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Abstract
Highly doped n-Al0.6Ga0.4N can be used to form tunnel junctions (TJs) on deep ultraviolet (UVC) LEDs and markedly increase the light extraction efficiency (LEE) compared to the use of p-GaN/p-AlGaN. High quality Al0.6Ga0.4N was grown by NH3-assisted molecular beam epitaxy (NH3 MBE) on top of AlN on SiC substrate. The films were crack free under scanning electron microscope (SEM) for the thickness investigated (up to 1 µm). X-ray diffraction reciprocal space map scan was used to determine the Al composition and the result is in close agreement with atom probe tomography (APT) measurements. By varying the growth parameters including growth rate, and Si cell temperature, n-Al0.6Ga0.4N with an electron density of 4×1019 /cm3 and a resistivity of 3 mΩ·cm was achieved. SIMS measurement shows that a high Si doping level up to 2×1020 /cm3 can be realized using a Si cell temperature of 1450 °C and a growth rate of 210 nm/hr. Using a vanadium-based annealed contact, ohmic contact with a specific resistance of 10−6 Ω·cm2 was achieved as determined by circular transmission line measurement (CTLM). Finally, the n-type AlGaN regrowth was done on MOCVD grown UVC LEDs to form UVC TJ LED. The sample was processed into thin film flip chip (TFFC) configuration. The emission wavelength is around 278 nm and the excess voltage of processed UV LED is around 4.1 V.
© 2021 Optical Society of America under the terms of the OSA Open Access Publishing Agreement
1. Introduction
AlGaN alloys have a wide range of applications in optical and electronic devices due to their attractive physical properties such as tunable wide bandgap, strong piezoelectricity, and a high breakdown field. Recently, deep ultraviolet (UVC) LEDs based on AlGaN materials have attracted significant interest for its disinfection applications such as water purification [1–4] and sterilization of critical care patients rooms [5,6].
Realizing highly conductive n-AlGaN with controlled doping is essential to achieve high performance devices. UV LEDs, for example, can benefit from the use of n-AlGaN based tunnel junctions (TJs) [7]. This is because the efficiency of UV LEDs is primarily limited by the light extraction efficiency which is usually below 30% [8–17]. TJs can potentially double the LEE by enabling the use of highly reflective MgF2-based mirrors in thin-film flip-chip UVC LEDs [17,18–21] and thus eliminate the need to use absorptive p-GaN [22,23]. TJs also facilitate current spreading and hole injection in UV LEDs [7,16,17].
Si-doped AlGaN has several challenges in comparison to than Si:GaN for several reasons. First, the ionization energy of Si dopant in AlGaN is higher than in GaN and increases with increased AlN content [23,24–34]. Ionization energy of Si in AlxGa1-xN increases from ∼15 meV for x = 0 to 60-280 meV for x = 1, with a value for 25-30 meV for x = 0.65 [24,35]. Second, Si also shows self-compensation at high doping level for AlxGa1-xN with x > 0.5, leading to reduced electron concentration and increased resistivity as Si doping level increases above a certain value (also known as the “knee behavior”) [23,24,26–28,36–40]. At high Al composition, compensating occurs, either through DX center formation for the Si dopants [23,28,35,40–50] or through group III (i.e., Al or Ga) vacancy complexes [26–28,49,51–56,56–62]. Finally, unlike GaN, the electron mobility in AlGaN is impacted by alloy scattering.
Most of the studies on Si-doped AlGaN are done by metal organic chemical vapor deposition (MOCVD) [24–27,30,35–37,52,63–71], with some studies by plasma-assisted molecular beam epitaxy (MBE) [23,28,31,41–43], and few by NH3-assisted MBE (NH3 MBE) [29,38]. NH3 MBE has several advantages such as avoiding repassivation of p-type layers, the capacity to grow n-GaN with excellent transport properties [72], incorporating high levels of Si without compromising the material quality [73], sharp p+/n+ interface [74], less memory effect compared to MOCVD especially for Mg [75], and achieving high-performance TJ devices [74,76–79]. It is thus of our interest to investigate Si-doping of AlGaN by NH3 MBE.
