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Abstract
High performance InGaN-based laser diodes (LDs) monolithically grown on Si is fundamentally interesting and highly desirable for photonics integration on Si platform. Suppression of point defects is of crucial importance to improve the device performance of InGaN-based LDs grown on Si. This work presents a detailed study on the impact of point defects, such as carbon (C) impurities and gallium vacancies (VGa), on the device characteristics of InGaN-based LDs grown on Si. By suppressing the VGa-related defect within the waveguide layers, reducing the thermal degradation of InGaN-based quantum wells, and controlling the C impurity concentrations within the thick p-type cladding layers, the as-fabricated InGaN-based LDs grown on Si exhibited a significantly reduced threshold current density of 2.25 kA/cm2 and an operation voltage of 4.7 V.
© 2019 Optical Society of America under the terms of the OSA Open Access Publishing Agreement
1. Introduction
The success of electrically pumped InGaN-based laser diode (LDs) directly grown on Si has opened up a new era of the GaN optoelectronics integration on Si platform [1–5]. By virtue of the carefully engineered AlN/AlGaN multilayer buffer, the major challenges of growing crack-free high-quality GaN on Si due to the huge mismatch in both lattice constant (~17%) and coefficient of thermal expansion (∼54%) between GaN and Si can be well resolved [6,7]. Upon the crack-free and high quality GaN film grown on Si, we have demonstrated the first room-temperature electrically injected InGaN-based LDs on Si [8–10].
However, the performance of the as-fabricated InGaN-based LDs grown on Si is to be improved, including the relatively high threshold current density (> 4 kA/cm2) and high threshold voltage (> 8 V) [8]. The high threshold voltage was mainly related to the non-optimized p-type doping profile with unintentionally incorporated carbon (C) impurities, which act as donors and compensate the Mg acceptors within the p-type cladding layers, leading to a low hole concentration and hence a large series resistance. The large series resistance increases the power consumption and produces large amounts of Joule heat within the LDs, causing a steep rise in junction temperature and a rapid drop of optical output power under the high injection [11]. Moreover, the substantial non-radiative recombination caused by the point defects also converts the injected carriers into Joule heat and further elevates the junction temperature. High junction temperature promotes the diffusion of point defects into the active region of the LDs [12,13], which in turn reduce the internal quantum efficiency (IQE) and further increase the threshold current. In addition to the detrimental influence on the operation voltage, point defects also adversely affect the threshold current of the InGaN-based LDs grown on Si by the band-tail optical absorption within the waveguide layers, resulting in a large internal optical loss during the photon feedback oscillation within the resonance cavity [14,15]. The high voltage and threshold current eventually lead to immature LDs with a very limited lifetime (< 1 minute) [8]. It has been reported that the lifetime of InGaN-based LDs could be improved by six orders of magnitude (from 1 s to 300 h), when the voltage was reduced from 8 to 4 V [16,17]. Therefore, it is of great significance to study how to suppress point defects for the improvement of InGaN-based LDs performance.
In this work, we study in detail the impact of point defects on the optical and electrical characteristics of InGaN-based LDs grown on Si, which has led to a substantial reduction in both threshold current density (2.25 kA/cm2) and voltage (4.7 V). And the optical output power of the LDs can be well maintained after the on-bar test for a few hundreds of hours.
