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Compared with conventionally grown thin InGaN wells, thick InGaN wells with digitally grown InN/GaN exhibit superior optical properties. The activation energy (48 meV) of thick InGaN wells (generated by digital InN/GaN growth from temperature-dependent integrated photoluminescence intensity) is larger than the activation energy (25 meV) of conventionally grown thin InGaN wells. Moreover, thick InGaN wells with digitally grown InN/GaN exhibit a smaller σ value (the degree of localization effects) of 19 meV than that of conventionally grown thin InGaN wells (23 meV). Compared with green light-emitting diodes (LEDs) with conventional thin InGaN wells, the improvement in 20-A/cm2 output power for LEDs containing thick InGaN wells with digitally grown InN/GaN is approximately 23%.
© 2014 Optical Society of America
1. Introduction
Considerable progress has been achieved in the design and fabrication of GaN-based light-emitting diodes (LEDs), particularly for LEDs using InGaN/GaN multiple quantum wells (MQWs) as active regions [1–3]. To enhance luminous efficiency, significant effort has been exerted to improve the material quality [4,5], light-extraction efficiency [6], and metal-semiconductor ohmic contacts [7] of blue LEDs. InGaN-based blue LEDs can achieve external quantum efficiency of more than 70% [8,9]. However, the quantum efficiency of InGaN-based LEDs is significantly lower in the green to yellow (500-nm to 580-nm) spectral range, which is known as the “green-yellow gap” [8,9]. Therefore, for applications involving white LEDs with direct mixing of blue, green, and red lights, a high number of InGaN-based green LEDs are required to match the high-efficiency InGaN-based blue LEDs. For InGaN wells, green GaN-based LEDs require a high-In mole fraction, which results in an enlargement in the lattice-mismatch-induced strain of InGaN wells. V-shaped defects are easily formed in high-In MQWs and are triggered by threading dislocations in the buffer layer. These defects are formed because of strain relaxation associated with stacking faults or In segregation, and such defects typically terminate on the sample surface with V-shaped defects [10–15]. These defects can degrade the efficiency of GaN-based green LEDs. Recently, several studies have considered improving the interface abruptness and optical properties of green LEDs using trimethylindium (TMIn) preflow prior to the growth of InGaN quantum wells [16–20]. TMIn preflow results in smooth InGaN well surfaces, decreases V-shaped defects in InGaN/GaN MQWs [18], and improves the emission efficiency of InGaN/GaN MQW green LEDs. Moreover, several studies have also reported enhancements in surface migration of adatoms during the growth of III-nitride alloys using digital techniques [21,22]. In the present study, we demonstrate efficiency-improved green LEDs containing thick InGaN wells with digitally grown InN/GaN. We discuss the optical characteristics of thick InGaN wells with digitally grown InN/GaN in the InGaN/GaN MQW region. In addition, we discuss the effect of digital InN/GaN growth in thick InGaN wells on the characteristics and fabrication process of InGaN/GaN MQW green LEDs.
2. Experiments
All samples were grown on a 2 inch (0001) patterned sapphire substrate (PSS) using a 19 × 2 inch Thomas Swan close-coupled showerhead metal-organic chemical vapor deposition system. The PSS was prepared using a periodic convex pattern on a (0001) sapphire, which was fabricated using inductively coupled plasma reactive ion etching. The (0001) sapphire was etched on the cone-shaped photoresist layer. The pattern’s diameter, spacing, and height on the PSS were 3.5, 2.0, and 1.3 µm, respectively. During LED epitaxy, TMIn, trimethylgallium (TMGa), trimethylaluminum, and NH3 were used as source materials of In, Ga, Al, and N, respectively. Bicyclopentadienyl magnesium and silane were used as the p-type and n-type doping sources, respectively. The reactor temperature was initially raised to 900°C to grow a 10-nm-thick in situ AlN nucleation layer on the PSS. The reactor temperature was then raised to 1050°C to grow a 2-µm-thick undoped GaN (u-GaN) epitaxy on the in situ AlN nucleation layer. A 2-µm-thick n-GaN was then grown following the growth of the u-GaN layer at 1050°C. We ramped down the reactor temperature for the standard 12-pair InGaN (3 nm)/GaN (16 nm) MQW epitaxy. The low and high growth temperatures for the InGaN wells and the GaN barriers of the 12-pair InGaN/GaN MQWswere adopted. The growth temperature of both the InGaN QW and the barrier spacer layer was approximately 730°C, whereas the difference in growth temperature between the InGaN well and the GaN barrier was approximately 150°C. Conventional thin InGaN wells for green InGaN/GaN MQWs were prepared by flowing TMIn, TMGa, and NH3 simultaneously into the reactor for 80 s. By contrast, thick InGaN wells for green InGaN/GaN MQWs were prepared using digital InN/GaN growth. The InN and GaN growth time and the loops for InN/GaN growth switching were set at 2 s and 40 times, respectively. Conventional green LEDs and green LEDs containing thick InGaN wells with digitally grown InN/GaN were denoted LED I and LED II, respectively. The source-switching sequences and growth temperature profile during the growth of InGaN/GaN MQWs with conventional InGaN wells (i.e., LED I) and with thick InGaN wells derived from digital InN/GaN growth (i.e., LED II) are shown in Fig. 1.Notably, the growth time for the GaN barrier was identical for both LED samples. After the growth of InGaN/GaN MQWs, a 20-nm-thick Mg-doped AlGaN layer and a 150-nm-thick Mg-doped p-GaN layer were grown for each LED. Using indium tin oxide as the transparent conductive layer, the processing of LED chips involves methods similar to the processes reported in previous studies [9]. All fabricated LEDs presented a dominant emission wavelength of 511 nm and an area of 580 µm × 1160 µm. The current-voltage characteristics of the experimental LEDs were measured using an HP 4156C semiconductor parameter analyzer and a Keithley 2400. LED output power was measured with a calibrated integrating sphere at room temperature.
