The fabrication process and surface treatment of the nLEDs consisting of ITO/p-GaN/MQW/n-GaN are shown in Fig. 1a–c. We used commercially grown 4” epitaxial wafers, where 8 pairs of InGaN/GaN MQW LED structures were grown on a c-plane sapphire substrate. A nanorod pattern with the diameter of about 600 nm was constructed via nano-imprint lithography, and then it was dry etched by inductively coupled plasma reactive ion etching (ICP-RIE). Subsequently, potassium hydroxide (KOH) wet etching was conducted to remove the sidewall damages from dry etching process.23,28 The profile of the nanorod is transformed from trapezoidal to vertical cylinder because of the high etching-barrier index of the m-plane surface29, and the surfaces of the nanorods become smooth (Fig. 1b). The average diameter and length of the InGaN/GaN nanorods is 580 nm and 4 μm, respectively.
Because our pixel comprised of laterally self-aligned nLEDs, the p-GaN and n-GaN should be isolated by an insulator. In addition, a single insulating layer is insufficient to simultaneously achieve the high efficiency, good reliability, and processability required for display products. Therefore, we applied two insulating layers on the nanorods. The inner and outer layers were for sidewall passivation and etch stopping during pixel processing, respectively. We used SiO2 (60 nm thickness) and Al2O3 (20 nm thickness) as the inner and outer layers, respectively, because of their high transparency, low leakage current, and high bandgap energy30,31. Optimization of the inner insulating layer, which is directly deposited on the GaN surface, is of utmost importance for achieving highly efficient nLEDs. Even when the same insulating material is used, the LED efficiency strongly depends on the process conditions of the ALD such as the substrate temperature, process time, and use of plasma. This dependence directly corresponds to the creation of VGa complex defects on the GaN surface during deposition, as will be demonstrated below.
To minimize the surface damage during the passivation process, we developed a wet-chemical process for synthesizing the SiO2 passivation layer via a sol–gel method25-27. Figure 1c and the inset image show the deposition of SiO2 and a simplified schematic of the sol–gel reaction, respectively. Performing the sol–gel process at room temperature could minimize the atomic reaction between the GaN surface and the SiO2 layer and enhance the optical properties of the nanorods by passivating the dangling bonds.
The TEM image in the inset of Fig. 1c shows a well-defined and uniform SiO2 layer on the nanorods from the top to bottom. The thickness of the layer increased linearly with an increase in the reaction time until 60 min and saturated at 23 nm owing to the limited quantity of sol particles (Extended Data Fig. 2). In addition, upon doubling the thickness of SiO2 layer up to 46 nm by repeating the sol–gel reaction, the PL intensity increased because the structural stress and surface defects of the SiO2 layer decreased with an increase in the thickness32. To evaluate the passivation uniformity, we recorded the PL spectra from three different positions on the 4” wafer (Extended Data Fig. 2). We found that the variation in the PL intensity with respect to the position is minor.
The PL images and the fluorescence excitation–emission spectra (Fig. 1d) and the panchromatic (λ=300 – 700 nm) CL images (Fig. 1e) are compared for the nanorods passivated by the different methods; conventional plasma enhanced ALD–deposited SiO2 (left) and sol–gel-deposited SiO2 (right). When the sol-gel process is applied for the SiO2 passivation, the enhancement of the blue emission over yellow emission is clearly observed in Fig. 1d. The intensity of blue emission of the sol–gel SiO2-passivated nanorods is approximately 13 times higher than that of ALD SiO2-passivated nanorods (Fig. 1f). In addition, the panchromatic CL images of the individual nanorod LEDs in Fig. 1e clearly visualizes that the existence of the non-radiative recombination region depending on the passivation method.33 The outer rim corresponds to the yellow emission from defects including VGa-related defects34,35, and the inner circle corresponds to the blue emission of the MQW regions. In the case of ALD SiO2–passivated nanorods (Fig. 1e. left), the outside of the MQWs is dark owing to the non-radiative recombination centre (NRC). Additionally, the nanorods exhibit significant differences in brightness because of the different positions of the NRC. However, for the sol-gel SiO2–passivated nanorods (Fig. 1e. right), the overall brightness is uniform and the above-mentioned abnormal light emission is absent. Furthermore, the reduced carrier lifetime of the sol–gel SiO2-passivated nanorods verifies that the surface defects of GaN are reduced than that of the ALD–deposited SiO2 (Fig. 1g).
Figure 2a shows the EL and PL composite images of a single nLED within a pixel according to the passivation method used. The pixel is composed of the nanorods connected in parallel with transparent conductive metals18. Because the EL of the laterally aligned nLED is emitted along the radial and longitudinal directions owing to waveguide reflection, a bright blue and a relatively dark emission are observed at the top and bottom of the nanorods, respectively. According to the light extraction calculation for our pixel design, the total light extraction efficiency (LEE) is approximately 25%, and it comprises 71% radial emission and 17% p-side and 10% n-side waveguide emission. The EL intensity profiles obtained from a region horizontally across the top of the nLEDs confirm that the EL intensity of the sol–gel SiO2–passivated nanorod is higher than that of the ALD SiO2–coated nanorod (Fig. 2b).
