N-polar iii-nitride semiconductors on silicon carbide

Abstract

A method of fabricating a heterostructure includes providing a substrate, the substrate including silicon carbide, the substrate having a carbon (C)-face, forming a III-nitride-based buffer layer of the heterostructure, the III-nitride-based buffer layer being in contact with the C-face of the substrate, and implementing a molecular beam epitaxy (MBE) procedure to grow a III-nitride-based layer such that the III-nitride-based layer is supported by the III-nitride-based buffer layer. The molecular beam epitaxy (MBE) procedure is implemented under nitrogen (N)-rich conditions

Description

N-POLAR lll-NITRIDE SEMICONDUCTORS ON SILICON CARBIDECROSS-REFERENCE TO RELATED APPLICATION

[0001] This application claims the benefit of U.S. provisional application entitled “N-Polar Ill-Nitride Semiconductors on Silicon Carbide,” filed October 26, 2023, and assigned Serial No. 63 / 593,393, the entire disclosure of which is hereby expressly incorporated by reference.STATEMENT REGARDING FEDERALLY SPONSORED RESEARCH OR DEVELOPMENT

[0002] This invention was made with government support under Contract No. 2026484 awarded by the National Science Foundation, and under Contracts Nos. W911 NF-22-2-0176 and W911 NF-23-1-0142 awarded by the Army Research Office. The government has certain rights in the invention.BACKGROUND OF THE DISCLOSUREField of the Disclosure

[0003] The disclosure relates generally to Ill-nitride heterostructures.Brief Description of Related Technology

[0004] In recent years, Ill-nitride semiconductors have deeply penetrated the realms of photonic and electronic devices, catalyzing disruptive technological advancements in high- power and high-frequency electronics, optoelectronics, and clean energy. Within this expansive domain of applications, two notable areas have garnered significant interest. First, the utilization of Ill-nitride semiconductors in deep ultraviolet (UV) optoelectronics, including light-emitting diodes (LEDs) and lasers, has emerged as a compelling solution for sterilization and water / air purification. Secondly, there has been significant progress in developing next-generation ultra-wide bandgap high-electron mobility transistor (HEMT) devices that possess high breakdown fields and efficient thermal management.

[0005] The successful implementation of these promising technologies relies on the utilization of tunable aluminum content of AIGaN layers with a precise control over dopinglevels spanning multiple orders of magnitude. The upper range of doping enables the realization of high electrical conductivity in injection layers for deep UV LEDs and lasers, whereas the lower range facilitates the creation of low-doped drift regions in HEMT devices.

[0006] Wurtzite Ill-nitride semiconductors, including GaN, InN, AIN, and their alloys, can crystalize along the c-axis with either metal or nitrogen polarity. While the majority of studies have focused on metal-polar Ill-nitride semiconductors, there has been a recent surge in interest regarding the distinct advantages offered by their nitrogen (N) polar counterparts. In the context of HEMT devices, an N-polar GaN / AIGaN heterostructure naturally forms a back barrier, contributing to enhanced carrier confinement, reduced contact resistance, and improved gate control compared to metal-polar devices. In connection with optoelectronic devices like LEDs and laser diodes, an N-polar configuration offers notable advantages such as reduced electron overflow and enhanced electrical efficiency, surpassing the performance of their metal-polar counterparts.

[0007] Unfortunately, to date, the realization of high-quality N-polar AI(Ga)N has remained extremely challenging. For instance, N-polar AI(Ga)N grown on substrates like sapphire and silicon are plagued with the presence of high densities of dislocations, due to the large lattice mismatch and the formation of inversion domains and poor surface morphology. Furthermore, while the growth and characteristics of N-polar AIN on a native AIN substrate or sputtered AIN template have been reported, the direct epitaxy of N-polar AI(Ga)N on an AIN bulk substrate encounters severe limitations primarily due to the prohibitively high cost associated with the substrate, making this approach unscalable.

[0008] Prior investigations utilizing MOCVD techniques have demonstrated the successful attainment of N-polar Ill-Nitrides on 4H-silicon carbide (SiC) substrates with a C-face orientation. In contrast to MOCVD, molecular beam epitaxy (MBE) growth presents significant advantages, including significantly reduced impurity incorporation, precise control over doping, and the flexibility to tune the composition especially in Al-rich AIGaN. Previous reports have demonstrated the epitaxy of metal-polar AIGaN on the Si-face of 4H-SiC substrates by MBE and MOCVD, and reports of p-type conduction in AIN have utilized MBE. However, growth of N-polar Ill-nitride materials is more involved than its metal-polar counterpart, primarily due to the small adatom diffusion length owing to a high binding energy on the N-polar surface. This promotes the formation of islands and other surface irregularities.SUMMARY OF THE DISCLOSURE

[0009] In accordance with one aspect of the disclosure, a method of fabricating a heterostructure includes providing a substrate, the substrate including silicon carbide, the substrate having a carbon (C)-face, forming a Ill-nitride-based buffer layer of the heterostructure, the Ill-nitride-based buffer layer being in contact with the C-face of the substrate, and implementing a molecular beam epitaxy (MBE) procedure to grow a Ill-nitride- based layer of the heterostructure such that the Ill-nitride-based layer is supported by the Ill- nitride-based buffer layer. The molecular beam epitaxy (MBE) procedure is implemented under nitrogen (N)-rich conditions.

[0010] In accordance with another aspect of the disclosure, a method of fabricating a heterostructure includes providing a substrate, the substrate including silicon carbide, the substrate having a carbon (C)-face, growing a Ill-nitride-based template layer of the heterostructure, the Ill-nitride-based template layer being in contact with the C-face of the substrate, and implementing a molecular beam epitaxy (MBE) procedure to grow a Ill-nitride- based layer of the heterostructure such that the Ill-nitride-based layer is in contact with the Ill-nitride-based template layer. Providing the substrate includes annealing the substrate in a growth chamber in which the MBE procedure is implemented. Growing the template layer is implemented in the growth chamber in which the Ill-nitride-based layer is grown. The molecular beam epitaxy (MBE) procedure to grow the Ill-nitride-based layer is implemented under nitrogen (N)-rich conditions.

[0011] In accordance with yet another aspect of the disclosure, a method of fabricating a heterostructure includes providing a substrate, the substrate comprising silicon carbide, the substrate having a carbon (C)-face, forming a Ill-nitride-based buffer layer of the heterostructure, the Ill-nitride-based buffer layer being in contact with the C-face of the substrate, and implementing a molecular beam epitaxy (MBE) procedure to grow a Ill-nitride- based layer of the heterostructure such that the Ill-nitride-based layer is supported by the Ill- nitride-based buffer layer. The substrate is not miscut such that growth of the Ill-nitride- based buffer layer and the Ill-nitride-based layer is on-axis.