Fig. 1 Schematic of the source-switching sequences of the MQW growth conditions for LEDs (a) I and (b) II. Axis scales are not in proportion.
3. Results and discussion
The measured and simulated X-ray θ−2θ diffraction (XRD) spectra of LEDs I and II are shown in Fig. 2. The XRD spectra of both LEDs I and II exhibit distinct satellite peaks, which indicate the abrupt interfaces between the InGaN wells and the GaN barrier layers of green InGaN/GaN MQWs. The XRD results indicate that the period thickness of green InGaN/GaN MQWs containing either InGaN wells with digitally grown InN/GaN or conventional wells is approximately 19 nm. The full-width half maxima (FWHM) of the −1-order satellite peaks of LEDs I and II are 93.5” and 89”, respectively. The broadening of the FWHM of the satellite peaks is attributed to the interface roughness or to fluctuations in the alloy composition [23]. Compared with LED I, LED II presents a −1-order satellite peak with a smaller FWHM. Therefore, the structural properties and the compositional abruptness of LED II were remarkably improved as a result of digital InN/GaN growth for InGaN wells. The In composition of the InGaN well is obtained by fitting the XRD data, and the In contents in LEDs I and II are 27.5% and 16.7%, respectively. Compared with conventional InGaN wells, digital InN/GaN growth reduces the In content in InGaN by nearly half. Figure 3 shows transmission electron microscopy (TEM) images of the InGaN/GaN MQWs of LEDs I and II. The period thickness values of the InGaN/GaN pair of LEDs I and II are 21.5 nm and 22.8 nm, respectively. However, the thickness of the InGaN wells in LED II (7.2 nm) is found to be more than two times the thickness in LED I (3.3 nm).
Fig. 2 θ-2θ scan X-ray diffraction spectra of LEDs (a) I and (b) II. The satellite peaks are labeled with numbers. The lower curve is the simulation resulting from fitting the XRD data.
Furthermore, the InGaN/GaN MQWs containing thick InGaN wells with digitally grown InN/GaN have clearer interfaces between InGaN wells and GaN barriers than conventionally grown InGaN/GaN MQWs, and this finding is consistent with the XRD results. The morphology of InGaN/GaN MQW samples with and without digitally grown InN/GaN in thick InGaN wells was studied by atomic force microscopy (AFM). The pit density from AFM images (not shown here) of the InGaN/GaN MQWs containing thick InGaN wells with digitally grown InN/GaN (1.8 × 108 cm−2) is less than the pit density of the InGaN/GaN MQWs with conventional InGaN wells (2.9 × 108 cm−2). Furthermore, InGaN/GaN MQWs containing thick InGaN wells with digitally grown InN/GaN have less surface roughness (10.6 nm) than InGaN/GaN MQWs with conventional InGaN wells (16.2 nm). Therefore, compared with conventional InGaN/GaN MQWs, InGaN/GaN MQWs containing thick InGaN wells with digitally grown InN/GaN show superior crystal quality.
A photoluminescence (PL) measurement was performed on both LED samples to study their optical characteristics using a 25-mW HeCd laser as the excitation source. The room temperature (RT) PL spectra of both LEDs are shown in Fig. 4.The PL emission peak wavelength of both LEDs is approximately 506 nm. The FWHM of the PL spectrum of LED II (20.2 nm) is less than the FWHM for LED I (22.0 nm). Similarly, the PL peak intensity of LED I is less than the PL peak intensity of LED II. The enhanced intensity and reduced PL FWHM of LED II can be attributed to the improved crystal quality of green InGaN/GaN MQWs containing thick InGaN wells with digitally grown InN/GaN. Although LED II presents a 16.7% In content in its InGaN wells, an InGaN well more than 7 nm thick can push the emission wavelength to 511 nm, which approaches the emission wavelength of LED I. The emission wavelength is sensitive to the In content, and the well thickness is also affected. We have simulated both a conventional well and a thick well using APSYS software. The result shows that the wavelengths of LED I and II are about 510 and 520 nm, although this result is not shown here. This simulation result is quite similar to our experiment’s findings.
