GROWTH OF III-NITRIDES VIA SUBLIMATION AND METALORGANIC VAPOR PHASE EPITAXY
RAST HI-NITRIDOV S SUBLIMACIJO IN METALORGANSKO PARNO FAZNO EPITAKSIJO
ROBERT F. DAVIŠ', C. M. BALKAS1, M. D. BREMSER1, O. H. NAM1, W. G. PERRY', B. L. WARD2, Z. SITAR1, T. ZHELEVA1, L. BERGMAN2,1. K. SHMAGIN3, J. F. MUTH3, R. M. KOLBAS3, R. J. NEMANICH2
'Department of Materials Science and Engineering, North Carolina State University, Box 7907 Raleigh, NC, 27695-7907 2Department of Physics, North Carolina State University, Bo\ 8202, Raleigh, NC 27695-8202 3Department of Electrical and Computer Engineering, North Carolina State University, Box 7911, Raleigh, NC 27695-7911
Prejem rokopisa - received: 1997-10-01; sprejem za objavo - accepted for publication: 1997-10-21
Single crystals of GaN < 3 mm in length were grown by sublimation/recondensation of GaN in 760 Torr NH3 at 1100°C. Platelets of A1N < 1 mm thick were similarly grown between 1950 and 2250°C using an Al source. Monocrystalline GaN and AlxGa,.xN(0001) (0.05 < x < 0.96) films were grown via MOVPE on a(6H)-SiC(0001) wafers with and without, respectively, a 1000 A A1N buffer layer. Photoluminescence (PL) spectra of GaN showed bound and free excitonic recombinations. Selective growth of hexagonal pyramid arrays of undoped GaN and Si-doped GaN vvas achieved on 6H-SiC(0001)/AIN/GaN multilayer substrates using a patterned Si02 mask. Field emission of these arrays exhibited a tum-on field of 25 V/(im for an emission current of 10.8 nA at an anode-to-sample distance of 27 jim. Lateral growth and coalescence of GaN have been achieved using stripes oriented along <1100> at 1100°C and a triethylgallium flow rate of 26 mmol/min. Approximately 109 cm"2 dislocations, originating from the underlying GaN/AlN interface, were contained in the GaN grovvn in the vvindovv regions. The overgrovvth regions contained a very low density of dislocations.
Key words: gallium nitride, aluminum nitride, single crystals, thin films, photoluminescence, Raman spectroscopy, cathodoluminescence, selective growth, lateral overgrovvth
Kristali GaN < 3 mm dolžine so bili izdelani s sublimacijo/rekondenzacijo GaN pri 760 tor. NH3 in 1100°C. Ploščice A1N < 1 mm debeline so bile na podoben način izdelane pri 1950 do 2250°C z uporabo Al kot izvora. Monokristalini filmi GaN in AlxGai.x N(001) z (0.05 S x < 0.96) so bili izdelani z MOVPE na a(6H)-SiC(0001) vaferjih z in brez 1000 A A1N bufer sloja. Fotoluminiscentni (PL) spektri GaN so pokazali vezi in proste excitonske rekombinacije. Selektivna rast združb heksagonalnih piramid nedopiranega GaN in GaN dopiranega s silicijem je bila dosežena na 6H - Si(0001)/AlN/GaN večslojnih substratih z uporabo vzorčastih SiOj mask. Prag polja emisije teh združb je imel jakost 25 V/)jm za emisijski tok 10.8 nA pri razdalji 27 pm med anodo in preizkušancem. Lateralna rast in koalescenca GaN je bila dosežena z uporabo lamel orientiranih vzdolž <1100> pri 1100°C in pretoku trietilgalija 26 mmol/min. V GaN, ki je nastal v področju oken je bilo približno 109 cm"2 dislokacij, ki so izvirale iz mejne površine GaN/AlN. Področja večje rasti so imela majhno gostoto dislokacij.