[0012] In connection with any one of the aforementioned aspects, the methods described herein may alternatively or additionally include or involve any combination of one or more of the following aspects or features. The molecular beam epitaxy (MBE) procedure implemented to grow the Ill-nitride-based layer has a lll / V ratio of about 2:1 or higher. The molecular beam epitaxy (MBE) procedure implemented to grow the Ill-nitride-based layer has a lll / V ratio of about 3:1 . Providing the substrate includes cleaning the substrate via asolvent cleaning process. Providing the substrate includes annealing the substrate in a growth chamber in which the MBE procedure is implemented at a temperature sufficient to remove oxide from the C-face of the substrate. The MBE procedure to grow the Ill-nitride- based layer is implemented at a growth temperature that falls in a range from about 600 degrees Celsius to about 800 degrees Celsius. The Ill-nitride-based layer includes AIGaN. The MBE procedure to grow the Ill-nitride-based layer is implemented at a growth temperature that falls in a range from about 800 degrees Celsius to about 1100 degrees Celsius. The Ill-nitride-based layer includes AIN. Forming the buffer layer includes growing the buffer layer via molecular beam epitaxy in a growth chamber in which the Ill-nitride- based layer is grown. The Ill-nitride-based buffer layer includes AIN. The MBE procedure is configured to dope the Ill-nitride-based layer with silicon. The Ill-nitride-based layer includes AIN. The Ill-nitride-based layer includes AIGaN. The substrate has a 4H-SiC structure. The MBE procedure is configured such that the Ill-nitride-based layer is a ternary compound having an aluminum composition higher than another Group III element composition of the ternary compound. The substrate is a miscut SiC substrate. The Ill-nitride-based layer is in contact with the Ill-nitride-based buffer layer. The Ill-nitride-based template layer includes AIN. The MBE procedure is implemented at a growth temperature and with a 11 l / V ratio such that the Ill-nitride-based layer is atomically smooth. The 11 l / V ratio is about 2:1 or higher. The growth temperature falls in a range from about 600 degrees Celsius to about 800 degrees Celsius. The growth temperature falls in a range from about 800 degrees Celsius to about 1100 degrees Celsius. The MBE procedure is configured such that the Ill-nitride layer is N-polar.BRIEF DESCRIPTION OF THE DRAWING FIGURES

[0013] For a more complete understanding of the disclosure, reference should be made to the following detailed description and accompanying drawing figures, in which like reference numerals identify like elements in the figures.

[0014] Figure 1 depicts (a) a schematic illustration of a heterostructure having an N-polar Ill-nitride-based layer supported by an SiC substrate (e.g., an AIGaN / AIN / SiC heterostructure) in accordance with one example, (b) RHEED patterns recorded at the end of AIGaN layer growth showing clear reconstruction lines indicative of a smooth surface morphology, (c) an AFM image of the AIGaN surface showing atomically smooth surface with an r.m.s roughness 0.9 nm for a 10 x 10 pm2scanning area, and (d) a SEM image showing featureless AIGaN surface, with an inset depicting a post-TMAH etching SEM image confirming the N-polar crystallographic orientation of the epitaxially grown AIGaN.

[0015] Figure 2 depicts SEM images of N-polar AIGaN surfaces grown at different growth temperatures along with RHEED patterns (insets), including (a) a low growth temperature that leads to the formation of pits, (b) a relatively high growth temperature that leads to a smooth surface morphology, and (c) an excessively high growth temperature that leads to a rough surface.

[0016] Figure 3 depicts graphical plots of XRD asymmetrical reciprocal-space maps (RSMs) for (a) AIN / SiC, (b) Alo 4Gao eN / SiC, and (c) AloesGaossN / SiC, showing the fully relaxed and fully strained lines and the corresponding position of the maximum reflection intensity of AIGaN reciprocal lattice points.

[0017] Figure 4 depicts graphical plots of (a) normalized photoluminescence (PL) intensity plots showing compositional tunability in N-polar AlxGai.xN, from x = 0% (GaN), 20%, 40%, 65%, to 100% (AIN), (b) temperature dependent PL spectra of the 65% Al content AIGaN sample at 12 K and 300 K, and (c) variations of the integrated PL intensity versus temperature for a 200 nm Alo esGao 35N epilayer.

[0018] Figure 5 depicts graphical plots of (a) carrier concentration and Hall mobility for example Si-doped AIGaN epilayers with varying Si cell temperature, and (b) bulk resistivity as a function of Silicon cell temperature for an example 200 nm 65% Al content AIGaN epilayer.

[0019] Figure 6 is a flow diagram of a method of fabricating a heterostructure having an N- polar Ill-nitride-based layer supported by a SiC substrate in accordance with one example.

[0020] Figure 7 depicts (a) a schematic illustration of a heterostructure having an N-polar Ill-nitride-based layer supported by an SiC substrate (e.g., N-polar AIN on C-face 4H SiC substrate) in accordance with one example, (b) RHEED patterns taken after AIN epitaxy, (c)- (e) SEM images for AIN samples A, B and C, respectively, and (f) an SEM image of Sample C after etching in 25% TMAH solution, confirming the N-polarity.

[0021] Figure 8 depicts (a) an atomic force microscopy (AFM) image of Sample C of Figure 7, and (b) a graphical plot of a 2Theta-Omega XRD scan, showing the presence of SiC and AIN peaks.

[0022] Figure 9 depicts (a) a BF-STEM image taken at the interface of N-polar AIN on c- SiC for sample C of Figure 7, in which dark lines are misfit dislocations that extend and annihilate at about 200 nm from the heterointerface, (b) a high resolution HAADF-STEM) image of the heterointerface, in which the interface is marked with a dashed line, and (c) aniDPC-STEM image of the interface with the atomic models overlaid (Al: pink; N: blue, Si: yellow; C: gray) and the interface marked with a dashed line.

[0023] Figure 10 depicts 4D-STEM mapping of strain tensors near the interface of AIN / SiC for sample C of Figure 7, including (a) an virtual bright-field image of the mapping region showing the presence of misfit dislocations, in which a dashed box indicates a reference region for strain calculation, (b) a schematic view of diffraction with unit vectors selected for strain calculation, in which gi and g2 correspond to

[1100] and

[0001] directions, respectively, giving the strain maps normal (EXJ and along (eyy) the growth direction, and (c) resulting strain maps. Note that SiC at the bottom is shown only to indicate the interface and its numbers may be ignored.