Figure 2c shows the EQE curves for the nLEDs for each surface passivation type. Each curve is obtained from a 60-pixel array, and each pixel comprises an average of 6 and 9 nLEDs for ALD SiO2- and sol–gel SiO2 passivation, respectively. The difference in the number of the nLEDs per pixel originates from the self-aligned process dispersion. The average values of the peak EQE of the sol–gel- and ALD SiO2–deposited nLEDs are 20.2% (standard deviation = 0.6%) and 8.9% (standard deviation = 0.1%), respectively. Furthermore, the IQE, which was obtained by dividing the EQE by the calculated LEE, was 81% and 36% for the sol–gel SiO2- and ALD SiO2-deposited nLEDs, respectively. The high EQE of the sol–gel SiO2-deposited nLED is remarkable because it is higher than the best EQE value of the μLED structure even at a larger diameter16. The primary reason for such a high EQE is the decrease in the GaN surface damage in the sol–gel SiO2-deposited nLED, as confirmed by the longer carrier lifetime in PL and decreased NRC region in CL in Fig. 1. Furthermore, this is observed in the J–V electrical characteristics of the LEDs.
The sol–gel SiO2–deposited nLEDs exhibit a lower leakage current than that of the ALD SiO2–deposited nLEDs at below-threshold voltages (Fig. 2d, e) owing to the parallel resistance component attributed to sidewall damage. This is consistent with the decrease in the ideality factor of sol–gel SiO2–deposited nanorods compared with that of ALD SiO2–deposited nanorods (Fig. 2f). Generally, the ideality factor of the LED structure changes significantly depending on its epi structure36; however, we have proven that the ideality factor can also be changed significantly by the passivation method. The decrease in the ideality factor of sol–gel SiO2–deposited nanorods indicates a decrease in the effective SRH recombination in the ABC model14,37, thereby increasing the radiative recombination and internal quantum efficiency to over 80%.
The significant change in the electrical and optical properties depending on the method of deposition of the side-wall insulating layer is closely related to the defects generated during the passivation process. These defects could prevent the passivation of the GaN surface dangling bonds, thereby degrading the device performance. Therefore, understanding the phenomenon occurring at the interface and controlling it are the key to manufacturing nLEDs with outstanding performance.
Figure 3 shows the analyses results of the interface between the sidewall of the nanorods and the insulator. First, we traced the evolution of the morphology and atomic structure of the sidewall in the MQW region after each fabrication step via high-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) imaging (Fig. 3a). After the dry etching of the epi layer, the nanorod developed a trapezoidal shape owing to the impact-energy difference on the surface. The white arrows indicate the amorphous region at the sidewall, which formed owing to successive ion bombardments. However, after KOH wet etching, the nanorod developed an inverted trapezoidal shape because the etch rate of p-GaN is slower than that of n-GaN38. Additionally, the semi-polar plane and m-plane appear alternately, resulting in a staircase-like sidewall morphology of the InGaN QWs. The white arrow indicates the native oxide layer on the nanorod surface. When the SiO2 layer is deposited via plasma enhanced ALD, the plasma creates point defects, such as nitrogen vacancies (VN), nitrogen interstitials (Ni), nitrogen split interstitials (N-N)N, VGa, and VGa complexes39,40, resulting in the amorphization of the surface of the nanorods. This amorphous region is observed on the entire surface of the GaN nanorod; however, it is predominantly observed on the semi-polar facet of the InGaN QWs (yellow arrows), probably owing to the high density of dangling bonds. With an increase of the thickness of the SiO2 layer from 2 nm to 60 nm (Extended Data Fig. 3e), the penetration depth of the plasma-induced amorphization increases to 4 nm at the InGaN QWs. However, when sol–gel SiO2 is deposited on the nanorods, there is no amorphous region formed at the interface in addition to the deposited SiO2.
Figure 3b, c and Extended Data Fig. 4b–g show the XPS results. The core level spectra of Ga 3d could be deconvoluted to the Ga–N, Ga–O, and Ga–Ga bonds. The Ga 3d state ratios (Fig. 3b) clearly indicate that the Ga–O bonds were the highest in the ALD SiO2–coated nanorods and the lowest in the sol–gel SiO2–deposited nanorods. Owing to the plasma-induced GaN defects, such as VN and ON, the Ga–O bonds increased, whereas the Ga–N bonds decreased. Furthermore, comparing the Si 2p core level spectrum of sol–gel SiO2 and ALD SiO2 (Extended Data Fig. 4e), the latter exhibits oxygen deficiency because of oxygen atoms contributing to the Ga–O bond. Moreover, we obtained electron spin resonance (ESR) spectra to analyse the concentration of the (N-N)N defects (Fig. 3d, e), which are known as ambipolar defects, i.e., with a deep acceptor in n-GaN and deep donor in p-GaN41. The (N-N)N0 defects decreased after the wet etching of the rods and increased after ALD SiO2 deposition. Because the defect concentrations are calculated by dividing the spins by the total weight of the nanorods and considering that most of the defects are concentrated at the nanorod surface, the actual differences in the defect concentrations of the nanorods after each fabrication step will be greater. Additionally, the (N-N)N0 defects in the sol–gel SiO2–coated rods were fewer than those in the wet-etched rod, demonstrating the passivation of surface dangling bonds without generating excess defects.