2. Experiments
The InGaN-based LDs were grown on Si(111) substrates by metal-organic chemical vapor deposition (MOCVD). The schematic diagram of the LDs can be found in Fig. 1(a). The reference sample A in our previous report [8] consisted of an AlN/AlGaN multilayer buffer, a 2.7-µm-thick n-type GaN layer, 1.2-µm-thick n-type Al0.05Ga0.95N lower optical cladding layer (CL), 80-nm-thick n-type GaN lower waveguide (WG) layer, 3 pairs of In0.12Ga0.88N (2.7 nm) / In0.02Ga0.98N (12 nm) multiple quantum wells (MQWs), 60-nm-thick undoped GaN upper WG layer, 20-nm-thick p-type Al0.2Ga0.8N electron blocking layer, 600-nm-thick p-type Al0.11Ga0.89N/GaN superlattice (SL) upper optical CLs and finally 30-nm-thick Mg doped p-type GaN contact layer. The structure of sample B was the same as that of sample A, except that the WG layers were replaced by a low In-content (2%) InGaN/GaN short-period SLs. The p-type CL of sample A was grown at 950 °C with a high growth rate of 17.4 nm/min. To suppress the thermal degradation of MQWs and reduce the concentration of C impurities in the upper CLs, the p-type CLs of sample B was grown at a relatively low temperature of 920 °C and a slow growth rate of 8.7 nm/min. Detailed comparison of the growth conditions between the two samples is shown in Table 1. It is noted that the adoption of InGaN/GaN SL WG would cause little effect on the thermal conductivity of the LD structure, because there was only a trace amount (1% in average) of In in such SL WG (sample B), and the total thickness of the In0.01Ga0.99N layer was only 70 nm (40 and 30 nm for the lower and upper waveguide, respectively), which was negligibly small than that of the AlGaN cladding layers (1.8 μm in total) and the n-GaN contact layers (2.7 μm).
Fig. 1 (a) Schematic diagram of the InGaN-based LDs grown on Si. (b) The cross-sectional HAADF STEM image of the InGaN MQW active region.
Table 1. Comparison of the growth conditions between samples A and B
The as-grown epitaxial wafers were processed into edge-emitting LDs by etching 4-μm-wide ridge into the p-AlGaN cladding layer with a cavity length of 800 μm. Both the rear and front facets were cleaved and coated with TiO2/SiO2 dielectric multilayers to reduce the mirror loss and the threshold current. The detailed LD fabrication and facet coating process can be found in our previous report [8]. The interface and thickness of the InGaN active region were studied by high-angle annular dark filed (HAADF) scanning transmission electron microscopy (STEM). The concentration depth profile of C impurity was studied by secondary ion mass spectroscopy (SIMS) measurement. Excitation of time-resolved photoluminescence (TRPL) was provided by a picosecond pulsed diode laser (EPL-375) with an excitation wavelength of 375 nm, a pulse width of 59.6 ps and a repetition rate of 100 MHz. The excitation wavelength was chosen to guarantee that only the InGaN MQWs were excited. The micro-photoluminescence (micro-PL) images were recorded by a Nikon A1 confocal laser scanning fluorescence microscopy with a 405 nm laser, and taken over the full visible wavelength range, using a sharp-cutoff filter of 405 nm to remove the excitation light. The on-bar electrical luminescence (EL) tests were performed by a fiber optic spectrometer under a pulsed current injection at room temperature, with a pulsed width of 400 ns and a repetition frequency of 10 kHz to minimize the self-heating effect on the device performance for a comparison between samples A and B. The current-voltage (I-V) characteristics of the LDs were measured with Keithley 4200 semiconductor characterization system.
3. Results and discussion
The structural difference between samples A and B was the lower and upper WG layers. As mentioned above, the GaN WG layers of reference sample A were grown at a high temperature of 1050 °C, while sample B featured InGaN/GaN SL WG layers with an average In content of 1% grown at 790 °C. Figure 1(b) is the HAADF-STEM image of the active region of sample B, from which the InGaN/GaN SL WG structure can be clearly identified. The thickness of QWs and quantum barriers (QBs) was determined to be 2.7 and 11.6 nm, respectively, which was in line with the design. In addition, the sharp interfaces between QWs and QBs along with a uniform thickness can be well observed, indicating a good control of the active region growth, which ensured a reduced light scattering and absorption loss.