In this study, we investigated n-type doping of AlxGa1-xN by Si with an Al fraction (x) around 60% by NH3 MBE. This Al composition was of interest to this study since negligible absorption of UVC wavelength occurs at this composition [23]. Besides, the maximum realizable carrier concentration can decreases with increased Al content x [23,24,26,28,30,36,42,43,52,63,65]. The relationship between n-AlGaN quality and growth parameters including growth rate and Si cell temperature was investigated. Using the optimized n-AlGaN condition, we demonstrate low n-AlGaN contact resistance and UVC TJ LEDs (around 278 nm).
2. Experimental
The templates used for this study were AlN grown by MOCVD on Si-face c-plane 6H-silicon carbide (SiC) substrates. The templates were crack-free, with a threading dislocation density (TDD) around 109 /cm2. Details on the low TDD AlN growth can be found elsewhere [80,81]. Before growth, the templates were coated with 500 nm of Ti on the back side by electron beam evaporation to assist the heat spreading and pyrometer temperature reading. Then, they were diced into 1 cm × 1 cm squares, solvent-cleaned, and indium-bonded to silicon wafers for transfer into the MBE. The samples were baked for 1 hour at 400 °C in vacuum before growth. The growth was done by Veeco 930 NH3-assisted MBE with NH3 as the group V source and solid effusion cells for Al, Ga, and Si. The flux for each group III elemental beam was performed using an ion gauge under NH3-free conditions. Temperature was measured using a pyrometer calibrated using the melting point of Al. The sample morphology was monitored in real time using reflection high-energy electron diffraction (RHEED). The growth temperature was set at 780 °C with a NH3 flow of 200 sccm. The group III fluxes were adjusted to realize a growth rate of ∼100-400 nm/hr and an Al content around 60%. The silicon doping level was controlled by growth parameters such as the Si cell temperature.
The composition and relaxation of the AlGaN films were evaluated on the PANalynatic Pixcel Diffractometer using reciprocal spacing mapping technique (RSM). The measurement was taken in an asymmetric scattering geometry with Cu K-α1 X-ray source, 2-bounce Ge (220) monochromator with a PIXcel3D detector. The lattice parameters of AlGaN was determined from the separation of its (1 0 $\bar{1}$ 5) peak and SiC (1 0 $\bar{1}$ 15) in RSM assuming the SiC is completely relaxed. Based on the Vegard’s law for the AlGaN lattice parameters and elastic constants, its alloy composition and in-plane strain can be therefore calculated. The thickness of the AlGaN films was determined by the spacing of fringes of the on-axis ω-2θ (0002) scan. Optical microscopy, atomic force microscopy (AFM), and scanning electron microscope (SEM) were used to characterize the morphology of samples. Atom probe tomography (APT) with a Cameca local electrode atom probe (LEAP) 3000X HR was used to evaluate the three-dimensional distribution of Al and Ga atoms and the alloy distribution [82].
The n-type doping optimization was carried out by growing two series of n-AlGaN samples with various Si cell temperature. Series 1 samples were grown at 109 nm/hr while Series 2 samples were grown at 210 nm/hr. The n-type doping of the AlGaN samples was evaluated by room temperature Hall measurements in a Van der Pauw geometry. The AlN on SiC templates underneath the n-AlGaN layer were confirmed to be insulating. To rule out the effect of interface charge, we also confirmed that a sample with undoped AlGaN grown on the AlN/SiC template is highly insulating. The Si doping level, as well as the C, H, and O impurity levels were measured by Cameca IMS 7f Auto secondary ion mass microscopy (SIMS). The morphology of n-AlGaN samples were evaluated by atomic force microscopy (AFM). Contacts to the n-AlGaN were formed using a V/Al/V/Au (10/150/20/300 nm, respectively) metal stack. The samples were annealed in N2 at 720 °C for 30 s to achieve an Ohmic contact. Circular transmission line measurement (CTLM) pattern was formed on the n-AlGaN with the annealed contact to measure the contact resistance.