Ključne besede: galijev nitrid, aluminijev nitrid, posamični kristali, tanki filmi, fotoluminiscenca, Raman spektroskopija, katodoluminiscenca, selektivna rast, pospešena lateralna rast
1INTRODUCTION
The realization of blue and green light emitting di-odes and blue lasers as well as prototypes of several mi-croelectronic devices produced from GaN-based materials containing copious line and planar defects has been most fortunate. These achievements also indicate that the employment of substrates on which homoepitaxial films can be grovvn would result in marked improvements in the properties of the devices fabricated in these films. At present, very thin films of GaN, A1N and AlxGa|.xN are deposited on foreign substrates as buffer layers on which the device-related III-V nitrides films are grovvn. Aluminum nitride is also a candidate material for selected pie-zoelectric applications and surface acoustic vvave (SAW) devices. Hovvever, the potential of A1N in these and other applications has been hampered by the lack of bulk single crystals, as discussed in several revievv articles'"4. Gallium nitride is a promising material for field emission because of its Iow electron affinity (2.7-3.3eV)5 6, reason-able thermal, chemical and mechanical stability and the
ability for controlled n-type doping. Recently, it has been reported6-7 that A1N and Al-rich AlxGa,.xN (x > 0.75) films exhibit a negative electron affinity vvhich suggests that these materials belong to a special class of field emitters.
Recent research conducted by the authors and de-scribed in the follovving sections represent important ad-vances in the determination of the process parameters necessary to achieve grovvth of GaN and A1N bulk single crystals via seeded sublimation/recondensation. Addi-tionally, we discuss the employment of a 1000 A, monocrystalline, high-temperature (HT) (1100°C) A1N buffer layer for metalorganic vapor phase epitaxy (MOVPE) thin film deposition vvhich has resulted in subsequently deposited GaN films void of oriented do-main structures and associated lovv-angle grain bounda-ries8 9. Monocrystalline films of AlxGai_xN (0.05 < x < 0.96) of the same quality as GaN vvith a HT-AIN buffer layer have also been achieved directly on 6H-SiC(0001) vvafers at 1100°C. We also report the selective grovvth of GaN and Si-doped GaN hexagonal pyramid arrays on
circular patterns etched in Si02 masks deposited on GaN/AlN/6H-SiC(0001) multilayer substrates and the field emission results from these arrays, as well as the deposition, the lateral overgrowth and subsequent coales-cence of the GaN stripes selectively grown in the same manner.
2 EXPERIMENTAL PROCEDURES
A.	GaN Bulk Growth
Growth of individual GaN crystals was achieved by evaporating 0.5" diameter and 0.25" high GaN pellets cold pressed from high purity GaN povvders produced in our laboratory10 in a stream of 99.9999% pure NH, gas. Experiments were conducted in a system that was spe-cifically designed for GaN growth. The growth system consisted of: (1) an outer vacuum chamber that con-tained the heat shields and electrical feed-throughs and (2) a reaction tube containing the source and seed mate-rials. Two independently controlled heaters vvere used to achieve the necessary temperature gradients for sublima-tion growth. Heater temperatures vvere monitored and controlled via the use of thermocouples placed adjacent to each heater. Source and substrate temperatures vvere monitored independently from within the reaction tube. Experiments vvere performed vvithin the source temperature range of 1100-1450°C. The seed heater was main-tained at 1100°C throughout each deposition. Growth pressures of 50-760 torr vvere investigated; hovvever, most experiments vvere conducted at 760 torr. An NH3 flow rate of 50 sccm was used for ali grovvth experi-ments. Blank BN seed holders vvere used as substrates and placed at a distance from the source vvherein the crystals vvith lovv aspect ratios could be obtained.
B.	AIN Bulk Growth
Aluminum nitride sublimation/recondensation ex-periments vvere also conducted in a resistively heated graphite furnace. Bulk AIN (99% dense) blocks produced via sintering vvithout additives vvere used as the source material. The source vvas positioned in the iso-thermal section of the furnace to ensure an essentially constant evaporation rate. Single crystal, 6H-SiC (0001) squares (10 mm x 10 mm) vvere used as seeds in ali ex-periments due to the relatively small lattice mismatch to AIN (0.9%) and high temperature stability. The seed crystals vvere heated in vacuum at =1150°C prior to crys-tal grovvth to desorb the surface oxide, hydrocarbons and any other contaminants.