It is known that high temperature growth is beneficial to improve the quality of p-AlGaN CL, which, however, usually causes a severe thermal degradation of the underlying InGaN/GaN MQWs [18]. As shown in the micro-PL image in Fig. 2(a), the reference sample A with the p-CL grown at a high temperature of 950 °C showed an inhomogeneous fluorescent morphology with patterns of bright and dark, indicating a severe thermal degradation of the active region [19]. In contrast, sample B with the p-CL grown at a reduced temperature of 920 °C presented a remarkably improved micro-PL image with a homogeneous emission pattern, as shown in Fig. 2(b), indicating that the MQWs of sample B were well preserved from degradation. According to Li et al., the thermal degradation of InGaN MQWs during the p-type layer growth initiates at the QW upper interface and is probably caused by the formation of In-rich clusters there [20]. Such thermal degradation usually produces VGa-related voids and metallic In precipitates in the active region [20,21]. It has been reported that the VGa defects and related complexes can act as efficient channels for Shockley-Read-Hall (SRH) recombination [22,23]. Therefore, a fast TRPL decay caused by the SRH non-radiative recombination with a reduced IQE can be expected for sample A, giving rise to junction temperature and threshold current. However, it is noteworthy that the thermal degradation for sample A has not yet reached the extent of yielding voids and In clusters (otherwise sample A cannot lase), but presents noticeable effects on the luminescence property [as shown in Fig. 2(a)]. Therefore, no obvious In clusters can observed in the TEM images of both samples A and B.
Fig. 2 Micro-PL images of the InGaN-based LDs grown on Si. (a) sample A and (b) sample B. (c) Room-temperature TRPL at the PL peak emission wavelength for the two samples.
To confirm the influence of thermal degradation, room-temperature TRPL measurements were performed for samples A and B. As shown in Fig. 2(c), a strong improvement in the effective lifetime of sample B (6.5 ns) can be observed as compared to sample A (3.5 ns), predicting a slower non-radiative and/or radiative recombination process. The transient TRPL characteristics usually can be studied using a two-exponential decay model, and the lifetime at the initial and final stages can be defined as [24]
where τinitial and τfinal are the lifetime of the initial and final stages of the TRPL response, τnr and τr the non-radiative and radiative carrier lifetime, A and B the SRH and radiative recombination coefficients. In Fig. 2(c), the unique initial stage and the non-exponential behavior at the long delays can be ascribed to the weak excitation and saturation of slow dynamics of the non-radiative recombination centers, respectively, which have already been observed in InGaN/GaN MQWs [25–27]. Given the weak excitation (0.03 W/cm2) and low carrier density (ΔN ~1.3 × 1017 cm−3) for each pulse, according to Eq. (1), the dominant recombination channel can be shifted to the non-radiative recombination (A coefficient dominates) in the samples [28]. Therefore, the fast TRPL decay (3.5 ns) in sample A can be partially ascribed to defect-assisted non-radiative recombination in the active region due to the thermal degradation of the MQWs. Both the TRPL and micro-PL studies confirmed that the InGaN/GaN MQWs of sample B was well preserved during the p-type CL growth at a reduced temperature.The IQE improvement of the active region in sample B was also related to the utilization of InGaN/GaN short-period SLs, replacing the GaN WG layers in sample A. The high-temperature growth and Si-doping of the GaN WG layers favors the formation of Ga vacancies VGa [29,30]. Hence, the Si-doped n-GaN WG layer grown at a high temperature of 1050 °C in sample A is believed to have a relatively high concentration of VGa. Both the VGa and Si impurities within the n-GaN WG layers are detrimental to InGaN-based LDs, because of their strong band-tail absorption during the light feedback oscillation in the resonance cavity [12,31,32]. Such absorption will increase the internal optical loss and reduce the IQE, resulting in an increase in junction temperature and threshold current for the LDs. In contrast, the InGaN/GaN SL WG layers with an average In content of 1% in sample B were grown at a low temperature of 790 °C. And the growth temperature reduction can partially suppress the VGa formation in the WGs. Moreover, the In atoms can effectively fill the VGa defects in the InGaN WG layers [26,33,34]. As a result, it is reasonable to believe that the InGaN WG layers of sample B have a significantly reduced concentration of VGa than the GaN WG layers of sample A. The suppression of VGa defects in the WG layers is beneficial to reduce the internal optical absorption loss, and improve the IQE and the threshold current of the InGaN-based LDs grown on Si. It is challenging to perform the optical waveguide loss measurements by the on-bar test of the as-fabricated bare LDs with no device packaging, because the self-heating effect would affect the measurement of the gain spectra with fine details. The optical waveguide loss measurement will be carried out after proper packaging of the LDs with good heat dissipation and lead-out electrodes, which is currently in development and will be reported elsewhere.
