Abnormal StranskiâKrastanov Mode Growth of Green InGaN Quantum

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Abnormal Stranski-Krastanov Mode Growth of Green InGaN Quantum Dots: Morphology, Optical Properties, and Applications in Light-Emitting Devices Lai Wang, Lei Wang, Jiadong Yu, Zhibiao Hao, Yi Luo, Changzheng Sun, Yanjun Han, Bing Xiong, Jian Wang, and Hongtao Li ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b16767 • Publication Date (Web): 06 Dec 2018 Downloaded from http://pubs.acs.org on December 8, 2018

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ACS Applied Materials & Interfaces

Abnormal Stranski-Krastanov Mode Growth of Green InGaN Quantum Dots: Morphology, Optical Properties, and Applications in Light-Emitting Devices Lai Wang*, Lei Wang, Jiadong Yu, Zhibiao Hao, Yi Luo, Changzheng Sun, Yanjun Han, Bing Xiong, Jian Wang, Hongtao Li

Beijing National Research Center for Information Science and Technology (BNRist),

Department of Electronic Engineering, Tsinghua University, Beijing 100084, P. R. China

* Address correspondence to [email protected]. KEYWORDS: quantum dots, QW-QD coupled nanostructure, morphology, tunneling injection, photoluminescence, electroluminescence

Abstract

Stranski-Krastanov (SK) growth mode is widely adopted for the self-assembled growth of semiconductor quantum dots (QDs), wherein a relatively large critical thickness is essential and a thick wetting layer (WL) is formed beneath the QD layer. In this paper, we report the metal organic

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vapor phase epitaxy (MOVPE) of green InGaN QDs, employing a growth interruption method to decrease the critical thickness and improve the morphology of QDs. The QDs exhibit similar photoluminescence properties with those grown by conventional SK mode, implying the existence of a WL. We experimentally verify that the formation of QDs, whether based on the SK mode or the growth interruption method, conforms to the phase separation theory. However, the density of QDs grown by the interruption method exhibits abnormal dependence on the strain when a quantum well (QW) is inserted beneath the QD layer. Furthermore, the underlying QW not only influences the morphology of the QDs but also plays as a reservoir of electrons, which helps enhance the photoluminescence and the electroluminescence of the QDs. The method of QD growth with improved morphology and luminescence by introducing the QW-QD coupled nanostructure is universally applicable to similar material systems. Furthermore, a 550-nm green light-emitting diode (LED) and a 526-nm superluminescent light-emitting diode (SLED) based on the nanostructure are demonstrated.

INTRODUCTION InGaN/GaN quantum dots (QDs) have shown great potential for many applications, including GaN-based light-emitting diodes (LEDs) and laser diodes (LDs).1-6 The traditional quantum-well (QW) based LEDs and LDs have been suffering from the “Green Gap” problem, which is caused by the strong piezoelectric field and the poor crystal quality of high-indium-composition InGaN.7-8 QDs have been proved as one of the most promising solutions to the “Green Gap” problem.9-12 2 ACS Paragon Plus Environment

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The formation of QDs helps improve the crystal quality as the strain is relaxed by threedimensional growth rather than dislocation formation within the QDs.13 Considering the fact that the lattice constant of QD is consistent with that of relaxed InGaN, QDs are more favorable for indium incorporation, and hence for long-wavelength light emitting. Besides, the release of strain can significantly suppress the quantum confined Stark effect (QCSE), which has long plagued InGaN QW devices.14-16 In addition, the QD-based LDs may achieve higher differential gain and lower threshold than QW-based ones as QDs have higher density of states (DOS) than QWs.17-18 QDs also enjoy many other advantages over quantum wells, including shorter carrier lifetime and weaker Auger recombination. All these advantages lead to the great prospects of InGaN QDs for green and longer wavelength light emitting devices. Stranski-Krastanov (SK) growth mode is the most widely used method to grow self-organized QDs. During the growth, the InGaN growth transforms from two-dimensional to three-dimensional growth as the thickness of InGaN reaches the critical thickness hc, and finally InGaN QDs are formed upon a wetting layer (WL).19 However, the growth of long-wavelength InGaN QDs using conventional SK mode by metal organic vapor phase epitaxy (MOVPE) requires a relatively thick WL. Furthermore, for high-indium-composition QDs growth, the strain tends to be relaxed preferentially by generating dislocations rather than three-dimensional growth.19-20 In order to solve these problems, a growth interruption method was proposed, and the thickness of the WL can be significantly reduced, compared with the case of traditional SK mode.18