Finally, the optimized n-AlGaN condition was used to form a UVC hybrid tunnel junction (TJ) LED. Mesa-defined MOCVD-grown UVC LED with 0.5 nm p-GaN on top was patterned with SiO2 to allow selective area growth of MBE-grown n-Al0.59Ga0.41N on top of both the 0.5 nm MOCVD-grown p-GaN and n-AlGaN region. The Mg concentration in the 0.5 nm p-GaN and underlying 50 nm of AlGaN:Mg layers was ∼1×1020 cm−3 and ∼2-3×1018 cm−3, respectively. Only 0.5 nm of p-GaN was used to minimize the absorption by GaN. The lower 4 nm n++ layer forms the n-side of the TJ and was grown with high Si cell temperature (1450 °C) to achieve high donor concentration ND. The 500 nm n+ layer was grown with the optimized condition with Si cell temperature at 1425 °C for the highest n-type carrier concentration and the lowest resistivity to facilitate current spreading while maintaining low resistivity. The upper 4 nm n++ layer serves as the contact layer and was grown with the same condition as the lower n++ layer. The growth rate used for the regrowth was 426 nm/hr. The sample was then processed into thin film flip-chip (TFFC) LEDs using the vanadium-based contact as both n and p-side contact. Details of the MOCVD epitaxial structure and the TFFC UV LED fabrication techniques can be found elsewhere [18,82,83]. The electroluminescence (EL) spectrum of the UV TJ LED was acquired after regrowth and before further processing. The on-wafer testing was done using a fiber optic and UV-Vis spectrometer (Ocean Optics) at 10 mA with an integration time of 50 ms. The fiber was placed above the LED sample, in proximity to the device. The IV characteristics of the processed UV TJ LED was measured and compared with the reference UV LEDs without the TJ regrowth and with 5nm p-GaN.
3. Results
3.1 Epitaxial growth of AlGaN
The epitaxial growth of undoped AlGaN was carried out to map out the growth condition to achieve the desired crystal quality and alloy composition (∼60%). Figure 1(a) shows the SEM image of a sample with a 110 nm thick AlxGa1-xN layer, x = 0.55, and a growth rate of 218 nm/hr. The sample was crack free and no pits were observed. No cracks were observed for the grown undoped AlGaN and n-AlGaN up to a thickness of 1 µm (the highest thickness investigated) as shown in the optical microscope (OM) image in Fig. 1(b). Figure 1(c) shows the AFM image of MOCVD-grown AlN on SiC template. The surface morphology was smooth with a root mean square (RMS) roughness of 0.3 nm. Figure 1(d) shows the AFM image of 100 nm AlxGa1-xN layer grown by MBE on the template with x=0.65 and a growth rate of 216 nm/hr. Very few (∼107 /cm2) pits were observed on the sample surface. The granular features in the AFM image of AlGaN are commonly observed in NH3 MBE films [72]. The RMS roughness was 0.48 nm. The comparison of the sample’s AFM images before and after the AlGaN growth suggests that the MBE growth did not degrade the surface roughness.
Fig. 1. (a). SEM image of a crack-free undoped AlGaN layer (110 nm), with Al=55.5%, and a growth rate of 218 nm/hr. (b). Optical microscope image of a crack-free Al0.59Ga0.61N layer (1 µm) grown at 210 nm/hr with both doped and undoped sub-layers showing no cracks. (c). AFM image of the AlN on SiC template, which shows step flow growth. The RMS roughness is 0.3 nm (d). AFM image of 100 nm undoped AlGaN layer grown by MBE on the template with Al%=65.5% and a growth rate of 216 nm/hr. The surface was smooth and the RMS roughness was 0.48 nm.