[0024] Figure 11 depicts graphical plots of (a) photoluminescence spectra of Samples A and C measured at room temperature, (b) photoluminescence spectra of Sample C measured at room temperature and 20 K, (c) integrated photoluminescence intensity versus inverse temperature, and (d) Raman spectrum, in which a dashed arrow indicates the peak position (657.67 cm1) corresponding to strain-free AIN.

[0025] The embodiments of the disclosed devices and methods may assume various forms. Specific embodiments are illustrated in the drawing and hereafter described with the understanding that the disclosure is intended to be illustrative. The disclosure is not intended to limit the invention to the specific embodiments described and illustrated herein.DETAILED DESCRIPTION OF THE DISCLOSURE

[0026] Methods of fabricating heterostructures having an N-polar Ill-nitride-based layer are described. The heterostructures are grown on the carbon (C)-face of a SiC substrate. As described herein, a buffer or template layer may be grown on the SiC to serve, e.g., as a nucleation layer for the N-polar Ill-nitride-based layer.

[0027] The N-polar Ill-nitride-based layer is grown under nitrogen (N)-rich conditions. As described herein, the N-rich conditions may be established despite a lll / V ratio higher than 1 - e.g., greater than about 2. Such higher lll / N flux ratios do not necessarily result in metal rich conditions at the growth front. Indeed, N-rich conditions can also be maintained in connection with such high lll / N flux ratios at the growth front under certain conditions. For instance, at relatively high temperatures, metal atoms will be desorbed from surfaces much faster than N atoms, leading to N-rich conditions at the growth front.

[0028] The Ill-nitride-based layer is grown via molecular beam epitaxy (MBE). Compared to metal-organic chemical vapor deposition (MOCVD), MBE offers several advantages for the epitaxy of AIN, including simple and more controllable growth chemistry, reduced impurity incorporation, efficient dopant incorporation, easier polarity control, and atomically sharp interfaces.

[0029] In some cases, the disclosed methods are directed to growing an Ill-nitride-based layer composed of, or otherwise including, AIGaN. In other cases, the disclosed methods are directed to growing an Ill-nitride-based layer composed of, or otherwise including, AIN. Growth of these and other Ill-nitride-based layers on a SiC substrate is useful due to its excellent thermal conductivity, wide bandgap, low cost, and minimal lattice mismatch (e.g., about 1% between AIN (3.112 A) and 4H-SiC (3.08 A)), which allows for the seamless epitaxial growth of high-quality Ill-nitride layers with minimal defects and dislocations.

[0030] Described herein are a number of examples demonstrating the growth and characterization of high-quality N-polar AIGaN films on C-face 4H-silicon carbide (SiC) substrates by MBE. Upon optimization of the growth conditions, the example N-polar AIGaN films exhibited a crack free, atomically smooth surface (e.g., r.m.s. roughness of about 0.9 nm), and high crystal quality with low density of defects and dislocations. The N-polar crystallographic orientation of the epitaxially grown AIGaN film is unambiguously confirmed by wet chemical etching. Precise compositional tunability of the N-polar AIGaN films over a wide range of Al content and a high internal quantum efficiency of about 74% for the 65% Al content AIGaN film at room temperature were also attained. Furthermore, the examples demonstrated controllable silicon (Si) doping in high Al content (e.g., 65 %) N-polar AIGaN films with the highest mobility value (about 65 cm2 / V-s) observed, corresponding to an electron concentration of 1.1 x 1017cm3. A relatively high mobility value of 18 cm2 / V-s was also sustained for an electron concentration of 3.2 x 1019cm3, with an exceptionally low resistivity value of 0.009 D em. Such polarity-controlled epitaxy of AIGaN on SiC may be useful for fabricating heterostructures with high quality N-polar Ill-nitride semiconductors for a wide range of electronic, optoelectronic, and other devices and applications.

[0031] Also described herein are examples involving the molecular beam epitaxy of high- quality N-polar AIN grown (e.g., directly grown) on the C-face of a SiC substrate. The N-polar AIN examples were grown under conditions to exhibit an atomically smooth surface with root mean square (RMS) roughness of about 0.5 nm. Detailed scanning transmission electron microscopy (STEM) studies reveal that most dislocations are terminated / annihilated within approximately 200 nm of the AIN grown directly on SiC substrate. Epitaxial films grown on SiC with a thickness of 500 nm show X-ray diffraction (XRD) (002) and (102) rocking curvepeaks with full-width-half-maximum values of about 300 and 1000 arcsec, respectively, which compare favorably with previously reported heteroepitaxial AIN with similar thicknesses. Detailed deep UV photoluminescence studies further show that the N-polar AIN examples exhibit strong excitonic emission with a room temperature quantum efficiency exceeding 50%.

[0032] The N-polar Ill-nitride-based layers of the disclosed heterostructures may be useful as a wide-bandgap semiconductor in a wide variety of devices and applications, including, for instance, high electron mobility transistor (HEMT) devices, ultraviolet light emitting diode (LED) devices, and lasers (e.g., laser diodes), as well as various applications in high-power and high-frequency electronics. The disclosed N-polar Ill-nitride-based heterostructures and nanostructures are capable of achieving superior performance radio-frequency electronics, ultrahigh-efficiency micro-LEDs, and highly stable artificial photosynthesis including solar H2generation and CO2reduction. With its ultrawide bandgap of 6.2 eV, high breakdown field, large exciton binding energy (70 meV), thermal conductivity of 3.2 W / cm-K at room temperature, and quadratic and cubic nonlinearity, the Ill-nitride-based layers (e.g., AIN layers) of the disclosed heterostructures are useful for high power, high frequency, and high temperature electronics, acoustic resonators and filters, ultraviolet (UV) optoelectronics, integrated photonics, and other devices and applications. Thus, although described herein in connection with emissive devices, the heterostructures described herein are useful in connection with a wide variety of electronic and optoelectronic devices. The heterostructures may be used in various non-emissive optoelectronic devices, such as detectors. Moreover, other types of emissive devices may also be formed, including, for instance, various types of laser devices.

[0033] Although described herein in connection with examples involving N-polar AIGaN and AIN layers, the disclosed methods and heterostructures may be directed to other Ill-nitride- based layers. A wide variety of Ill-nitrides and Ill-nitride alloys may be grown. For instance, the disclosed methods may be directed to forming a heterostructure with an N-polar layer of InGaN, ScAIN, YAIN, or AIBN.