Despite the many advantages described above, there remain some challenges to be conquered for self-organized growth of QDs by MOVPE: namely, the nonuniformity of both the size and the distribution of QDs, as well as the low density of QDs. The gap between QDs will provide leakage path for carriers, especially when the QD density is low. Furthermore, the relaxation time of hot carriers from the continuous state to the quantum state in QDs is longer than that in QW, which will also increase the probability of carrier leakage.21-23 One of the methods to enhance the carriercapture capability of QDs is to increase the number of QD layers. But this will result in inhomogeneous carrier distribution and lead to increased absorption loss in LDs.24 A novel tunneling injection nanostructure was proposed to suppress the carrier leakage.25 The structure consists of two coupled parts: a blue InGaN QW layer and a green InGaN QD layer, separated by a tunneling barrier. The QW beneath the QD layer is carefully nanoengineered so that the QCSE effect is strong enough to reduce the recombination rate of carriers, thus carriers can be injected into the QD layer by phonon-assisted tunneling and recombine there.26 In this work, we experimentally verify that the QD formation in both the SK mode growth and the growth interruption method is closely related to the phase separation process of InGaN (the In composition becomes inhomogeneous and agglomerate as the InGaN grows).27-28 We report the MOVPE growth of the QW-QD coupled nanostructure, using the growth interruption method.20 The underlying QW can affect the strain in the QD layer by adjusting the thickness of the GaN tunneling barrier. It is found that the formed QDs exhibit similar optical properties with the ones grown by SK mode. However, atomic force microscope (AFM) measurement reveals an abnormal 4 ACS Paragon Plus Environment

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relationship between the QD density and the strain, compared to the traditional view.

29-32

This

abnormal phenomenon is explained unambiguously by the spinodal decomposition of InGaN during growth interruption.33 In addition, the QW-QD coupled structure is adopted as the active region of light-emitting diodes and proved effective for suppressing carrier leakage. A green superluminescent light-emitting diode (SLED) is also demonstrated based on the nanostructure to show its potential in long wavelength visible light-emitting devices. RESULTS AND DISCUSSION Dynamics of QD Formation. Generally, QDs are considered as the result of strain relaxation during the epitaxy of InGaN. For the SK-mode, two-dimensional growth converts to threedimensional growth to form QDs when the InGaN layer reaches a critical thickness, due to the strain accumulation originating from the lattice mismatch. However, we found that it is probably the decomposition and re-incorporation of InGaN that leads to the formation of QDs. Early studies generally suggest that the InGaN decomposition is mainly attributed to the small In-N bonding energy34, which allows the In-N bond to easily decompose and migrate, so the indium composition of InGaN QDs is higher than that of WL. But recent studies have shown that the decomposition region of InGaN may be a metastable or unstable region formed by phase separation.33 In order to understand the mechanism of QD formation, a series of InGaN samples with different thicknesses and different In compositions (by varying the NH3 flow) were grown by the conventional SK mode. The surface morphology of the InGaN samples, illustrated in Figure 1b, fits well with the theoretically calculated diagram of the InGaN phase separation by Tessarek, C. 5 ACS Paragon Plus Environment


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et al., as shown in Figure 1a.33 The similarity between the morphologies of these samples and the calculated results of phase separation indicates that the InGaN actually undergoes decomposition, migration, and re-incorporation to form QDs during the growth process, implying that the growth of InGaN QDs by SK mode is consistent with the theory of phase separation. More specifically, as the thickness of InGaN increases, the quantity of QDs increases significantly, with the QDs mainly distributed in the pits and much less in the flat area, as shown in Figure 1b.2\4\5. Figure 1c shows the magnified image of the local area marked by a red dotted box in Figure 1b.2, and it can be clearly seen that QDs are formed in the pits. This preferential formation of QDs is due to the inhomogeneity of the In composition in the InGaN layer (phase separation), causing these regions to first reach the critical thickness and form QDs. According to Figure 1a, phase separation intensifies as InGaN thickens. Therefore, as the thickness of InGaN increases, the number of region with high In composition increases, thus the number of QDs increases. On the other hand, increasing the NH3 flow can suppress In-N decomposition and In desorption, thereby increasing the In composition. As shown in Figure 1b.1-3, as the In composition of InGaN increases, the phase separation intensifies, and the pits generated by InGaN decomposition increase. In addition, the number of QDs in the pit increases dramatically as the NH3 flow increases. This indicates that during the QD formation process, the decomposed In atoms are partially desorbed, and increasing the NH3 flow inhibits the desorption, causing more In atoms to coalesce


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into the pits. This also confirms that the pits are caused by InGaN decomposition, and QDs are formed by subsequent recrystallization.

