Comparison of structural and optical properties of blue emitting In0.15Ga0.85N/GaN multi-quantum-well layers grown on sapphire and silicon substrates

Six periods of 2-nm-thick In0.15Ga0.85N/13-nm-thick GaN blue emitting multi-quantum-well (MQW) layers are grown on sapphire (Al2O3) and silicon (Si) substrates. X-ray diffraction, Raman spectroscopy, atomic force microscopy, temperature-dependent photoluminescence (PL), Micro-PL, and time-resolved PL are used to compare the structural and optical properties, and the carrier dynamics of the blue emitting active layers grown on Al2O3 and Si substrates. Indium clustering in the MQW layers is observed to be more pronounced on Al2O3 than those on Si as revealed through investigating band-filling effects of emission centers, S-shaped peak emission energy shifts with increasing temperature, and PL intensity-peak energy spatial nonuniformity correlations. The smaller indium clustering effects in MQW on Si are attributed to the residual tensile strain in the GaN buffer layer, which decreases the compressive strain and thus the piezoelectric polarization field in the InGaN quantum wells. Despite a 30% thinner total epitaxial thickness of 3.3 μm, MQW on Si exhibits a higher IQE than those on Al2O3 in terms of internal quantum efficiency (IQE) at temperatures below 250 K, and a similar IQE at 300 K (30% vs 33%). These results show that growth of blue emitting MQW layers on Si is a promising approach compared to those conventionally grown on Al2O3. © 2019 Author(s). All article content, except where otherwise noted, is licensed under a Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/). https://doi.org/10.1063/1.5078743


I. INTRODUCTION
Indium gallium nitride (InGaN)/gallium nitride (GaN) multiple quantum well (MQW)-based blue light emitting diodes (LEDs) and their white LED derivatives have transformed the lighting industry thanks to their superior power efficiency, longevity, and compactness. The lack of commercially viable native GaN substrates necessitates InGaN/GaN LEDs to be grown on foreign substrates such as sapphire (Al 2 O 3 ). 1 LEDs can also be grown on Si(111) substrates, which are widely available in very high quality and wafer sizes as large as 450 mm at lower than one-fifth of the price per unit area. In the latter, the absorptive silicon substrate can then be removed after a flip-chip bonding scheme is employed. This approach can significantly reduce the cost and increase the market penetration of LED. 2 In addition, recent advances in III-nitrides growth on silicon have mitigated the thermal expansion and lattice mismatch issues with the use of stress management layers such as AlN buffer layers and SiN interlayers. [3][4][5] Recent studies on electrically-injected and packaged LEDs showed instances where LED on Si outperforms those on Al 2 O 3 . 6,7 However, such studies have not been done for ARTICLE scitation.org/journal/adv LEDs on Al 2 O 3 and Si substrates with identical blue emitting MQWs, inhibiting direct comparation of structural and optical properties of the blue emitting MQWs. In this work, identical blue-emitting MQW layers without p-GaN caps are grown on c-plane Al 2 O 3 (MQW-Al 2 O 3 ) and Si(111) (MQW-Si) substrates. Structural studies using X-ray diffraction (XRD), Raman spectroscopy, and atomic force microscopy (AFM), and optical studies using Microphotoluminescence (µPL), temperature-dependent PL (TDPL), time-resolved PL (TRPL) are carried out in order to compare the structural and optical properties of MQWs on Al 2 O 3 and Si substrates.  respective averages using a continuous wave HeCd 325 nm laser. MQW-Al 2 O 3 and MQW-Si exhibit integrated intensity uniformities of 5.1% and 4.1%, and average peak emission energies of 2.791 eV ± 3.7 meV (444 nm) and 2.823 eV ± 1.4 meV (439 nm), respectively. One noteworthy observation is that the strongly blue-shifted regions of MQW-Al 2 O 3 coincide accurately with the high PL intensity regions, and conversely the red-shifted clusters overlap with the low PL intensity spots. Closer inspection of the spectra of these locations shows that the brighter regions emit additional photons only in the high energy shoulder when compared to the dimmer regions. This is attributed to the screening of the quantum confined stark effect (QCSE) and filling of the states in the emission centers. 8,9 The lack of such observation in MQW-Si suggests the local energy minima are not as pronounced. Such local energy minima have been reported in the literature as a manifestation of indium clustering. To further explore the differences between the samples, surface analysis using atomic force microscopy (AFM) is conducted. Figure 3 shows 20 µm × 20 µm AFM scans of (a) MQW-Al 2 O 3 and (b) MQW-Si with identical height scales. The surface morphology of MQW-Al 2 O 3 shows a surface with step-flow terraces that have step heights of ∼ 0.3 nm, which corresponds to a monolayer of GaN. The threading dislocations in Fig. 3a has a density of ∼ 2.87 × 10 8 cm -2 . The root-mean-square (RMS) surface roughness is measured as 1.17 nm. The surface of MQW-Si shows a different surface morphology that is covered by hexagonal spiral hillocks with a mixed type dislocation in the center that are ∼ 5 µm apart and with a density of 6.56 × 10 6 cm -2 . 