Temperature-dependent PL measurements were performed on both samples to determine the temperature dependence of the peak energy and integrated intensity for the InGaN/GaN MQWs of LEDs I and II. The temperature dependence of the PL spectrum peak positions for LEDs I and II is shown in Fig. 5.The PL spectrum peak positions for both samples exhibit a blueshift within the temperature range of 75 K to 200 K and also present an S-shaped curve with respect to temperature. This finding does not agree with the semiconductor band gap behavior predicted by the Varshni [24] or the Bose-Einstein [25] formula. The temperature-induced blueshifts of LEDs I and II (shown in Fig. 5) are approximately 17 meV and 15 meV, respectively. However, for LEDs I and II, the curves of the PL spectrum peak position with respect to temperature in the high-temperature region can be fitted using the formula reported by Eliseev et al. [26], which is a combination of the Varshni formula and the band-tail model:
where T is the temperature in Kelvin. The first term describes the energy gap at zero temperature, and α and β are known as Varshni’s fitting parameters. The third term is obtained from the localization effect, in which σ indicates the degree of the localization effect. A large value of σ indicates a strong localization effect. Boltzmann’s constant is kB. From fitting the equation to the experimental data, the values of σ obtained for LEDs I and II are 23 meV and 19 meV, respectively. LED I exhibits a larger σ than does LED II, and this difference implies a wide energy-scale distribution for band potential profile fluctuations. In InGaN/GaN MQW heterostructures, In compositional inhomogeneity has been proposed as the origin of the localization effect [27,28]. The inhomogeneity size of nanocrystallites in InGaN/GaN heterosystems causes potential fluctuations capable of spatially localizing excitations. The larger σ value of LED I can be attributed to the high In content of conventionally grown thin InGaN wells, which can lead to more severe composition inhomogeneity and thickness variation than observed in LED II.Fig. 5 PL spectrum peak position and FWHM with respect to temperature for LEDs I and II. The solid lines are fitted to the experimental data points using Eq. (1).
Figure 5 also plots the temperature-dependent PL FWHM. At a low temperature, carriers are randomly located within the potential minima. As the temperature increases, carriers are thermally activated and redistributed into the lowest potential minima while the FWHM decreases. When temperature increases, the carriers populate higher-energy states as the FWHM increases. LED I has a larger FWHM than does LED II at low temperatures. The larger value of the FWHM at low temperatures indicates a more inhomogeneous distribution [29].
The Arrhenius plot of the integrated PL intensity with respect to the measurement temperatures for LEDs I and II is shown in Fig. 6.The evolution of InGaN/GaN MQWs’ peak energy with respect to temperature presents a characteristic S shape, which is attributed to the evolution from localized states to extended band-tail states. Consistently, two activation energies are expected. For the two-recombination channel, the data were fitted using the following equation:
where I(0) is the integrated intensity of the emission at 0 K. Activation energies EA1 and EA2 are carrier delocalization and defect-related thermal quenching in the low- and high-temperature ranges, respectively [30,31]. The EA1 and EA2 of conventional InGaN wells are 6 meV and 25 meV, respectively. The EA1 and EA2 of thick InGaN wells with digitally grown InN/GaN are 14 meV and 48 meV, respectively. As these results indicate, thick InGaN wells with digitally grown InN/GaN demonstrate a larger EA1 than that of conventional InGaN wells. The larger EA1 indicates a larger carrier delocalization energy, which allows for greater efficiency in electron-hole pair recombination. Additionally, the thick InGaN well with digitally grown InN/GaN has a larger EA2 activation energy than that of conventional InGaN wells. The EA2 activation energy in high-temperature regions should be related to the lateral diffusion of the carriers into the dislocations, where they recombine nonradiatively [32]. Compared with conventional InGaN wells, the larger EA1 and EA2 activation energies of thick InGaN wells with digitally grown InN/GaN indicate a larger barrier for nonradiative recombination. Moreover, this finding implies that the quality of LED II (which contains 7.2-nm-thick InGaN wells with digitally grown InN/GaN) is superior to the quality of LED I (which contains conventional 3.3-nm-thick thin InGaN wells). Moreover, the suppressed nonradiative recombination in LED II can lead to higher internal quantum efficiency (17.9%) than that of LED I (15.9%) at 300 K, as estimated from Fig. 6 [33].Fig. 6 Arrhenius plots of the integrated PL intensity at different temperatures for LEDs I and II. For clarity, the curves have been shifted vertically. The solid lines are fitted to the experimental data points using Eq. (2).