Figure 1. (a) Theoretically calculated diagram showing the regions with different In composition formed by InGaN phase separation, where the black regions indicate high In composition. (Reproduced with permission from ref. 33. 2011 American Physical Society.①)

①

(b) Surface

https://doi.org/10.1103/PhysRevB.83.115316 7 ACS Paragon Plus Environment


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morphology of InGaN samples on 3 m GaN with different thickness and In composition (by varying the NH3 flow) without growth interruption, which means the QDs are grown by SK mode. The green ellipses in AFM images highlight the typical QDs. Increasing the NH3 flow during growth results in Images 1-3, while increasing the InGaN thickness leads to Images 2, 4, and 5. (c) Enlarged image of the area marked by the red dotted box in Figure 1b.2, which contains two QDs. (d) and (e) AFM image and scanning electron microscopy (SEM) image of QD samples grown by growth interruption method, respectively. Regions with similar morphology are formed after InGaN decomposition, as marked by the red ellipses. The green ellipses mark the typical QDs, formed after the decomposition of InGaN. In order to enhance the decomposition and recrystallization of the unstable and semi-stable regions after phase separation, facilitate QD formation, and decrease the thickness of the WL in SK mode, a growth interruption method is introduced. The growth interruption method is a modified SK mode.35 It is believed that during the growth interruption, the InGaN surface atomic decomposition will intensify, and atoms can migrate and recombine on the surface to form stable QDs due to the breaking of the adsorption-desorption equilibrium. Figure 1d-e show the AFM and SEM images of typical QD samples obtained with the growth interruption method. There are many gully regions arising from the decomposition of InGaN around the QDs, which are very similar to the results in Figure 1a-b. In Figure 1d, obviously, QDs appear in pitted areas, which also implies that QDs are derived from the decomposition of InGaN, exactly the same as SK mode. The growth 8 ACS Paragon Plus Environment

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interruption method provides a generic solution for growing QDs in an InGaN-like phase separation system. According to the phase separation theory, strain can affect the growth process of QDs by affecting the size of the spinodal region, and thus it is theoretically feasible to control the QD morphology via strain-controlled phase separation.33 Three samples, labeled as A, B and C, were grown by MOVPE, with the epitaxial structure shown in Figure 2.

Sample A contains pure QD

active region, while Samples B and C include a 4.5 nm InGaN blue QW beneath the QD layer. The GaN tunneling barrier between QD layer and QW in Samples B and C are 4.5 nm and 1.5 nm, respectively. The InGaN QDs were grown using a two-step growth interruption method, similar to our previous reports.20,

36-37

Details of the growth is given in the METHODS AND

EXPERIMENTAL SECTION. The surface morphologies of Samples A, B and C were measured by a Bruker Dimension Icon AFM, as shown in Figure 2d-f, respectively. According to the AFM images, the width and height of the QDs are around 70 and 20 nm, respectively, which are typical sizes of uncapped QDs grown by growth-interruption method.13, 20, 36-37 The densities of QDs are 2.8×108, 1.4×109, and 2.7×109 cm-2 for Samples A, B, and C, respectively, determined by counting the number of QDs in the AFM images.



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Figure 2. Schematic illustration of the epitaxial structure for Samples (a) A, (b) B, and (c) C. The QW-QD coupled tunneling structure consists of a 4.5 nm InGaN QW, a 1.5 nm (4.5 nm) GaN barrier, and a green InGaN QD layer. The surface morphologies of Samples (d) A, (e) B and (f) C measured by AFM, and the densities of QDs are 2.8×108, 1.4×109, and 2.7×109 cm-2 for Samples A, B, and C, respectively. As the GaN tunneling barrier between QDs and QW in Samples B and C are 4.5 and 1.5 nm, respectively, the strain in QD layer is strongest in Sample A and weakest in Sample C. The strain here refers to the strain of the samples before the QDs begin to grow but after the growth of the 10 ACS Paragon Plus Environment

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GaN tunneling barriers. In the traditional view, the QD density should be determined by strain of the epitaxial layer, with larger strain resulting in higher QD density.29-32

In our case, the thinner

the GaN barrier between the InGaN QW layer and the QD layer is, the smaller the strain in InGaN QD layer will be. Consequently, the highest QD density should be expected in Sample A while the lowest in Sample C. However, our experimental results are contrary to that expectation. This abnormal phenomenon implies the principle of QDs formation in growth interruption method is fully different with the traditional strain-induced view, even though similar WL and QDs are achieved. This abnormal dependence of the QD density on strain can be explained by phase separation theory. The spinodal region desorbs and provides the necessary atoms to form QDs during the growth interruption, which can be affected by indium composition, temperature, and strain state. As the thickness of the barrier between the QW layer and the QD layer decreases, the strain decreases, and thus the unstable spinodal region increases.33 Therefore, the increased QD density in Samples B and C can be attributed to the enhancement of phase separation due to strain relaxation of the QD layer induced by the InGaN QW. The enhanced phase separation of InGaN increases the unstable spinodal region, leading to the increased QD density. The introduction of the nanostructure to adjust the morphology of QDs is believed applicable to all phase separation material systems. Optical Properties of the Nanostructure. In order to gain a better understanding of the influence of the nanostructure on the optical properties of QDs, three samples: A1, B1, and C1, 11 ACS Paragon Plus Environment