We also carried out DLTS analysis on a bulk LED chip and ALD SiO2–coated nLED–array chip (Extended Data Fig. 5). In the bulk LED chip, the main defects are electron traps owing to VN and Ni with activation energies of 0.56 and 0.67 eV, respectively42. In the nLED-array chip, although the electron trap concentration is similar to that of the bulk LED chip, five types of hole traps (0.14, 0.39, 0.51, 0.56, and 0.93 eV) are observed, and the defect concentration is greater than 1015 cm-3 for the deepest defect level (Ev+0.93 eV). VGa-related complexes (complexes of VGa and oxygen that replace nitrogen sites, i.e. VGa–ON, VGa–ON–2H) are reportedly responsible for these hole traps43,44. According to previous reports,38,45 these VGa complexes increase the SRH coefficient, thereby increasing the non-radiative recombination rate and significantly decreasing the IQE of optoelectronic devices.
Figure 4 shows a detailed comparison of two representative EEL spectra obtained from the bulk and the surface region of the InGaN QWs. For the bulk InGaN, the N-K energy-loss near-edge structure (ELNES) can be decomposed into four contributions, denoted by A, B, C, and D46. For the InGaN passivated by sol–gel SiO2, the observed spectral features of N-K ELNES are almost similar for the bulk and surface regions (Top panel in Fig 4b). Compared with those of InGaN passivated via the sol–gel method, pronounced changes are observed in the spectral features of the N-K ELNES (red line) of the surface region of InGaN passivated by plasma enhanced ALD: the appearance of the small peak between the second and third peak, and the chemical shift of the third peak, labelled as C, to a lower energy (red shift).
To theoretically verify the observed spectral features of the N-K ELNES, density functional theory (DFT) calculations were performed by considering the VGa–ON–2H complex in a 100-atom-based 3×4×2 GaN supercell, which corresponds to the most energetically favourable VGa point defect complexed with O and H44. The major changes caused by the VGa–ON–2H complex are that the third peak (C) at approximately 405 eV (bottom-most panel of Fig. 4b) is shifted toward a lower energy by approximately 0.4 eV, and a small shoulder peak appears between B and C at approximately 403 eV, which is consistent with the experimental measurements. The red shift and the formation of the shoulder peak suggest that the VGa–ON–2H complex locally breaks the wurtzite symmetry by changing the relative distance between the Ga and N atoms and then modifying the hybridization between the Ga and N 2p states, resulting in a chemical shift of the N-K edge and the formation of a shoulder peak. Therefore, we conclude that the VGa–ON–2H complex is the dominant defect generated in the sidewalls of the nanorods during the plasma ALD of SiO2, wherein a hydrogen-containing precursor is used. According to a report44, the VGa–ON–2H complex dominantly causes SRH recombination in the InGaN QWs system, and it significantly decreases the IQE of the nLEDs owing to the relaxation of the biaxial stress at the sidewall.
Although the sol–gel method provides superior opto-electrical properties, it has a few disadvantages such as the presence of residual reaction by-products, which have to be overcome to achieve good reliability. Therefore, suitable post-treatments need to be conducted to complete the reaction and eliminate the reaction residue. We have found that this can be achieved while maintaining the high EL by baking the sol–gel SiO2 film (Extended Data Fig. 6). With these preliminary results, research on further improvement via process optimization is currently under way.
In concluion, we developed highly efficient top-down-processed nLED pixels via sol–gel SiO2 passivation. In each pixel, the nanorods were aligned on pre-patterned electrodes using the dielectrophoretic force and connected in parallel via transparent electrodes. We achieved a peak EQE of 20.2% with a low leakage current at below-threshold voltages and an ideality factor of 1.54. Additionally, by performing various analyses on the surface and the interface between the nanorods and the insulating layer via DLTS, XPS, ESR, and STEM-EELS, we demonstrated that passivation via conventional plasma ALD induces amorphization at the InGaN QW surfaces and creates point defects on the sidewalls of GaN nanorods, including high concentrations of ambipolar N split interstitials and VGa-complexes, which increases the SRH recombination and generates an NRC region at the InGaN sidewall of the nanorods. Moreover, the sol–gel process is advantageous for insulating the GaN surfaces because the SiO2 nanoparticles are adsorbed on the GaN surface after the sol–gel reaction. Thus, the atomic interaction with the GaN surface is minimized and only dangling bonds of the surfaces are passivated, resulting in a low leakage current, decrease in the NRC regions, and high EQE of over 20%. We strongly believe that our findings could accelerate the implementation of nLEDs in next-generation displays.