with decreasing NP diameter. This indicates that, among the factors influencing the luminous of NP LEDs with the same NP density, the loss of active area is gradually dominant with decreasing NP diameter compared to the enlarged recombination rate and LEE. As also shown in Fig. 7, the total integrated PL intensity enhancement factors at

300 K were 2.3, 1.7 and 0.9 for NP320, NP250 and NP180, respectively. The enhancement factors at 300 K were smaller than the corresponding enhancement factors at 4 K for all NP samples. This should be attributed to the residual etching defects on the sidewall of NPs that play roles as non-radiative recombination centers at 300 K [22–24]. Moreover, the enhancement factor of NP180 was smaller than that of the planar sample. The possible reason is that the NP180 possesses a higher residual surface defects density due to the larger surface/volume ratio. Finally, let's complete this section with a brief discussion. Our experiments were well designed with the expectation that this study might shed light on the fundamental understandings and useful guidance for fabricating high-efficiency strong phase-separated dual-wavelength InGaN/GaN MQW NP LEDs. To further enhance the luminous efficiency of NP LED, it is critical to simultaneously optimize NP density and NP diameter, which is beyond the scope of this paper. 4. Conclusions High-density strong phase-separated blue/green dual-wavelength InGaN/GaN MQW NP LEDs were fabricated by FIB dry etching. We have demonstrated the diameter-dependent emission characteristics of those NP LEDs using μ-PL measurement. Diameter-dependent blue shift of PL peak positions indicated that the degree of strain relaxation of MQWs embedded in the NP LEDs increased with decreasing NP diameter, verified by μ-Raman spectroscopy. Blue peak showed a larger blue shift than the green peak for all the NP samples, because the QCSE in the quasi-QDs was weaker than that in the InGaN matrix due to the smaller sizes of the strong localized quasi-QDs. Furthermore, despite smaller active volume, the NP LEDs exhibited stronger PL intensity compared to the planar sample due to the higher recombination rate and LEE. The total integrated PL intensity enhancement factors decreased with decreasing NP diameter, indicating that, among the factors influencing the luminous intensity of NP LEDs with the same NP density, the loss of active volume was gradually dominant compared to the recombination rate and LEE. And the enhancement factors at 300 K were smaller than the corresponding factors at 4 K due to the residual etching defects acting as non-radiative recombination centers. Both

the NP density and the NP diameter should be optimized, in order to further enhance the luminous efficiency of NP LED. Acknowledgments This work was supported by the Key Laboratory of Functional Crystal Materials and Device (Shandong University, Ministry of Education) (Grant No. JG1401), the Major Research Plan of the National Natural Science Foundation of China (Grant No. 91433112), the Project of Shandong Province Higher Educational Science and Technology Program (Grant No. J15LJ01), and the National Natural Science Foundation of China (Grant Nos. 51672163 and 11304172).

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Fig.1. Flowchart for the fabrication of NP LED. (a) an GaN LED structure grown on sapphire substrate using MOCVD; (b) an Au film ~100 nm was evaporated onto the GaN LED; (c) the FIB milling; (d) the Au film was removed using chemical etch, post-processing.

Fig.2. (a) Diagram of the bitmap file used for FIB patterning. Pixels in the black regions were exposed to FIB. Tilt view SEM images of InGaN/GaN NPs array with diameters of (b) 320 nm (c) 250 nm and (d) 180 nm.

Fig.3. Close-up views of room-temperature μ-Raman scattering spectra of the E2(H) phonon modes of the Planar, NP320, NP250 and NP180. The black lines are the experimental data, the green lines are the Voigt fitting curves and the red lines are the sum of the Voigt fitting curves.

Fig.4. PL spectra of four samples measured at 4 K (a) and 300 K (b), respectively.

Fig.5. Measured blue and green PL peak energies versus NP diameters at 4 and 300 K.

Fig.6. Measured peak energy shift versus NP diameters at 4 K (a) and 300 K (b).

Fig.7. Enhancement factor for total integrated PL intensity versus NP diameters at 4 and 300 K.