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were grown. The active region of Sample A1 consists of 5 periods of InGaN QDs without tunneling heterostructures, as shown in Figure 2a. For Samples B1 and C1, the active region consists of 5 periods of InGaN QDs with tunneling heterostructures shown in Figure 2b and Figure 2c, respectively. Details of the growth can be found in the METHODS AND EXPERIMENTAL SECTION. Figure 3 demonstrates the bright field transmission electron microscope (BFTEM) images of the three samples, which indicates the actual thickness of each layer closely matches the designed value. The thickness of the InGaN QW layer is 4.8 nm for Sample B1 and 4.6 nm for Sample C1, while the thickness of the barrier layer is about 1.5 nm for Sample B1 and 4.5 nm for Sample C1, and the thickness of the upper barrier layers is around 11.5 nm. QD layers can be clearly identified in the BETEM images. The height of the QD fluctuates around 2.4 nm, while the lateral size varies from 20 to 50 nm.

Figure 3. Bright field transmission electron microscope (BFTEM) images for samples A1 (a), B1 (b), and C1 (c), respectively. The darker region represents InGaN, while the brighter region


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represents GaN. The actual thickness of each layer of the three samples is in good agreement with the designed ones. Detailed optical property characterization of the structure was performed by temperaturedependent photoluminescence (TDPL) measurement of the three samples, as shown in Figure 4ac. For Samples A1 and C1, as the temperature rises from 8 to 300 K, only the luminescence from the QD layers can be observed. The spectra consist of two peaks. The one with a shorter wavelength can be attributed to the emission from the InGaN WL, while the other peak with a longer wavelength is attributed to the emission from the InGaN QDs.38-40 For Sample B, an extremely intense peak centering around 430 nm together with one LO phonon replica can be observed at relatively low temperatures (below 150K), which is referred to the luminescence of the QWs. The TDPL spectra reveal that carriers can tunnel from the QW into the QDs when the GaN barrier in the nanostructure is sufficiently thin. Figure 4d-f demonstrate the variation of the center wavelength, the full width at half maximum (FWHM), and the integrated intensity of the QD emission peak of Samples A1 and C1 with temperature, respectively. As the temperature rises, the center wavelength and the FWHM of the QDs peak of both samples decrease slightly and then begin to increase, while the integral intensity shows the opposite trend. These trends exhibit good agreement with the optical properties of the QDs grown in the SK mode, implying that the growth interruption method is just a modified SK



mode.

41-44

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The behavior of Sample B1 is not compared, as the existence of the QW peak at

relatively low temperatures (below 150K) complicates the luminescence of QDs.

Figure 4. TDPL spectra of Samples A1 (a), B1 (b), and C1 (c). As the temperature rises from 8 to 300 K, only QD emission can be observed for Samples A1 and C1, while intense QW luminescence peak exists in Sample B at relatively low temperatures (below 150K). The temperature dependence 14 ACS Paragon Plus Environment

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of (d) the center wavelength, (e) FWHM, and (f) the integral intensity of the QD peaks of Samples A1 and C1. The TDPL measurement of the three samples at 8 and 300 K are specially shown in Figure 5a and b, respectively. As Figure 5a shows, two peaks are found in all spectra of the three samples at 300 K. The peak with a shorter wavelength is derived from the InGaN WL and the other peak with a longer wavelength is emission from the InGaN QDs. According to the phase separation theory, the inserted InGaN QW, which can partially relax the strain, enhances the InGaN decomposition and decrease the area of WL. This explains the decrease of PL intensity of WL with the increase of QD density upon comparing the TDPL spectra of the three samples in Figure 5a.



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Figure 5. Normalized PL spectra of Samples A1 (green), B1 (red) and C1 (blue) measured at (a) 300 K and (b) 8 K, respectively. The inset in (a) is a schematic diagram of the conduction band lineup of the QW-QD coupled nanostructure. The normalization method we use is dividing the spectra of samples A1, B1 and C1 by their peak intensities, respectively, so that all the peaks are equal to 1. However, as shown in Figure 5b, the normalized spectra of these three samples at 5 K is quite different from those obtained at 300 K. This phenomenon indicates there are more physical mechanisms that affect the luminescence, including the tunneling injection of carriers from the QW and the quantum confined Stark effect (QCSE). To begin with, the intense peak centering around 435 nm in the low-temperature spectrum of Sample B1 can be identified to be the emission from the InGaN QW by comparing the spectra of Samples A1 and B1. Hence, the absence of the QW emission peak in the spectrum of sample C1 can be only explained by the enhancement of tunneling injection due to the decreased GaN barrier thickness. Meanwhile, Sample C1 exhibits a blue shift of the QDs peak compared with Sample A1. The blue shift may be caused by the increase of carrier density in QDs due to the tunneling injection, while the contribution of strain relaxation of QDs induced by the introduction of InGaN QW should also be taken into account. The internal quantum efficiencies (IQE) of the QD emission in Samples A1, B1 and C1 are 9.8%, 11.6% and 12.4%, respectively, where the IQE is calculated as the ratio of the integrated PL intensity