10 The treading dislocation of MQW-Si has a density of ∼ 3.20 × 10 8 cm -2 . The long-range order of these hillocks is reflected in the high RMS roughness of 2.81 nm. To further explore the differences between the crystals structural analysis using X-ray diffraction (XRD) and Raman spectroscopy are carried out. Figure 4 shows the (a) symmetric ω/2θ XRD spectra and (b) Raman spectroscopy of the MQW-Al 2 O 3 (blue, solid line) and on Si (red, dashed line). In Fig. 4a, the (004) GaN reflections are observed at θ ≈ 73 • with InGaN/GaN MQW fringe peaks on the left for both samples and AlGaN buffer layers' peaks on the right for MQW-Si. The 0 th peak position of InGaN implies an average indium composition of ∼ 2% across the MQW. The thickness of each well-barrier period is calculated from the fringe spacing using the relationship: 11

II. EXPERIMENTS AND DATA ANALYSIS
where n i and n j are the order of the i th and j th fringe peaks, respectively, and λ is the wavelength of the X-ray (Cu-K α 1 , 0.154056 nm). This yields a well-barrier pair thickness of 14.8 nm. Asymmetric scan on the (102) reflection is conducted to evaluate composition of the Al 1-x Ga x N buffer layers. The lattice constants, a and c, are calculated using the d-plane spacing (d hkl ), which is the interplanar spacing of the lattice plane (hkl) that satisfy Bragg's condition: and Poisson-Vegard's law: where c m (x) and a m (x) are the measured lattice constants, c 0 (x) and a 0 (x) are the relaxed lattice constants, ν(x) is the Poisson's ratio, δ a and δ c are the lattice bowing parameters. 12,13 The compositions of the Al 1-x Ga x N buffer layers are calculated in the order from the substrate to the surface to be 1.00, 0.83, 0.57, and 0.29. Additionally, the lattice constants and strain of the GaN layers are also calculated and tabulated in Table I The results are tabulated in Table II, which shows MQW-Si contains approximately four times more screw-and edgetype dislocations than MQW-Al 2 O 3 . Raman spectroscopy is conducted with a Horiba Confocal Raman Microscope with an 1800 l/mm grating and a λ = 532 nm laser using a Z(XX)Z configuration to calculate strain in the GaN layers of the InGaN MQW. As shown in Fig. 4b, MQW-Al 2 O 3 shows three distinct Raman shifts: 417 cm -1 , which is attributed to the A 1g phonon mode from the Al 2 O 3 substrate, 14 and 569.7 and 735.3 cm -1 , which are associated with the E H 2 and A 1 (LO) phonons of GaN. 15,16 MQW-Si shows four Raman shifts: a saturated peak at 520 cm -1 from the Si substrate, 565.7 and 732.1 cm -1 from the E H 2 and A 1 (LO) phonons of GaN, and a weak 646 cm -1 peak, which is likely a E H 2 phonon mode associated with a AlGaN buffer layer. 17 The E H 2 and A 1 (LO) Raman shifts of GaN in the MQW-Al 2 O 3 exhibits a blueshift in energy from its relaxed value (567.5 cm -1 , 734.0 cm -1 , respectively); this indicates a compressive strain experienced in the GaN layer. MQW-Si exhibits the opposite shifts, which indicates the presence of a tensile strain. The E H 2 peak is used to calculate the stress based on its deviation from the relaxed value due to its high strain sensitivity. 18 The in-plane biaxial stress is then calculated using: 19 where ∆ω is the Raman shift and σ xx is the stress. The in-plane strain is calculated using: 20 where C ij are the elastic constants of GaN (C 11 = 390 GPa, C 12 = 145 GPa, C 13 = 106 GPa, and C 33 = 398 GPa), 21 which give a proportionality factor of 478 GPa. The calculated inplane stress and strain values are tabulated in Table II. Raman spectroscopy shows a compressive strain (-0.107%) in MQW-Al 2 O 3 and tensile strain (+0.088%) in MQW-Si, in line with XRD results. Figure 5 shows the PL spectra of (a) MQW-Al 2 O 3 and (b) MQW-Si from 6 ∼ 300 K obtained using a 10-mW frequencytripled laser (λ = 266 nm) and a helium-bath cryostat, and the extracted temperature dependent (c) internal quantum efficiency (IQE) and (d) peak emission energy. MQW-Al 2 O 3 exhibits InGaN multiple quantum well peak (MQW) centered at 2.812 eV at 300 K. At low temperature (T < 100 K), a shoulder on the low-energy (2.728 eV) side of the MQW peak and a peak in the ultraviolet range (3.486 eV) appear. The low energy peak is attributed to the localized indium clusters, 22 which is caused by the phase separation as a result of lattice mismatch, 23 whereas the UV peak is attributed to near-bandedge (NBE) emission from excitons recombining in the GaN layer. 24 The PL spectra for MQW-Si show three major emission centers: the MQW at 2.853 eV, donor-acceptor-pair (DAP) transition of GaN at 3.288 eV, 24 and NBE emissions from GaN at 3.446 eV. These emissions are convolved with the Fabry-Perot interference caused by the thinner films on the silicon substrate. The DAP emission exhibits the commonly seen LOphonon replicas that are spaced at 90 meV apart. 25 The DAP recombination indicates the elevated impurity concentration in MQW-Si, while the stronger GaN NBE emission suggests the carrier confinement is poorer due to the lower crystal quality.