Figure 2(a) shows the ω-2θ scan of 100 nm Al0.60Ga0.40N film grown at 218 nm/hr. The clear thickness fringes indicate high crystal quality and the spacing of the fringes was used to calculate the film thickness and growth rate. Figure 2(b) shows an RSM scan of 400 nm Al0.60Ga0.40N film grown with the same condition. The AlN layer is fully relaxed [80] while the AlGaN film for this sample is coherent with the AlN layer as determined by the identical Qx values for the off axis (10$\bar{1}$5) AlGaN (bottom) and AlN (top) peaks. The separation between the SiC and the AlGaN peak was used to calculate the Al% composition.
Fig. 2. (a). ω-2θ scan of 100 nm Al0.60Ga0.40N film grown at 218 nm/hr with clear thickness fringes in log scale. The spacings between the fringes were used to calculate the film thickness and growth rate. (b). RSM scan of 400 nm fully strained Al0.60Ga0.40N film grown at the same condition as in 2 (a). The separation between the SiC (in the middle) and the AlGaN peak (bottom) was used to calculate the Al% composition and relaxation ratio. The same Qx between the AlN (top) and AlGaN peak shows that the AlGaN film is fully strained.
At fixed group III fluxes ΦIII (ΦAl: 1.35×10−7 torr; ΦGa: 1.19×10−7 torr) and an NH3 flow rate of 200 sccm, the growth temperature was varied in the range of 750-800 °C. Figure 3 shows that the Al% composition and growth rate were rather stable in this temperature range. The Al% changed within 3.5% and the growth rate varied within a 10% range.
The relationship between alloy composition and group III fluxes were also investigated. At a NH3 flow rate of 200 sccm and a growth temperature of 800 °C, the Al flux ΦAl was varied while keeping the Ga flux fixed at 6.08×10−8 torr. The measured Al% values are shown in Table 1.
Table 1. The relationship between group III fluxes and Al%. $S = \frac{{{S_{Al}}}}{{{S_{Ga}}}}$ is the ratio of incorporation factor between the two group III elements.
When the other growth parameters were kept the same, the Al content was determined by the group III fluxes. An incorporation factor (often referred to as a sticking coefficient) s can be defined as:
where N is the flux of atoms incorporated into the AlGaN layer, and $\mathrm{\Phi }$ was the supplied flux for this element. The Al% composition is thus determined by:The value of S was quantified for a given growth temperature and used to calculate the Al content for different Al and Ga fluxes when the change in group III flux were small. The calculated Al content value agreed well with the measured value for all samples investigated in this work. Typical range of S is 1.15-2.20, suggesting that Al atoms more readily incorporate into the AlGaN film than the Ga atoms which is consistent with the higher Al-N bond energy (2.2 eV) compared to the Ga-N bond (1.93 eV) [84].
APT measurements of the AlGaN layer are shown in Fig. 4. The Al% composition measured by APT is 59% and is in close agreement with the 58% Al composition calculated by XRD. The Al content was constant in the growth direction for the thickness studied (60 nm). The alloy was found to be random, and no Al rich clusters were observed [82].
Fig. 4. 3D APT reconstruction of 60 nm of Al0.59Ga0.41N showing Al and Ga atoms. (b) 2D distribution of the Aluminum fraction measured from the dashed rectangle shown in (a).