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measured at 300 K to the one measured at 8 K (IQE = I300K / I8K), and a Gaussian fitting is adopted to determine the integrated intensity of the QD and the WL peaks individually. In our previous study, TRPL spectra have confirmed the existence of carrier tunneling injection in the nanostructure.45 One typical phenomenon is that as the carriers tunnel into the QDs, the carrier lifetime of the coupled QW gets shorter, whereas that of the coupled QDs gets longer. Thus, the increase of IQE of Samples B and C can be attributed to the combined effect of the tunneling injection and the weakening of QCSE (the strain can be partially relaxed by the inserted InGaN QW). Applications in Light-emitting Devices. GaN-based light-emitting devices suffer from two obstacles for a long time. One is the “Green Gap”,

8-9, 46

and the other is the “droop” issue (the

“droop” refers to the phenomenon that the efficiency of InGaN QW LEDs deteriorates dramatically as the injection current increases, for which one of the most widely known reasons is the leakage current). 47-48 In traditional c-plane long-wavelength InGaN QWs, strong polarization field causes severe QCSE effect which seriously constrains the IQE. In addition, high crystal quality epitaxy for high In-component InGaN QWs has always been a major road blocker. The QW-QD coupled tunneling injection structure is expected to help alleviate both problems. As described in the Introduction, the crystal lattice in QDs is fully relaxed, which leads to improved crystal quality and weakened QCSE, which helps enhance the efficiency for green-yellow illumination. Furthermore, the unique nanostructure is expected to relieve the leakage current of



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the device, wherein the QW beneath the green QD layer plays a role as a reservoir of electrons, thus facilitating electron tunneling into the QDs. LEDs are fabricated using Samples A1 and C1 to study the influence of the QW-QD coupled nanostructure on the transportation of electrically injected carriers and to demosntrate the potential applications of the nanostructure in solid state lighting. The schematic LED structure based on Sample C1 is shown in Figure 6. The LED structure based on Sample A1 is exactly the same, except for the active region.

Figure 6. Structure diagram of the LEDs. Apart from the active region, Samples A1 and C1 have the same structure and are fabricated by the same process. P-electrodes are evaporated on p-GaN, and the samples are etched down to n-GaN to fabricate n-electrodes. Figure 7a shows the emission wavelength variation of Samples A1 and C1 as the injection current increases from 2 to 100 mA. The much smaller blue-shift of Sample C1 implies that QCSE in the sample is much weaker, which can be attributed to the strain relaxation due to the insertion of the InGaN QW. The integrated intensity of electroluminescence (EL) of Samples A1 and C1 as a function of the injection current at room temperature is plotted in Figure 7b. It is evident that the 18 ACS Paragon Plus Environment

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EL is significantly enhanced by introducing the coupled nanostructure, as a result of suppressed QCSE and enhanced capture capability of the QDs via quantum tunneling.

Figure 7. (a) Peak wavelength variation of Samples A1 and C1 as a function of injection current at room temperature. The much smaller blue-shift of sample C1 implies that the QCSE of sample C1 is much weaker, compared to sample A1. The inset reveals an improved IV characteristics for Sample C1. (b) The integrated EL intensity of Sample A1 (black) and Sample C1 (red) at room temperature. The inset shows the spectra of Sample C1 as the injection current increases from 1 to 10 mA. (c) Contrastive photos of Samples A1 and C1 under different injection currents. Figure 8a and Figure 8b show the temperature-dependent EL spectra of Samples A1 and C1, respectively, as the temperature ranging from 5 to 200 K under a constant injection current of 100 mA. At 5 K, as shown in Figure 8c, the EL spectra of Samples A1 and C1 both contain several peaks. By comparing with the PL spectra in Figure 5, the peaks around 480 nm and 570 nm are 19 ACS Paragon Plus Environment


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identified to the emission from the WL and the QDs, respectively. Meanwhile, no corresponding emission in the range of 370 ~ 450 nm was found in the PL spectra. In Figure 8d, the EL spectra ranging from 370 to 450 nm are re-plotted as a function of photon energy, with the emission peaks locating at 3.29, 3.20, 3.11 and 3.01 eV. According to previous reports, the emission peak near 3.29 eV is attributed to the radiative recombination of donor-acceptor pairs in GaN, while the 3.20 eV and 3.11 eV peaks are the first and second order phonon replica, respectively, and the peak near 3.01 eV was reported for Mg-doped GaN.49-52 Therefore, the EL peak in the 370 ~ 450 nm range is originated from the p-GaN, which indicates that part of the carriers recombine in the pGaN. Considering the difficulty of hole activation at low temperature, the surplus electrons in the active region can escape from the active region into the p-GaN region and subsequently recombine with holes. Therefore, the emission around 370 ~ 450 nm is caused by carrier leakage and will be referred to as the ‘leakage peak’.