The employment of metalorganic precursors during the MOCVD growth favors the incorporation of C impurities, especially at a reduced growth temperature (920 °C) and a relatively high growth rate (17.4 nm/min) of the p-AlGaN CL [18,35]. The C impurities prefer to occupy Ga sites as donors in p-AlGaN and compensate the Mg acceptors, leading to a high series resistance and operation voltage of the InGaN-based LDs [36]. To suppress the unintentional incorporation of C impurities, a decreased growth rate of 8.7 nm/min was adopted for the p-AlGaN CL growth for sample B. The growth rate reduction gives a longer residence time for the surface C atoms to form hydrocarbons with hydrogen atoms [37]. The optimized growth condition for the p-AlGaN CL led to a reduction of C concentration by one order of magnitude, as shown in Fig. 3. The reduction of C impurities is expected to cut the series resistance, enabling a lower operation voltage and a concomitant cooler junction of the LDs. Additionally, the decrease of C impurities also helps reduce the light absorption within the p-AlGaN CL, thereby contributing to the IQE improvement and the threshold current reduction of the InGaN-based LDs grown on Si. Further optimization of the growth temperature, growth pressure, and growth rate is currently underway to further drop the C incorporation for the p-AlGaN CL.
Fig. 3 SIMS depth profile of C impurity concentration of the p-AlGaN cladding layer for samples A and B.
The EL spectra of sample B below and above the threshold in Fig. 4(a) demonstrate a lasing at a wavelength of 418 nm with a full width at half maximum of 0.4 nm. The lasing wavelength presents a small difference of 4.0 nm in comparison with sample A (413.4 nm) which may be caused by the unintentional difference in the growth temperature of the MQWs, since the incorporation of indium in InGaN QW is very sensitive to the growth temperature. But the small difference in the lasing wavelength would not cause any significant influence on the device performance. The far-field patterns (FFPs) are shown as the insets. The FFP at 86 mA (1.2 times threshold) exhibited an elliptical pattern elongated along the growth direction due to the asymmetric optical confinement along the latitudinal and longitudinal directions [38]. The FFPs and the narrowed EL spectra provided direct evidence of the stimulated emission for sample B. The power-current-voltage (L-I-V) characteristics for samples A and B are shown in Fig. 4(b). The threshold voltage of sample B (4.7 V) decreased by 3.5 V in comparison with the reference sample A (8.2 V), demonstrating a substantial reduction of series resistance due to the effective suppression of C impurity incorporation. Moreover, sample B exhibited a low threshold current of 72 mA (2.25 kA/cm2), which was less than half of sample A (150 mA, 4.7 kA/cm2). The reduction of threshold current for sample B can be ascribed to the suppression of thermal degradation of the MQWs by optimizing the growth condition of p-AlGaN CL [Figs. 2(a) and 2(b)], the mitigated internal optical absorption by suppressing point defects within the WGs and p-CLs (Fig. 3), and the decreased junction temperature by reducing the operation voltage [Fig. 4(b)] and non-radiative recombination in the active region [Fig. 2(c)].
Fig. 4 (a) EL spectra of sample B above (1.2 times) and below (0.8 times) the threshold current. The insets are the corresponding FFPs. (b) On-bar L-I-V characteristics under pulsed injection for samples A and B.