3.2 Doping optimization for Si-doped AlGaN
The doping optimization of the n-AlGaN was done by varying the growth rate and Si cell temperature while keeping the Al content and other growth parameters fixed. Figure 5(a) shows the n-type carrier concentration, resistivity, and mobility of n-AlGaN (Series 1) with an Al content of 61.6% and a growth rate of 109 nm/hr. The Si cell temperature was varied from 1375 to 1450 °C with 25 °C steps. The Si-doped AlGaN grown with a Si cell temperature of 1450 °C was insulating and no valid Hall data was acquired. Using a Si cell temperature of 1400 °C, both the lowest resistivity (10.5 mΩ·cm) and the highest n-type carrier concentration (1.2×1019 /cm3) was achieved. Similar doping series (Series 2) was done at a higher growth rate of 210 nm/hr for AlGaN with the Al content fixed at 59.4% as shown in Fig. 5(b). At this growth rate, the optimum Si cell temperature was at 1425 °C. The carrier concentration was 4×1019 /cm3, the mobility was 48 cm2·V·s−1, and the resistivity was 3 mΩ·cm. The morphology for this sample was measured by AFM and shown in Fig. 6. No pits were observed on the surface. The RMS roughness was 0.6 nm.
Fig. 5. Doping optimization for n-AlxGa1-xN by varying the Si cell temperature using a growth rate of (a) 109 nm/hr at x = 0.62 (series 1) and (b) 210 nm/hr at x 0.59 (series 2). The n-AlGaN grown at 109 nm/hr with the Si cell at 1450 °C was resistive and thus not shown in (a). The star shape denotes the optimum growth condition for the 210 nm/hr growth rate.
Fig. 6. AFM image of the optimum n-AlGaN growth condition (star sign in Fig. 5). The RMS roughness is 0.6 nm. There were no pits on the surface.
Increasing the AlGaN:Si growth rate resulted in a wider optimal doping window, higher carrier concentration, and lower resistivities. Both doping profiles showed the “knee behavior” commonly observed for Si:AlGaN [23,24,26–28,36–40]. As the Si cell temperature increased, the resistivity decreased while the carrier concentration increased at the low doping region and the opposite happened at the high doping region. The optimum growth condition was at the Si cell temperature of 1400 °C and 1425 °C for Series 1 and 2, respectively (refer to Fig. 5).
The Si, C, H, and O concentration (Fig. 7) were measured by SIMS for Series 2. The Si concentration increases with increased Si cell temperature. At a Si cell temperature of 1425 and 1450 °C, Si concentration of 1×1020 and 2×1020 /cm3 were achieved. The oxygen and carbon level were at 4-7×1017 and 2×1016 /cm3, respectively, and were stable at all Si cell temperatures. When the Si cell temperature increased from 1400 °C to 1425 °C and then 1450 °C, the hydrogen incorporation increased from 1.5×1017 /cm3 to 7.5×1017 /cm3 and then to 1.8×1018 /cm3. Suggesting that there could be increased H or H-related compensation at high doping levels.
Fig. 7. Si, C, H, and O concentration in n-AlGaN series 2 samples measured by secondary ion mass microscopy (SIMS). The n-type carrier concentration is included for comparison. The star shape denotes the optimum doping condition.
CTLM measurements were done on a V/Al/V/Au contact deposited on a Si-doped Al0.59Ga0.41N. The growth of the AlGaN layers was done at the same time as the TJ regrowth and have the same epi structure as the MBE part in Fig. 9(b). The carrier concentration and resistivity for the n+ layer was 1.9×1019 /cm3 and 12 mΩ·cm respectively. Figure 8 shows the IV curve for the 10 nm gap CTLM pattern. The contact changed from Schottky to Ohmic after annealing in N2 gas at 720 °C for 30 s. The resistance of the ohmic contact was 7.9 Ω as measured by 2-point probe. CTLM measured on the sample showed that the specific contact resistance was 5.3×10−6-8.9×10−6 Ω·cm2. The transfer length was 2.3 µm. The sheet resistance was around 125 Ω/□. The contact resistance is low enough that it does not introduce any appreciable voltage penalty in the IV characteristics of the different LEDs in the later part of the study.
Fig. 8. IV curve for CTLM pattern with 10 nm gap on n-AlGaN with vanadium-based contact before and after annealing in N2 gas at 720 °C for 30 s.