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Figure 8. Temperature-dependent EL spectra of (a) Sample A1 and (b) Sample C1. (c) Lowtemperature EL spectra and (d) the leakage peaks of Samples A1 and C1 at 5 K. The leakage peak of Sample C1 is weaker, which can be attributed to the effectively suppressed leakage current, due to the introduction of the tunneling structure. The temperature-dependent EL spectra of Sample C1 confirms the existence of carrier leakage process. In Figure 8b, when the temperature increases from 5 to 70 K, the leakage peak and the QD peak is enhanced simultaneously. Such enhancement of luminescence can be attributed to the increase of hole concentration due to thermal activation of impurities. It also suggests that some electrons escaped directly without recombination in the p-GaN at low temperature. As the 21 ACS Paragon Plus Environment


temperature increases, the leakage peak weakens and then disappears. This can be attributed to the suppressed carrier leakage as more holes are activated and injected into the active region, as well as the increased nonradiative recombination at higher temperatures. Consequently, as the temperature increases, the intensity of QD luminescence enhances and reaches maximum around 160 K. When the temperature is over 160 K, the intensity of QD luminescence begins to decrease due to the increased nonradiative recombination, which is similar to the temperature-dependent PL measurement. Meanwhile, a blue shift of the QD peak is observed as the temperature increases from 5 to 160 K, which is attributed to the band filling effect and QCSE shielding due to the carrier concentration increase in the active region. As the temperature continues to increase, a small red shift of the QD peak is observed, which can be explained by band-gap shrinkage effect. In Figure 8a, the temperature-dependent EL spectra of Sample A1 shows a similar but more complex pattern. It is mainly because the WL of Sample A1 exhibits a stronger luminescence, which means carrier transportation between the WL and the QDs complicates the carrier transportation and recombination in the LED. As shown in Figure 8c, the introduction of the QW-QD coupled structure help suppress carrier leakage, leading to a weakened leakage peak. Apart from improving the performance of green LEDs, the nanostructure is also expected to be used in more advanced applications, such as SLEDs or LDs. SLEDs are edge-emitting devices offering wide spectral bandwidth and good coupling to external devices at the same time. As spatial coherent but time incoherent light sources, SLEDs have outstanding feature that combines the advantages of both LEDs and LDs, making them ideal light sources for fiber-optical gyroscopes 22 ACS Paragon Plus Environment

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(FOG), optical coherence tomography (OCT), fiber-optic sensors and equipment testing (FOT).5354

Currently, GaN-based LEDs suffer from the “efficiency droop” effect, leading tosignificant decrease of efficiency at high current density injections, and the small modulation bandwidth of LEDs also limits their applications in visible light communication (VLC). Meanwhile, visible LDs for laser display are associated with serious speckle noise and safety concerns. On the other hand, SLEDs combine the broad spectrum of LEDs and the high power density, beam polarization, and beam directionality of LDs. Thus, a visible SLED would be a wonderful light source for VLC and pico-projection applications. So far, there are a few reports on SLEDs in the UV or blue range, but green SLEDs with wavelength beyond 500nm are rarely reported.55-57 We demonstrate a green SLED emitting at 526 nm based on the InGaN QW-QD tunneling structure. QDs are adopted as the active region since they are expected to realize a much broader optical spectrum than QWs, and better temperature characteristics as well.58-60 The structure of the SLED is shown in Figure 9a.



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Figure 9. (a) Schematic of the SLED and its (b) emission spectra under different injected currents. The inset shows the polarization of the emitted light. (c) Power-current (PI) curve of the SLED. The insets are photos of the SLED viewed from different perspectives under an injection current of 40 mA. The spectra of the SLED are plotted in Figure 9b. At 10 mA, the center wavelength is around 510 nm, with an FWHM of 51 nm. When the current increases, the emission around 526 nm increases rapidly, together with a decrease of the FWHM. At 70 mA, the 526 nm peak becomes 24 ACS Paragon Plus Environment

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dominant, with an FWHM of 26 nm. Further increase of the injection current results in increased emission intensity, but the FWHM remains almost constant and the peak exhibits a slight blueshift. The inset in Figure 9b shows that the light from the SLED is linearly polarized, with the polarization direction perpendicular to the c-axis, i.e. the growth direction. The degree of polarization (DOP), defined as

𝐼0° ― 𝐼90°

𝐼0° + 𝐼90°,

is found to be as high as 78.7%.