[0034] Although described herein in connection with examples involving a 4H-SiC substrate, the SiC substrate may have a different polytype or structure in other cases. For instance, the substrate may be composed of, or otherwise include, SiC having a 6H structure. The substrates may or may not be miscut. In non-miscut cases, the growth of the Ill-nitride-based layers may thus be on-axis.

[0035] Examples of Epitaxial Growth of N-Polar AIGaN on C-Face 4H-SiC Substrates. Samples and examples involving the epitaxial growth of AIGaN were cleaned prior to loading into a MBE growth chamber. In these cases, the substrates of the samples and examples were cleaned via a solvent cleaning process. The substrates were then further cleaned (e.g., baked and outgassed at 200 °C and 600 °C) in the MBE load-lock and preparation chamber, respectively. Such surface cleaning before MBE growth of AIN may be useful for preventing polarity inversion on oxidized regions. The samples were then annealed for 10 mins inside the growth chamber at a thermocouple temperature of 1000 °C to further promote the desorption of surface oxides, resulting in the generation of a pristine growth front.

[0036] In these examples, a Vecco GENxplor MBE system equipped with a radio frequency (RF) nitrogen (purity 99.9999%) plasma source and Knudsen effusion cells for Ga (purity 99.99999%) and Al (purity 99.99995%) was utilized to epitaxially grow a 200 nm AIN buffer layer, or template layer, followed by 200 nm AIGaN films on the C-face of 4H-SiC substrates with a miscut of 4°, as schematically shown in Figure 1 , part a.

[0037] Due to the relatively small lattice mismatch of approximately 1% between AIN and SiC, AIN serves as a useful template for the epitaxial growth of AIGaN. Furthermore, this layer serves to inhibit the incorporation of undesirable impurities, such as Si, O, and C, in the subsequently grown AIGaN epilayers. Reflection high energy electron diffraction (RHEED) was utilized to monitor the entire growth. Figure 1 , part b, illustrates the observed RHEED patterns of the AIGaN epilayer along two different azimuths. The presence of bright, streaky lines accompanied by a 2x2 reconstruction signifies the attainment of an atomically smooth surface.

[0038] To further characterize the AIGaN epilayers, atomic force microscopy (AFM) measurements were conducted, revealing a root mean square (r.m.s) roughness of 0.9 nm across a scanning area of 10 pm x 10 pm for the 65% Al content AIGaN sample, as depicted in Figure 1 , part c. The growth of N-polar Ill-nitrides is characterized by adatoms exhibiting a restricted diffusion length on N-polar surfaces, which can be attributed to their high binding energy. This phenomenon often gives rise to island formation, leading to a rough surface that renders device fabrication infeasible. In these growth examples, the lll / N ratio and growth temperature were selected (e.g., optimized) to identify a useful (e.g., optimum) growth window that facilitates the achievement of an atomically smooth surface without the formation of islands.

[0039] Due to the presence of a 4° miscut in the SiC substrates, the surface morphology is prone to step bunching, often resulting in a high surface roughness (e.g., r.m.s greater than3 nm) as has been observed in previous reports of epitaxy of AIN on similar miscut SiC substrates. In this context, the measured surface roughness of 0.9 nm for a large scanning area (10 pm x 10 pm) in the example N-polar AIGaN films, depicted in Figure 1 , part c, is exceptionally and unexpectedly low.

[0040] Figure 1 , part d, shows an SEM image of an example N-polar AIGaN surface, further confirming the absence of any significant pits or features on the surface. The N-polar crystallographic orientation of the epitaxially grown AIGaN film is unambiguously confirmed by wet chemical etching, as shown in the inset in Figure 1 , part d. The example for parts b-d of Figure 1 was grown at a temperature of about 800 °C along with a 11 l / V ratio of about 3:1 , which is further addressed below.

[0041] As shown in the SEM images of the 65% Al content AIGaN samples and example in Figure 2, parts a-c, the surface morphology evolves drastically with variations in growth temperature. A relatively low growth temperature, in the same range used for metal polar AIGaN growth by MBE, leads to the formation of pits, as seen in Figure 2, part a. Owing to the low adatom mobility on the N-polar surface, the metal atoms coalesce in some regions, which leads to localized low growth rates in these regions compared to the bulk. On the other hand, a very high growth temperature, although avoiding the formation of pits, leads to a rough surface morphology due to high desorption of Al atoms. This is also reflected in the RHEED patterns observed during the AIGaN growth as shown in the insets in Figure 2. Thus, to obtain an atomically smooth surface morphology, a growth temperature of about 800 °C was used along with a lll / V ratio of 3:1, via which the adatom mobility is enhanced owing to the high temperature but the desorption of the metal atoms is relatively low. This is underscored by the presence of bright and streaky RHEED patterns, depicted in Figure 2, part b, and the inset, along with the emergence of reconstruction lines.

[0042] The growth temperature may vary in other cases (e.g., in accordance with the 11 l / V ratio and / or other growth parameters). For instance, the growth temperature may fall in a range from about 600 degrees Celsius to about 800 degrees Celsius in other cases involving AIGaN. Other growth temperatures may be used in connection with the growth of other N- polar Ill-nitride layers.

[0043] In order to comprehensively examine the strain state of the AIGaN films grown on AIN / SiC, the reciprocal-space maps (RSMs) derived from X-ray diffraction patterns were measured. Figure 3, parts a-c, displays the RSM spectra acquired around the SiC (108) / AIN (104) peaks for three separate examples: (i) AIN, (ii) AIO 4Gao6N / AIN, and (Hi) Alo 65Ga035N / AIN epitaxially grown on C-face 4H-SiC. The vertical and tilted dashed linesrepresent fully strained and fully relaxed positions of the AIGaN films grown on AIN layer, respectively. The presence of a well-defined center in the mosaic spread of both the AIN and AIGaN spots indicate that the growth of the AIN and AIGaN epilayers was characterized by a high degree of crystalline alignment and structural coherence. It can be observed that with increasing Al content from 40 % to 65 %, the out-of-plane reciprocal space lattice constant Qzmoves closer to the AIN peak. Moreover, with increase in Al content, it is observed that the maximum reflection intensity of AIGaN reciprocal lattice points (RLPs) gets closer to the fully strained position, indicating more pseudomorphic behavior.

[0044] AIGaN is well-suited for deep UV (DUV) optoelectronics due to its direct ultrawide bandgap and the ability to cover the UV-A, UV-B, and UV-C spectra through precise alloy composition tuning. Figure 4, part a, shows the room temperature photoluminescence characteristics of the AIGaN examples.