Fig. 9. (a). Schematic of the UV TJ LED. The upper image shows the top view of the LED. Both the n and p side epitaxial structure were terminated with MBE grown n-AlGaN. SiO2 was used for both the mask for selective area growth and isolation layer for the LED. The lower image shows the cross section of the LED. Selective area MBE n-AlGaN regrowth was done on the on top of the n-AlGaN and p-GaN region of the MOCVD grown UVC LED sample to form hybrid TJ UVC LED. The SiO2 layer served as the mask for the selective area growth. (b). Schematic of the epitaxial structure for the p-side of the UVC tunnel junction (TJ) LED.
3.3 UVC TJ LEDs
Schematic of the top and cross-section view of the UVC TJ LED after regrowth is shown in Fig. 9(a). Figure 9(b) shows the schematic of the TJ UV LED epitaxial structure.
Figure 10(a) shows the EL spectrum of the TJ LED wafer before LED processing, where the EL was obtained from the UV light emitted below indium dots and its surrounding perimeter (100 µm in diameter). The EL emission wavelength (around 278 nm) confirmed hole injection through the TJ structure into the LED active region. The EL intensity was low due to the low extraction efficiency and substrate absorption.
Fig. 10. (a) Electroluminescence spectra of the UVC LED with TJ (after regrowth and before further processing). Note the low intensity is likely due to the absorption by SiC. (b) I-V characteristics for the processed TJ LED and an identical LED without regrowth but with 5 nm of p-GaN on top of the MOCVD structure.
Figure 10(b) shows the IV characteristics of the processed TJ LED as compared to an identical LED without regrowth but with 5 nm of p-GaN on top of the MOCVD structure, which was reported elsewhere [81]. The TJ LED turned on, and had an excess voltage of 4.1 V at 20 A/cm2 compared to the LED with 5 nm of p-GaN. The excess voltage is high compared to blue TJ LEDs grown with similar approach (<1 V) [74,85,86]. This is partially due to the wider band gap of AlGaN, the higher activation energy of carriers, and the lower doping concentration (NA and ND) in AlGaN material. More optimizations are needed to reduce the excess voltage.
4. Conclusion
In this work, we demonstrated high carrier concentration, low resistivity n-AlGaN by NH3 MBE with the Al% around 60%. The films had smooth morphology and were crack free for the thickness investigated (up to 1 µm). XRD RSM scan was used to determine the alloy composition and the result agrees well with the APT result. The higher growth rate had a wider optimum doping window when varying the Si cell temperature. At a growth rate of 210 nm/hr, with a Si cell temperature of 1425 °C, n-AlGaN with a carrier concentration of 4×1019 /cm3 and a resistivity of 3 mΩ·cm was achieved. SIMS measurements shows that the Si incorporation increases with increased Si cell temperature and the hydrogen incorporation increased at high cell temperature. Ohmic contacts were achieved on the n-AlGaN using annealed vanadium-based contact stack. The optimized n-AlGaN was employed to show that a transparent tunnel junction can significantly reduce the thickness of p-GaN in UVC LEDs, however, further research is needed to reduce the excess voltage (4.1 V).
Funding
National Science Foundation (DMR 1720256, ECS-03357650); KACST-KAUST-UCSB Solid State Lighting Program SSLP; Solid State Lighting and Energy Electronics Center, University of California Santa Barbara.
Acknowledgements
The authors would like to thank Dr. Thomas Mates for help performing SIMS measurements. The authors would also like to thank Dr. Youli Li for insightful conversations about XRD. A portion of this work was carried out in the UCSB nanofabrication facility, with support from the NSF NNIN network, as well as the UCSB Materials Research Laboratory (MRL), which is supported by the MRSEC Program of the NSF under Award No. DMR 1720256; a member of the NSF-funded Materials Research Facilities Network (www.mrfn.org).
Disclosures
The authors declare no conflicts of interest.
Data availability
Data underlying the results presented in this paper are not publicly available at this time but may be obtained from the authors upon reasonable request.
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