Figure 9c shows the power-current curve of the SLED. At low current, only spontaneous emission exists, and the power increases linearly with the injection current. When the current is over 40 mA, the power begins to increase super-linearly. At 70 mA, the power of light exiting from one edge of the SLED reaches 1.06 mW. METHODS AND EXPERIMENTAL SECTION MOVPE Growth of Samples. All of the samples are grown on 2-inch (0001) sapphire substrates in an AIXTRON 2000HT MOVPE system. Samples A, B, and C all contain a 3-μmthick undoped high-temperature GaN layer (HT GaN, grown at 1035 C) , a 10-nm-thick undoped low-temperature GaN layer (LT GaN, grown at 740 C), and 2 periods of QD layers (grown at 650 C by the interruption method) with the top QD layer uncapped for morphology measurement. Sample A contains pure QDs active region, while Samples B and C include a 4.5 nm InGaN blue QW beneath the QD layer. The GaN tunneling barrier between QDs and QW in Samples B and C are 4.5 nm and 1.5 nm, respectively. The underlying InGaN QW is grown under 740 oC, with an estimated In concentration of 12%, and the flow rates of triethylgallium (TEGa) and



trimethylindium (TMIn) during growth are 67 sccm and 60 sccm, respectively. A flow chart for the growth process could be found in the Supporting Information. The epitaxial structure of Samples A1, B1, and C1 includes a 4-µm-thick Si-doped GaN layer grown at 1035 C, an n-type 10-pairs of In0.03Ga0.97N (3 nm)/GaN (3 nm) superlattices (SL) layer grown at 740 C, a 5-periods active region, a 30-nm-thick Mg-doped Al0.1Ga0.9N layer grown at 900 C, and a 150 nm Mg-doped GaN layer grown at 900 C. The active region of Sample A1 consists of 5 periods of InGaN QDs without tunneling heterostructures, as shown in Figure 2a. For Samples B1 and C1, the active region consists of 5 periods of InGaN QDs with tunneling heterostructures shown in Figure 2b and Figure 2c, respectively. Flow charts for the growth process and the LED fabrication process (for A1 and C1) can be found in the Supporting Information. The SLED consists of a 1 μm n-GaN layer grown at 1035 C, a 750 nm n-Al0.08Ga0.92N layer grown at 1000 C, a 150 nm n-In0.03Ga0.97N waveguide layer grown at 740 C, 5 periods of the nanostructure figured in Figure 2c, a 10 nm p-Al0.3Ga0.7N EBL grown at 900 C, a 150 nm p-GaN waveguide layer grown at 900 C, a 500 nm p-Al0.08Ga0.92N layer grown at 900 C, and a 200 nm p-GaN contact layer grown at 900 C. A flow chart for the growth and fabrication process is shown in the Supporting Information. PL and EL Measurement. The PL spectra are measured in a cryogenic chamber with a temperature ranging from 8 to 300 K (the chamber is cooled by a liquid helium compressor). A Coherent 405-nm laser with a power density of 4 W/cm2 is used as the excitation source, which 26 ACS Paragon Plus Environment

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can excite carriers in the QWs and the QDs only while avoiding stimulating carriers in the GaN barriers. Temperature-dependent EL spectra of Samples A1 and C1 are measured with a Lakeshore CRX-4K cryogenic probe station (cooled by a liquid helium compressor) with a temperature ranging from 5 to 200 K. A Keithley 2430 digital source-meter and an Ocean Optics HR4000 spectrometer are both controlled by a programmed computer to carry out the EL measurement. CONCLUSIONS In summary, we experimentally verify that the QD formation in both the SK mode growth and the growth interruption method is closely related to the phase separation process of InGaN. Furthermore, a QW-QD coupled tunneling injection nanostructure is introduced. The QD density exhibits abnormal dependence on the strain in the nanostructure, and this is explained by the strainregulated phase separation of InGaN. The nanostructure shows an improved IQE due to the tunneling effect and the weakening of QCSE. Green InGaN LEDs based on the QW-QD coupled nanostructure are fabricated, which exhibit enhanced EL and suppression of carrier leakage, thanks to the enhanced electron capture of QDs due to tunneling injection. A 526-nm green InGaN SLED based on the QW-QD tunneling structure is also demonstrated, with a degree of polarization as high as 78.7%. ASSOCIATED CONTENT Supporting Information



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The Supporting Information is available free of charge. Flow chart for the MOCVD growth and fabrication process of samples A, B, and C, flow chart for the growth process of samples A1, B1, and C1 and LED fabrication process of samples A1 and C1, and flow chart for the growth and fabrication process of SLED sample. (PDF) AUTHOR INFORMATION Corresponding Author *Tel.: +86 10 62798240. E-mail: [email protected] (Lai Wang) ORCID Lai Wang: 0000-0001-7262-0020 Lei Wang: 0000-0003-0418-7573 Author Contributions The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. Notes The authors declare no competing financial interest. ACKNOWLEDGMENTS


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The authors are grateful to Prof. Jianping Liu in Suzhou Institute of Nano-Tech and NanoBionics, Chinese Academy of Sciences for the fabrication of SLED. The authors are also highly indebted to the National Key Research and Development Plan (Grant No. 2016YFB0401803), S&T Challenging Project (Grant No. 2016003), the National Natural Science Foundation of China (Grant Nos. 61574082, 61621064, 61822404, and 51561165012), China Postdoctoral Science Foundation (Grant No. 2018M640129), and Tsinghua University Initiative Scientific Research Program (2015THZ02-3).