The top-view optical microscopy image of the as-fabricated GaN-on-Si LDs is shown in Fig. 5(a). The yield of the optimized LDs (sample B) was estimated by measuring 120 devices on twelve bars of the as-fabricated LDs. The optical power measurements were performed under pulsed injection with a pulsed width of 400 ns and a repetition frequency of 10 kHz. As shown in Table 2, the lowest, average and median threshold currents for the measured LDs were 72, 179 and 180 mA, corresponding to a threshold current density of 2.25, 5.6 and 5.6 kA/cm2, respectively. The results above presented an encouraging progress of InGaN-based LDs grown on Si, compared with the reference sample A, which showed the lowest, average and median threshold current of 4.4, 11 and 10.6 kA/cm2, respectively. The statistical histogram of threshold current for the as-fabricated InGaN-based LDs of sample B is shown in Fig. 5(b). 97% (116 in 120) devices in total could achieve lasing, indicating a decent yield of the as-fabricated InGaN-based LDs grown on Si. In addition, the on-bar lifetime test of the bare LDs of sample B was performed under pulsed mode (pulsed width of 400 ns and a repetition frequency of 10 kHz) at room temperature, which have sustained an on-bar test of 620 hours without any significant decay of the optical output power (not shown here), which is much longer than that of sample A. Device package of the as-fabricated LDs is currently underway to enhance heat dissipation, which is expected to further lower the junction temperature and improve the device performance.
Fig. 5 (a) Optical microscopy image of several InGaN-based LDs on bar after facet cleavage. (b) Statistical histogram of the threshold current for the as-fabricated InGaN-based LDs of sample B. The bell-shaped curve shows the normal distribution of the threshold current.
Table 2. The statistic results of threshold current (Ith) and threshold current density (Jth) for the as-fabricated LDs from samples A and B.
4. Conclusion
In summary, the performance of InGaN-based LDs grown on Si has been remarkably improved by suppressing the formation of point defects. By optimizing the growth temperature and growth rate of the p-AlGaN upper CLs, the thermal degradation of the InGaN/GaN MQWs was effectively suppressed, and the C impurity concentration in the p-AlGaN CLs was reduced by one order of magnitude, which led to a remarkable drop of the threshold voltage from 8.2 to 4.7 V. Additionally, the introduction of In atoms effectively filling the VGa defects within the WG layers, together with the reduced C impurities within the thick p-type CLs, cut down the optical absorption loss and enhanced the IQE. The improved InGaN-based LDs grown on Si exhibited a lasing wavelength of 418 nm with a low threshold current density of 2.25 kA/cm2 and operation voltage of 4.7 V. This work paves the way towards cost-effective high-performance InGaN-based LDs on Si platform.
Funding
National Key R&D Program (2016YFB0400100, 2016YFB0400104); National Natural Science Foundation of China (61534007, 61604168, 61775230, 61804162, 61874131); Key Frontier Scientific Research Program of the Chinese Academy of Sciences (CAS, QYZDB-SSW-JSC014); The CAS Interdisciplinary Innovation Team; Key R&D Program of Jiangsu Province (BE2017079); Key R&D Program of Guangdong Province (2019B010130001); Natural Science Foundation of Jiangsu Province (BK20160401, BK20180253); Natural Science Foundation of Jiangxi Province (20181ACB20002, 20181BAB211022); Suzhou Science and Technology Program (SYG201725, SYG201846); China Postdoctoral Science Foundation (2018M632408); The Open Fund of the State Key Laboratory of Reliability and Intelligence of Electrical Equipment (EERIKF2018001).
Acknowledgments
We would like to thank Dr. Mutong Niu from Platform for Characterization & Test of SINANO for the TEM characterizations, and Dr. Tong Liu, Dr. Zengli Huang and Prof. An Dingsun from Nano-X for the FIB-TEM specimen preparation. Technical supports from Nano Fabrication Facility of SINANO are also gratefully acknowledged.
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