[0045] By adjusting the Al / Ga flux ratio, the Al composition in the N-polar AIGaN films was systematically tuned from 0% (GaN), 20%, 40%, 65%, to 100% (AIN). This controlled variation was verified by the shift in the band edge emission peak position observed in the corresponding photoluminescence spectrum. Despite significant advancements in AIGaN- based materials and devices, the successful integration of AIGaN-based optoelectronic devices, especially those requiring high Al compositions, faces substantial challenges, primarily attributed to the formidable challenge of attaining high-quality material. AIGaN alloys, when grown heteroepitaxially on sapphire and silicon substrates, commonly encounter a notable presence of both extended and point defects owing to a large lattice mismatch. Employing SiC as the substrate, as described herein, is a useful approach to circumvent this challenge owing to a more favorable lattice approximation with AI(Ga)N.Additionally, the formation of three-dimensional (3D) island-like growth modes, particularly in N-polar growth, is influenced by the high adhesion coefficient and low surface migration velocity of Al atoms, leading to an elevated dislocation density in AIGaN epilayers. These dislocations typically serve as nonradiative recombination centers in active devices based on AIGaN. In the disclosed methods, a relatively high growth temperature along with a high metal flux helps to enhance the adatom mobility and yet maintain a residual metal flux to compensate for the adsorption at elevated temperatures, resulting in high-quality N-polar AIGaN films.

[0046] T o assess the radiative recombination efficiency of the epitaxially grown AIGaN examples, temperature-dependent photoluminescence measurements were conducted. Figure 4, part b, presents the PL spectra of Alo.65Gao.35N as the temperature decreases from 300 K to 12 K while maintaining a pumping power of 30 pW. The estimation of internalquantum efficiency (IQE) values was achieved by comparing the integrated intensity of the photoluminescence emission at room temperature (RT) to that at low temperature (12 K), assuming that IQE at 12 K is 100% due to the inactivation of defect-related nonradiative recombination centers.

[0047] As shown in Figure 4, part c, a high IQE of about 74% was attained by the Alo65Gao35N example, demonstrating a large radiative recombination rate and indicating an absence of significant defects or dislocations in the epitaxially grown films. The formation of Ga rich nanoclusters has been commonly observed in AIGaN epilayers. The resulting charge carrier localization and quantum-confinement contributes to the significantly enhanced photoluminescence intensity and high IQE. The formation of Ga-rich nanoclusters in the AIGaN epilayers also leads to significant asymmetric broadening of the PL spectrum to higher wavelengths for AIGaN epilayers compared to GaN or AIN, as observed in Figure 4, part b.

[0048] To integrate with existing electronic and optoelectronic platforms, successful doping and precise control of conductivity are useful. While achieving controllable n-type doping in low Al content AIGaN films is attainable, it proves to be challenging to attain highly conductive AIGaN with a high Al content. This difficulty primarily arises from factors such as dislocation densities, high donor activation energy, compensation effects caused by acceptor-like defects, negatively charged Al, Ga vacancies, and Si DX centers. Therefore, examples focusing on Si doping in high Al content (e.g., 65%) N-polar AIGaN films were fabricated and explored.

[0049] The carrier concentrations measured in example Alo.65Gao.35N epilayers are shown in Figure 5, part a, for five different Si cell temperatures, ranging from 1125 °C to 1225 °C. Effective modulation of the electron concentration was achieved, spanning several orders of magnitude, which greatly enhances the versatility of Si doped N-polar AIGaN in various application domains.

[0050] The relationship between mobility and Si cell temperature is also illustrated in Figure 5, part a. The highest mobility value of approximately 65 cm2 / V-s is observed corresponding to an electron concentration of 1.1 x 1017cm3. Notably, even at the higher range of doping, a relatively high mobility value of 18 cm2A / -s is sustained for an electron concentration of 3.2 x 1019cm3. This finding highlights the feasibility of maintaining desirable mobility characteristics across a wide range of doping levels. The Si-doped AIGaN films demonstrated exceptional electrical conductivity (shown in Figure 5, part b), as indicated by the lowest resistivity value of 0.009 Q.cm reported for AIGaN in a similar compositionalrange. Table I below presents a comparison with some of the lowest resistivity Si doped high Al content AIGaN films reported in the literature.Table I: Comparison of resistivity for high Al content AIGaN films.Ref. Composition Thickness pbuik (Q.cm) Growth methodThis i Alo e5Gao 35N 200 nm 0.009 MBE disclosure i

[0051] The results shown in Table I demonstrate that the Si-doped high Al content N-polar AIGaN epilayers disclosed herein exhibit remarkably low resistivity values within the specified thickness range, which also indicates that the formation of VAI + nSiAicomplexes and other self-compensating factors is suppressed. Moreover, such a low resistivity obtained for the Si doped AIGaN films suggest a transition to a metallic conductivity state in the higher doping range, as has been demonstrated recently for highly doped n-AI07Ga03N films (greater than 1019cm3) with a vanishing of the effective dopant activation energy, indicating almost all the Si dopants are effectively activated in this range of doping.

[0052] Figures 1 -5 provide examples of the fabrication of high quality N-polar AIGaN films on C-face 4H-SiC substrates with compositions varying across the entire alloy range, which exhibit atomically smooth surface, high luminescence emission efficiency in the deep UV, and excellent charge carrier transport properties. The achievement of N-polar AIGaN supported by, e.g., directly on, SiC wafers, with exceptional structural, electrical, and optical attributes, is useful for high-power, high-frequency, and high-temperature electronics, as well as applications involving high efficiency deep UV optoelectronics and photonics.

[0053] Figure 6 depicts a method 600 of fabricating a heterostructure device having an N- polar Ill-nitride-based layer in accordance with one example. The method 600 may be used to manufacture any of the heterostructures described herein or another type of heterostructure. The method 600 may include additional, fewer, or alternative acts. For instance, the method 600 may or may not include one or more acts directed to forming other device structures (act 622).

[0054] The method 600 may begin with an act 602 in which a substrate is prepared or otherwise provided. The substrate may be, be formed from, or otherwise include 4H-SiC. Other SiC substrates may be used, including, for instance, 6H-SiC.

[0055] Preparation of the substrate may include an act 604 in which one or more chemical or thermal cleaning procedures are implemented. Alternatively or additionally, the act 602 include an act 604 in which the substrate is annealed (e.g., in the MBE chamber in which the layer(s) of the heterostructure are grown).