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the TOC graphic



Figure 1. (a) Theoretically calculated diagram showing the regions with different In composition formed by InGaN phase separation, where the black regions indicate high In composition. (Reprinted from Ref [33], with the permission of APS.) (b) Surface morphology of InGaN samples on 3 μm GaN with different thickness and In composition (by varying the NH3 flow) without growth interruption, which means the QDs are grown by SK mode. The green ellipses in AFM images highlight the typical QDs. Increasing the NH3 flow during growth results in Images 1-3, while increasing the InGaN thickness leads to Images 2, 4, and 5. (c) Enlarged image of the area marked by the red dotted box in Figure 1b.2, which contains two QDs. (d) and (e) AFM image and scanning electron microscopy (SEM) image of QD samples grown by growth interruption method, respectively. Regions with similar morphology are formed after InGaN decomposition, as marked by the red ellipses. The green ellipses mark the typical QDs, formed after the decomposition of InGaN. 177x147mm (300 x 300 DPI)

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Figure 2. Schematic illustration of single period of the active regions of Samples (a) A, (b) B and (c) C. The QW-QDs coupled tunneling structure consists of a 4.5 nm InGaN QW, a 1.5 nm (4.5 nm) GaN barrier, and a green InGaN QD layer; The surface morphologies of samples (d) A, (e) B and (f) C measured by AFM, and the densities of QDs are 2.8×108 cm-2, 1.4×109 cm-2 and 2.7×109 cm-2 for samples A, B and C, respectively. 165x131mm (299 x 299 DPI)



Figure 3. The bright field transmission electron microscope (BFTEM) images for sample A1 (a), B1 (b), and C1 (c), respectively. The darker place represents InGaN, while the brighter place represents GaN. The actual thickness of each layer of the three samples is in good agreement with the designed ones. 165x54mm (299 x 299 DPI)


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Figure 4. The TDPL spectra of samples A1 (a), B1 (b), and C1 (c), respectively. As the temperature rises from 8 to 300 K, only the QDs of samples A1 and C1 emit light (The luminescence consists of the peak of the WL of about 480 nm and the QD peak of 550 nm), while the luminescence peak of the QWs exists in sample B simultaneously at relatively low temperatures. As the temperature changes, the center wavelength, the full width at half maximum (FWHM), and the integral intensity of the QD peaks of samples A1 and C1 are shown in (d)-(f), respectively. 165x165mm (299 x 299 DPI)



Figure 5. Normalized PL spectra of samples A1 (black line), B1 (red line) and C1 (blue line) measured at (a) 300 K and (b) 8 K, with a schematic structure of the conduction band lineup. 85x117mm (299 x 299 DPI)


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Figure 6. Structure diagram of the LEDs. Apart from the active region, Samples A1 and C1 have the same structure and are fabricated by the same process. P-electrodes are evaporated on p-GaN, and the samples are etched down to n-GaN to fabricate n-electrodes. 85x47mm (300 x 300 DPI)



Figure 7. (a) Peak wavelength variation of Samples A1 and C1 as a function of injection current at room temperature. The much smaller blue-shift of sample C1 implies that the QCSE of sample C1 is much weaker, compared to sample A1. The inset reveals an improved IV characteristics for Sample C1. (b) The integrated EL intensity of Sample A1 (black) and Sample C1 (red) at room temperature. The inset shows the spectra of Sample C1 as the injection current increases from 1 to 10 mA. (c) Contrastive photos of Samples A1 and C1 under different injection currents.


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Figure 8. Temperature-dependent EL spectra of (a) Sample A1 and (b) Sample C1. (c) Low-temperature EL spectra and (d) the leakage peaks of Samples A1 and C1 at 5 K. The leakage peak of Sample C1 is weaker, which can be attributed to the effectively suppressed leakage current, due to the introduction of the tunneling structure.



Figure 9. (a) Schematic of the SLED and its (b) emission spectra under different injected currents. The inset shows the polarization of the emitted light. (c) Power-current (PI) curve of the SLED. The insets are photos of the SLED viewed from different perspectives under an injection current of 40 mA.


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Abnormal StranskiâKrastanov Mode Growth of Green InGaN Quantum

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