The difference of ∼41 meV in the main MQW emissions from the samples throughout the temperature range (Fig. 5d), despite the same physical structure shown by XRD, is attributed to the reduced Quantum Confined Stark Effect (QCSE) in the InGaN active layers. The compressive strain experienced by the GaN layer in the MQW-Al 2 O 3 leads to an even larger compressive strain in the InGaN wells. This leads to a stronger piezoelectric polarization field in the InGaN layer, and thus a reduction in the QW ground state energies. The tensile strain in the GaN layer of the MQW-Si leads to a reduced compressive strain in the InGaN layer, which in turn exhibits a reduced piezoelectric polarization. 26 The reduced QCSE is believed to increase the electron-hole wavefunction overlap, which should increase the efficiency of radiative recombination. Figure 5c shows the temperature-dependent IQE, which is calculated using the following: where I(T) is the integrated intensity of the PL emission at temperature T, and assuming at 6 K all the recombination mechanisms are radiative. The IQE of the LEDs are shown in Fig. 5c where I is the integrated intensity, I 0 and C are fitting constants, k is the Boltzmann's constant, T is temperature, and shows that the carriers in the local energy minima of MQW-Al 2 O 3 requires very little thermal energy to escape to other higher energy emission centers, while MQW-Si requires 2.6 times more energy to activate a carrier loss mechanism. The peak MQW emission energies from 6 to 300 K are plotted in Fig. 5d. The MQW emission from the MQW-Al 2 O 3 exhibits a distinct S-shape shift in peak energy with increasing temperature in addition to the overall bandgap shrinkage, which is described by the Varshni coefficient. 28 The peak energy remains stable between 6 and 100 K, but then rapidly blue-shifts by 10 meV at 100 K, and then red-shifts by 21 meV steadily from 170 to 300 K. The S-shaped peak shift has been reported in the literature as caused by the temperaturedependent carrier dynamics in the self-organized inhomogeneous InGaN local minima. 29 The coincidence of a blue shift in the peak energy and an increase in the IQE of the MQW of the MQW-Al 2 O 3 at 130 K suggests a change in the emission center as carriers gain enough thermal energy to freely move between the low energy minima and the higher energy minima as the temperature increases. 30 On the other hand, the MQW emission from MQW-Si monotonically red-shifts by 16 meV with increasing temperature. To verify the reduced QCSE and the higher density of states in MQW-Si, power dependent PL is conducted. Figure 6 shows the injection power dependent PL of (a) MQW-Al 2 O 3 and (b) MQW-Si with laser power varying from 0.1 to 10 mW at 6 K. With 0.1 mW of laser power (purple spectrum), MQW-Al 2 O 3 shows a main LED peak at 2.813 eV, a red shoulder at 2.755 eV, and a broad, undefined blue shoulder, whereas MQW-Si shows a main LED peak at 2.830. A powerlaw dependence of the PL peak intensities on the laser power is fitted to quantify the difference between the two LEDs: 31 where I is the intensity of the emission, P is the power of the incident laser, and m is the power dependence. 0.91 ± 0.05 (MQW-Si main peak). The sublinear power dependence observed in both emission centers for MQW-Al 2 O 3 indicates the band filling of the low energy minima. This happens when a high concentration of injected carriers occupies all the available states in the lowest emission center, causing some of the carriers to occupy a otherwise higher energy states and subsequently emit photons with higher energies. The high energy shoulder has been reported to be a combination of the filling of the band-tail states and the screening of the internal electric fields. 