[0056] The method 600 includes an act 610 in which a buffer or template layer of the heterostructure is grown on the substrate. As described herein, the buffer layer may be composed of, or otherwise include, AIN. The act 610 may include implementation of an MBE procedure. The MBE procedure may be implemented in an act 612 in a common (the same) growth chamber used to grow the N-polar Ill-nitride layer. The MBE procedure may or may not be implemented under nitrogen-rich conditions.

[0057] In an act 614, an MBE procedure is implemented to grow the N-polar Ill-nitride- based layer. As described herein, the MBE procedure may be configured such that the Ill- nitride-based layer is N-polar. For instance, the act 614 may include an act 616 in which the MBE procedure is implemented under N-rich conditions. The N-rich conditions may be established or influenced by the selection of a number of process parameters, including growth temperature and 11 l / V ratio. For instance, the N-rich conditions may be established despite the use of a 11 l / V ratio of about 2:1 or higher.

[0058] The MBE procedure may be configured to dope the N-polar Ill-nitride-based layer (e.g., with silicon) in an act 618. The Al (or other Group III) composition may be controlled in an act 620.

[0059] In some cases, the method 600 may include an act 622 directed to forming one or more structures of a device in which the heterostructure is integrated.

[0060] Figures 7-11 are directed to examples involving the fabrication of heterostructures having an N-polar AIN layer. Examples are schematically shown in Figure 7, part a. In these cases, AIN epilayers were grown on C-face SiC substrates utilizing a Vecco GENxplor MBE system equipped with a radio frequency plasma-assisted nitrogen source. As described above, the initial nucleation process is useful for achieving high-quality AIN and directly impacting the segregation and incorporation of undesired impurities, e.g., Si, O, and C, in the subsequently grown AIN epilayers.

[0061] The SiC substrates were prepared via multiple steps to promote the nucleation and epitaxy of AIN epilayers. In these examples, the C-face SiC substrate was thermally annealed at 920 °C for 30 min to acquire a clean epitaxial surface. A 30 nm thick AIN nucleation (or buffer) layer was then grown at a relatively low substrate temperature under slightly nitrogen-rich conditions to help suppress the segregation and incorporation of Si in the epilayers.

[0062] After growth of the buffer layer, 500-nm-thick AIN was grown at a substrate temperature of 985 °C. Reflection high energy electron diffraction (RHEED) was utilized to monitor the growth process. Bright, sharp, and streaky RHEED patterns are clearly observed from two different azimuths of AIN, when grown under optimized conditions, shown in Figure 7, part (b), indicating an atomically smooth surface.

[0063] The growth conditions for the N-polar AIN layers, however, are different from those of conventional metal-polar AIN. Detailed studies of the effect of growth parameters on the structural and optical properties of AIN on C-face SiC were performed. The growth details of a few representative samples are listed in Table II below, with the corresponding scanning electron microscopy (SEM) images shown in Figure 7, parts (c), (d) and (e), respectively. The use of N-rich conditions can significantly suppress the incorporation of undesired Si incorporation in the subsequently grown AIN epilayers. Under N-rich growth conditions (Sample A), however, the surface exhibits poor morphology, with the presence of discrete islands, shown in Figure 7, part (c), which is not suited for practical device applications. The utilization of metal-rich conditions (Sample B), on the other hand, leads to the presence of metal droplets, discontinuous growth, and extensive voids and pits on the surface, shown in Figure 7, part (d). It is also observed that the growth temperature plays an important role. A low growth temperature results in low adatom migration length. This further increases the surface roughness by generating higher density of hillocks and higher temperature reduces surface roughness. A relatively high growth temperature, on the other hand, can effectively promote surface atom migration, resulting in atomically smooth surface. Illustrated in Figure 7, part (e), a featureless, smooth surface is obtained for N-polar AIN grown directly on SiC by utilizing a relatively high growth temperature of 985 °C (Sample C). The N-polarity is further confirmed by the presence of pyramid shape morphology, shown in Figure 7, part (f), after etching the sample in 25% TMAH base solution.Table I. List of growth parameters for a few representative N-polar AIN epilayer samples.Sample _ III / V Ratio _ Growth Temperature (°C)

[0064]

[0065] The example N-polar AIN epilayers were characterized by atomic force microscopy (AFM), which shows a root mean square (RMS) roughness of 0.55 nm over a scanning area of 3 pm x 3 pm, illustrated in Figure 8, part (a). Structural characterization of AIN epilayers was further performed by X-ray diffraction (XRD) rocking curve measurement. Shown in Figure 8, part (b), is the XRD 2theta-omega scan, with the two peaks corresponding to SiC and AIN being clearly identified. Full-width-at-half-maximum (FWHM) values of 328 arcsec and 1000 arcsec for the (002) plane and (102) plane were measured, respectively, which compare favorably with previous reports of N-polar AIN epilayers of a similar thickness grown on a foreign substrate.

[0066] Structural properties of the N-polar AIN examples were further characterized by scanning transmission electron microscopy (STEM). Figure 9, part (a), shows a representative bright-field STEM (BF-STEM) image acquired at the interface of AIN and SiC. Misfit dislocations as dark lines are observed to be generated at the interface and extend into the AIN lattice. Significant annihilation of these dislocations occurs at about 200 nm away from the interface. This annihilation happens during AIN epitaxy growth. A high resolution high-angle annular dark-field STEM (HAADF-STEM) image of the AIN / SiC interface is present in Fig. 9, part (b). The AIN on the upper part shows the typical “zig-zag” pattern from a wurzite lattice viewed along

[1120] , whereas the 4H-SiC bilayer zig-zag stacking sequence is observed on the lower part. The transition of such lattice patterns is atomically sharp at the interface, without observable lattice discontinuity, indicating good epitaxy of AIN grown directly on SiC. An integrated differential phase contrast STEM (iDPC- STEM) image highlighting the interface is shown in Fig. 9, part (c). As a complement to the HAADF-STEM image, iDPC-STEM shows clearly the N columns in AIN and C columns in SiC (also overlaid with atomic structure models of the wurzite AIN and 4H-SiC), confirming that the AIN is indeed N-polar, which is consistent with the etching studies shown in Figure 7, part (f).