Thus far, we have observed the superior optical properties of thick InGaN wells with digitally grown InN/GaN compared with conventionally grown thin InGaN wells. We have also fabricated LED chips to ascertain the light output power of LEDs I and II. The 20-A/cm2 forward voltages (Vf) of LEDs I and II are 3.35 V and 3.34 V, respectively. The reverse leakage currents at −15 V of LEDs I and II are 53.7 µA and 4.1 µA, respectively. The lower −15 V reverse leakage current of LED II can be attributed to the improved crystal quality of green InGaN/GaN MQWs containing thick InGaN wells with digitally grown InN/GaN.
For all LEDs, the measured output power and external quantum efficiency (EQE) as a function of the injection current density is shown in Fig. 7.The light output powers and EQE with respect to varying injection current density of LED II were larger than the values for LED I. The light output powers of LEDs I and II driven at 120 mA (20 A/cm2) were 89.5 mW and 110.2 mW, respectively, which correspond to EQEs of 30.8% and 37.9%, respectively. Therefore, compared with LED I, the improvement in 20-A/cm2 output power for LED II is approximately 23%. LED II also shows a higher peak EQE (48.6%) than is shown by LED I (34.0%). Compared with LED I, thick InGaN wells with digitally grown InN/GaN significantly improved the low-current-density EQE of LED II. This improvement is attributed to the larger carrier delocalization energy and defect-related activation energy of LED II. However, the improved low-current-injection EQE of LED II resulted in efficiency droop (efficiency degradation from the peak of the EQE to the EQE at 60 A/cm2); LED II experienced a larger efficiency droop (40.9%) than did LED I (28.2%).
We also investigated the relationship between emission wavelength and injection current density (as shown in Fig. 8). LED II has a larger blueshift in its emission wavelength (the emission wavelength difference between the injection current of 1 and 60 A/cm2); for LED II, the blueshift is 15.5 nm, while the blueshift for LED I is 10.6 nm. With increasing injection current, the emission wavelength blueshift of GaN-based green LEDs could result from the quantum-confined Stark effect (QCSE) and the band-filling effect of the InGaN wells. Both thick InGaN wells with digitally grown InN/GaN and conventional InGaN wells show the localization effect. The band filling of the localized state in wells leads to the blueshift in the emission wavelength for both LEDs as injection current increases. However, the thick InGaN wells with digitally grown InN/GaN are about two times thicker than conventional InGaN wells. Therefore, the QCSE could be the dominant factor causing the heightened blueshift of the LEDs containing thick wells with digitally grown InN/GaN in comparison with conventional InGaN wells.
LED II notably demonstrates a larger 20-A/cm2 output power than that of LED I despite its thicker InGaN wells. The improved 20-A/cm2 light output power of LEDs containing thick InGaN wells with digitally grown InN/GaN should be attributed to the improved crystal quality and larger carrier delocalization energy of thick InGaN wells, which is achieved by digitally growing InN/GaN at relatively low growth temperatures. The results for the comparison between the light output powers of LEDs I and II are consistent with the results for the aforementioned optical properties.
4. Conclusion
In summary, we demonstrate the properties of LEDs containing thick InGaN wells with high crystal quality by digitally growing InN/GaN for InGaN wells. Compared with conventionally grown thin InGaN wells, thick InGaN wells with digitally grown InN/GaN exhibit superior optical properties. Obtained from the Arrhenius plot of the integrated PL intensity with respect to the measurement temperature, the activation energy for defect-related thermal quenching of thick InGaN wells with digitally grown InN/GaN (48 meV) is larger than the activation energy of conventionally grown thin InGaN wells (25 meV). Moreover, thick InGaN wells with digitally grown InN/GaN exhibit a lower σ value of 19 meV than is shown by conventionally grown thin InGaN wells (23 meV). Finally, compared with green LEDs with conventionally grown thin InGaN wells, the 20-A/cm2 output power of green InGaN/GaN MQW LEDs containing thick InGaN wells with digitally grown InN/GaN has been improved by approximately 23%.
Acknowledgments
The authors are grateful to the National Science Council of Taiwan for their financial support under Contract Nos. NSC101-2221-E-006-066-MY3 and 102-3113-P-009-007-CC2. This research was also made possible by the Advanced Optoelectronic Technology Center, National Cheng Kung University (as a project of the Ministry of Education of Taiwan), and by the financial support of the Bureau of Energy, Ministry of Economic Affairs of Taiwan, under Contract No. 102-E0603.
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