8,9 The power-law fitting of the MQW-Al 2 O 3 main peak is separated into two laser power regimes of 0.1 ∼ 0.5 mW and 0.5 ∼ 10 mW and the power dependences are m = 0.94 ± 0.09 and m = 0.65 ± 0.02, respectively. This shows that the main peak of MQW-Al 2 O 3 exhibits a band-to-band-like transition that at laser power below 0.5 mW, whereas at laser power above 0.5 mW it exhibits defect-like band-filling behavior with a featureless high energy shoulder. MQW-Si, on the other hand, shows an approximately linear (m ≈ 1) power dependence. The observed in MQW-Al 2 O 3 is not present in MQW-Si. This suggests the bands in the main emission centers of MQW-Si are not fully-filled, and therefore the carrier spilling to higher energy states is not as pronounced. To observe the carrier recombination mechanisms in the MQW time-resolved photoluminescence (TRPL) is conducted. Figure 7 shows the TRPL spectra and mono-exponential decay fit of TRPL spectra and mono-exponential decay fit of (solid blue) MQW-Al 2 O 3 and (dashed red) MQW-Si. MQW-Al 2 O 3 (blue) and MQW-Si (red) using a frequency-doubled Ti-Al 2 O 3 laser (λ = 380 nm) to directly create electron-holepairs in the quantum well, and an avalanche photodiode with a bandpass filter to observe the LED emission. Using single exponential decay fitting, the PL lifetime of MQW-Al 2 O 3 and MQW-Si are calculated to be 4.9 ± 0.3 and 7.3 ± 0.1 ns, respectively. Given the relationship of lifetimes with IQE 32 where τ rad and τ PL are the radiative and photoluminescence lifetimes, the radiative lifetimes are calculated. Nonradiative recombination lifetimes (τ nr ) are calculated via

ARTICLE scitation.org/journal/adv
for MQW-Si. The faster radiative lifetime in MQW-Al 2 O 3 is attributed to the higher carrier concentration in the smaller indium clusters, while the faster nonradiative lifetime is attributed the carrier escape from these local potential minima. On the other hand, the slower radiative lifetime in MQW-Si suggests the recombination centers are operating in a lower carrier concentration regime, which has a lower probability of carrier escape. The net result of a reduced carrier escape probability and a lower crystal quality is a ∼ 40% longer nonradiative lifetime.

III. DISCUSSION
It has been shown that the carrier dynamics are very different in MQW-Al 2 O 3 and MQW-Si with the latter exhibiting less structural and optical characteristics that are associated with indium clustering. These include: S-shaped temperaturedependent PL peak emission energy, defect-like injection intensity dependence of the PL peaks, and correlation between intensity and spectral non-uniformities. The suppressed indium clustering effects are attributed to the residual tensile strain (-0.088% ∼ -0.162 %) in the GaN epitaxial layer in MQW-Si. It has been reported in the literature that tensile strain is beneficial for indium atom incorporation during MOCVD growth in the InGaN MQW by strain compensation 33 and the increased indium incorporation. 34 The higher InGaN uniformity indicates the ability of MQW-Si to incorporate a higher indium mole fraction than MQW-Al 2 O 3 before phase separation becomes significant. This implies the possibility of achieving a higher indium content green InGaN LED using the MQW-Si platform. Additionally, the tensile strain in the GaN layer can help reduce the compressive in the In 0.15 Ga 0.85 N layers and thus the piezoelectric polarization field. The reduced overall polarization field in the InGaN layers is beneficial for high efficiency emitters via the greater electron-hole wavefunctions overlap. 35 Thus, for InGaN devices on Si it may be more beneficial to not eliminate all the tensile strain originated from the silicon substrate in order to compensate for the compressive strain in the active InGaN layers.

IV. CONCLUSION
In conclusion, MQW-Si is found to be superior to MQW-Al 2 O 3 in emission wavelength and intensity areal uniformity, higher carrier loss mechanism activation energy, and the suppression of band-filling effect, as the result of a more uniform indium incorporation. This is attributed to the residual tensile strain found in the GaN layer of MQW-Si that also reduces the QCSE in the InGaN layers. Despite a lower crystal quality, MQW-Si exhibits a higher IQE at temperatures below 250 K and a similar one at room temperature with an epitaxial thickness that is 30% thinner. Overall, the results presented herein demonstrates the advantages of growing blue InGaN/GaN MQW emitters on Si over Al 2 O 3 substrates.