[0067] Despite the relatively small lattice mismatch (about 1%) between AIN and 4H-SiC, there exists a large difference between thermal expansion coefficients, i.e., 4.2 x w6K1and 3.21 x w6K'1for AIN and SiC, respectively. Such a large difference in thermal expansion coefficient can induce significant tensile strain distribution in AIN epilayers when grown at elevated temperatures. To verify the strain distribution in the AIN layer, a 4D-STEMtechnique was used, in which a series of diffraction patterns were recorded while the electron beam was scanned from the top surface to the interface of AIN / 4H-SiC. Lattice distortions were calculated from the diffraction pattern at each scanning position by taking the top surface of AIN as reference (dashed box in Figure 10, part a). A virtual bright-field image was generated by integrating the approximate 10 mrad center region of each diffraction pattern (Figure 10, part a), showing misfit dislocations within about 200 nm region from the interface. The unit vectors, gi and g2corresponding to

[1100] and

[0001] directions (normal and parallel to the growth direction), respectively, as shown in Figure 10, part b, were used to calculate the strain maps presented in Figure 4, part c. It is evident that indeed the AIN layer suffers from tensile strains in both

[1100] (exx) and

[0001] (eyy) directions. Accompanied there are also some small shear strains (exy) and rotations (0). Note that the calculations in SiC may be ignored as the reference lattice is only valid for AIN. They are shown on the map only to indicate the location of the interface.

[0068] Optical properties of the N-polar AIN examples grown on SiC were studied by deep UV photoluminescence spectroscopy. As shown in Figure 11 , part (a), significantly improved optical properties are measured for the sample grown under high temperature, which is also consistent with the reduction of dislocation densities confirmed by STEM characterization shown in Figure 9. For the sample grown under optimized conditions, temperaturedependent photoluminescence measurements were performed. Shown in Figure 11 , part (b), are the photoluminescence spectra measured at 300 K and 20 K. Figure 11 , part (c), shows the temperature dependence of the integrated photoluminescence intensity, plotted vs. the inverse of temperature. Given the large exciton binding energy (about 70 meV) of AIN, it is expected that the photoluminescence emission originates from a bound state of an electron and hole, instead of the recombination of free charge carriers. The effect of Shockley-Read- Hall recombination is expected to be less important, due to the strong Coulombic interaction between electrons and holes. If one assumes that a near unity quantum efficiency at 20 K for the excitonic emission, a relatively high internal quantum efficiency of 52% at roomtemperature is derived by comparing the integrated photoluminescence intensity measured at room-temperature with that at 20 K. It shows the potential for applications in deep UV optoelectronics.

[0069] The measured photoluminescence emission peak is at 5.916 eV (209.6 nm) at room temperature, which is redshifted compared to that of strain-free AIN, i.e., 6.026 eV (205.75 nm). The presence of large tensile stress in AIN grown on SiC and the resulting redshift of the photoluminescence emission has been reported previously. While the 4D-TEM results showed intuitively the strain state of the AIN films, the shift in PL peak position gives a goodmeasure of the tensile strain in the epilayer. In this example, the tensile strain is calculated to be 2.22 GPa utilizing the following empirical equation, which is within a reasonable range compared with other reported values.

[0070] Eg[eV] = £g[0] + 3.6 x 10“3P - 1.7 x 10“6P2[kbar]

[0071] The strain distribution, however, may vary considerably with temperature. With further reducing to cryogenic temperature, a non-negligible red-shift in the photoluminescence emission will be induced by additional tensile strain, which can compensate, to some extent, the blue-shift expected for the photoluminescence emission with reducing measurement temperature. Shown in Figure 11 , part (b), the photoluminescence emission measured at 20 K only showed a very small (0.8 nm) blue shift, compared to that measured at room temperature, due to the two combined effects.

[0072] The biaxial strain of N-polar AIN grown on SiC was further examined by microRaman spectroscopy. Shown in Figure 11 , part (d), is the measured Raman spectrum, with the E2(high) mode emission at 653 cm1. The dashed line indicates the peak position (657.67 cm1) expected for strain-free AIN. Based on the theoretical predicted shift of 2.55 cm VGPa, the residual stress in our AIN epilayer is estimated to be 1.73 GPa, which is in good agreement with analysis based on photoluminescence studies. Despite the relatively large strain distribution, the FWHM of the E2mode is relatively narrow (about 5.1 cm1), which compares favorably with previously reported high quality AIN on SiC or sapphire, suggesting excellent material quality.

[0073] The examples described above demonstrate the molecular beam epitaxy and characterization of high-quality N-polar AIN directly on C-face 4H-SiC substrate. By controlling and optimizing the growth conditions, the N-polar AIN grown on SiC exhibits excellent structural and optical properties, including atomically smooth surface and high luminescence efficiency in the deep UV at room temperature. Most dislocations are terminated near the heterointerface. However, the AIN epilayers grown on SiC exhibit relatively large tensile strain, which leads to a considerable redshift of the photoluminescence and micro-Raman emission. The realization of high-quality N-polar AIN on SiC is useful in various applications in high power, high frequency, and high temperature electronics, high frequency acoustic resonators and filters, as well as deep UV optoelectronics and integrated photonics.

[0074] Described above are examples of fabrication methods including the growth of N- polar AIGaN on C-face SiC by MBE. The epitaxy and structural, optical, and electrical properties of examples of N-polar AIGaN on C-face 4H-SiC substrates with compositionsvarying across the entire alloy range were described. The epitaxially grown N-polar surfaces were atomically smooth and free of pits or cracks. The N-polar orientation of the epitaxial AIGaN layers was confirmed by wet chemical etching. The strain-state analysis by XRD reciprocal space mapping reveals an in-plane compressive stress in the epitaxially grown AIGaN films, that reduces with increasing Al content. The Al content in AIGaN was varied from 0% (GaN) to 20%, 40%, 65%, and 100% (AIN) to enable a precise control over the bandgap, leading to tunable photoluminescence (PL) characteristics. Temperaturedependent PL measurements conducted on the AIGaN alloy with a high aluminum (Al) content of 65% demonstrate a high internal quantum efficiency (IQE) of 74% at roomtemperature. Precise control over Si doping spanning multiple orders of magnitude has been achieved in high aluminum (Al) content (65%) AIGaN films. The measured Hall mobility reaches up to 65 cm2 / V-s at an electron concentration of 1 .1 x 1O17cm-3, and a relatively high mobility value of 18 cm2A / -s is sustained for an electron concentration of 3.2 x 1O19cm-3, accompanied by an exceptionally low resistivity of 0.009 Q cm

[0075] Also described above are examples of fabrication methods including the growth of N-polar AIN on C-face 4H-SiC substrates. The N-polar AIN films grown exhibited atomically smooth surfaces and strong excitonic emission in the deep UV with luminescence efficiency exceeding 50% at room temperature. Detailed scanning transmission electron microscopy (STEM) studies indicated that most dislocations were terminated / annihilated within approximately 200 nm AIN grown directly on SiC substrate, thanks to the relatively small (1%) lattice mismatch between AIN and SiC.

[0076] As used herein, the terms "atomically smooth" or "atomically smooth surface" may be used herein in connection with a layer of a heterostructure to indicate that the layer has a surface roughness (e.g., a root mean square, or RMS, roughness) less than or on the order of 1 nm. In some cases, the RMS roughness of such atomically smooth layers is less than 1% of the thickness of the layer. The surface roughness may vary in accordance with the growth conditions, parameters, and other aspects of the fabrication processes described and / or referenced herein and / or other processes.

[0077] The term "about" is used herein in a manner to include deviations from a specified value that would be understood by one of ordinary skill in the art to effectively be the same as the specified value due to, for instance, the absence of appreciable, detectable, or otherwise effective difference in operation, outcome, characteristic, or other aspect of the disclosed methods and devices.

[0078] The present disclosure has been described with reference to specific examples that are intended to be illustrative only and not to be limiting of the disclosure. Changes, additions and / or deletions may be made to the examples without departing from the spirit and scope of the disclosure.

[0079] The foregoing description is given for clearness of understanding only, and no unnecessary limitations should be understood therefrom.

Claims

What is Claimed is:

1. A method of fabricating a heterostructure, the method comprising: providing a substrate, the substrate comprising silicon carbide, the substrate having a carbon (C)-face; forming a Ill-nitride-based buffer layer of the heterostructure, the Ill-nitride-based buffer layer being in contact with the C-face of the substrate; and implementing a molecular beam epitaxy (MBE) procedure to grow a Ill-nitride-based layer of the heterostructure such that the Ill-nitride-based layer is supported by the Ill-nitride- based buffer layer; wherein the molecular beam epitaxy (MBE) procedure is implemented under nitrogen (N)-rich conditions.

2. The method of claim 1 , wherein the molecular beam epitaxy (MBE) procedure implemented to grow the Ill-nitride-based layer has a 11 l / V ratio of about 2:1 or higher.

3. The method of claim 1 , wherein the molecular beam epitaxy (MBE) procedure implemented to grow the Ill-nitride-based layer has a 11 l / V ratio of about 3:1 .

4. The method of claim 1 , wherein providing the substrate comprises cleaning the substrate via a solvent cleaning process.

5. The method of claim 1 , wherein providing the substrate comprises annealing the substrate in a growth chamber in which the MBE procedure is implemented at a temperature sufficient to remove oxide from the C-face of the substrate.

6. The method of claim 1 , wherein the MBE procedure to grow the Ill-nitride-based layer is implemented at a growth temperature that falls in a range from about 600 degrees Celsius to about 800 degrees Celsius.

7. The method of claim 6, wherein the Ill-nitride-based layer comprises AIGaN.

8. The method of claim 1 , wherein the MBE procedure to grow the Ill-nitride-based layer is implemented at a growth temperature that falls in a range from about 800 degrees Celsius to about 1100 degrees Celsius.

9. The method of claim 8, wherein the Ill-nitride-based layer comprises AIN.

10. The method of claim 1 , wherein forming the buffer layer comprises growing the buffer layer via molecular beam epitaxy in a growth chamber in which the Ill-nitride-based layer is grown.

11. The method of claim 1 , wherein the Ill-nitride-based buffer layer comprises AIN.

12. The method of claim 1 , wherein the MBE procedure is configured to dope the Ill- nitride-based layer with silicon.

13. The method of claim 1 , wherein the Ill-nitride-based layer comprises AIN.

14. The method of claim 1 , wherein the Ill-nitride-based layer comprises AIGaN.

15. The method of claim 1 , wherein the substrate has a 4H-SiC structure.

16. The method of claim 1 , wherein the MBE procedure is configured such that the Ill- nitride-based layer is a ternary compound having an aluminum composition higher than another Group III element composition of the ternary compound.

17. The method of claim 1 , wherein the substrate is a miscut SiC substrate.

18. The method of claim 1 , wherein the Ill-nitride-based layer is in contact with the Ill- nitride-based buffer layer.

19. A method of fabricating a heterostructure, the method comprising: providing a substrate, the substrate comprising silicon carbide, the substrate having a carbon (C)-face; growing a Ill-nitride-based template layer of the heterostructure, the Ill-nitride-based template layer being in contact with the C-face of the substrate; and implementing a molecular beam epitaxy (MBE) procedure to grow a Ill-nitride-based layer of the heterostructure such that the Ill-nitride-based layer is in contact with the Ill- nitride-based template layer; wherein: providing the substrate comprises annealing the substrate in a growth chamber in which the MBE procedure is implemented; growing the template layer is implemented in the growth chamber in which the Ill- nitride-based layer is grown; and the molecular beam epitaxy (MBE) procedure to grow the Ill-nitride-based layer is implemented under nitrogen (N)-rich conditions.

20. The method of claim 19, wherein the Ill-nitride-based template layer comprises AIN.

21. The method of claim 19, wherein the Ill-nitride-based layer comprises AIN.

22. The method of claim 19, wherein the Ill-nitride-based layer comprises AIGaN.

23. The method of claim 19, wherein the substrate has a 4H-SiC structure.

24. A method of fabricating a heterostructure, the method comprising: providing a substrate, the substrate comprising silicon carbide, the substrate having a carbon (C)-face; forming a Ill-nitride-based buffer layer of the heterostructure, the Ill-nitride-based buffer layer being in contact with the C-face of the substrate; and implementing a molecular beam epitaxy (MBE) procedure to grow a Ill-nitride-based layer of the heterostructure such that the Ill-nitride-based layer is supported by the Ill-nitride- based buffer layer; wherein the substrate is not miscut such that growth of the Ill-nitride-based buffer layer and the Ill-nitride-based layer is on-axis.

25. The method of claim 24, wherein the MBE procedure is implemented at a growth temperature and with a 11 l / V ratio such that the Ill-nitride-based layer is atomically smooth.

26. The method of claim 25, wherein the 111 A / ratio is about 2:1 or higher.

27. The method of claim 26, wherein the growth temperature falls in a range from about 600 degrees Celsius to about 800 degrees Celsius.

28. The method of claim 26, wherein the growth temperature falls in a range from about 800 degrees Celsius to about 1100 degrees Celsius.

29. The method of claim 25, wherein the MBE procedure is configured such that the Ill- nitride layer is N-polar.

30. The method of claim 25, wherein the MBE procedure is implemented under nitrogen (N)-rich conditions.

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