Ali AIN experiments vvere performed under a 100 sccm flow of ultra-high purity N2. The background pressure vvas maintained at 500 Torr by an automatic throttle valve. The temperature ranges of 2100-2250°C and 1950-2050°C vvere investigated consecutively. The lovver temperature range vvas employed primarily be-cause of the degradation of the furnace and the seed
crystals at the higher temperatures. In both cases, a temperature difference of 80-150°C vvas employed, depend-ing on the separation distance (1-40 mm) betvveen the source and the seed. The grovvth rate at a separation of 3 mm vvas =30 times higher than at a 15 mm separation. The AIN source vvas repositioned to the desired source height before each experiment.
C. Metalorganic Vapor Phase E pitaxy and Selective
Area Growth
As-received vicinal 6H-SiC(0001) vvafers11 oriented 3°-4° off-axis tovvard <1120> vvere cut into 7 mm squares. These pieces vvere degreased in sequential ultra-sonic baths of trichloroethylene, acetone and methanol and rinsed in deionized vvater. The substrates vvere sub-sequently dipped into a 10% HF solution for 10 minutes to remove the thermally grovvn oxide layer and blovvn dry vvith N2 before being loaded onto the SiC-coated graphite susceptor contained in a cold-vvall, vertical, pancake-style, MOVPE deposition system. The system vvas evacuated to <3xl0'5 Torr prior to initiating grovvth. The continuously rotating susceptor vvas RF inductively heated to the GaN (AlxGa,.xN) deposition temperature of 1050°C (1100°C) (optically measured on the susceptor) in 3 SLM of flovving H2 diluent. Hydrogen vvas also used as the carrier gas for the various metalorganic precursors. Deposition of AlxGai.xN vvas initiated by flovving various ratios of triethylaluminum (TEA) and triethylgallium (TEG) in combination vvith ammonia (NH3).
Selective grovvths of GaN and Alo.2Gao.8N vvere achieved at 1000-1050°C vvith TEG flovv rates = 26.1-70.0 pm/min. on stripe (vvindovv vvidth = 3-80 pm) and circular (diameter = 5 pm) patterned GaN/AlN/6H-SiC(0001) multilayer substrates. To produce these patterned substrates, a Si02 mask layer (thickness = 1000A) vvas subsequently deposited on each GaN film via RF sputtering or lovv pressure chemical vapor deposition. Patterning of the mask layer vvas achieved using standard photolithography techniques and etching vvith a buffered HF solution. The edges of the stripe patterned samples vvere parallel to <1120>. Prior to selective grovvth, the patterned samples vvere dipped in a buffered HC1 solution to remove the surface oxide of the underlying GaN layer. Incorporation of the n-type Si dopant into the GaN pyramids during grovvth vvas achieved using SiH4 at a flovv rate of 5.5 nmol/min.
Field emission measurements (FEM) vvere performed on the Si-doped GaN hexagonal pyramid arrays in a UHV-FEM system having a vvorking pressure of 2x10S Torr. Each array vvas placed beneath a five mm diameter movable Mo anode having a flat tip. The anode vvas controlled by a stepping motor such that one step yielded a translation of 0.44 pm. The current-voltage (I-V) measurements vvere taken from 2 to 40 pm for anode voltages in the range of 0 to 1100 V.
The lateral overgrowth of GaN was achieved in a manner similar to that of the films12. The GaN grew ver-tically to the top of the mask and then both laterally and vertically over the mask until the lateral growth fronts from many different windows coalesced and formed a continuous layer. The samples were characterized by scanning electron microscopy (SEM-JEOL 6400 FE), atomic force microscopy (AFM-Digital Instrument NanoScope III) and transmission electron microscopy (TEM-TOPCON 002B, 200KV).
3 RESULTS AND DISCUSSION
A. Billk GaN Growth
Colorless vvurtzitic GaN crystals < 3 mm in size were achieved by sublimation of the pressed GaN pellets in a stream of ammonia. Figure 1 shows an optical mi-crograph of =1 mm long, well faceted, transparent GaN crystals with low aspect ratios which are in contrast to the commonly observed13 needle-shaped crystals grown via vapor phase reaction. The GaN crystals primarily grew by spontaneous nucleation on the BN seed holders. The use of temperatures above 1200°C resulted in the rapid conversion of the GaN source into Ga metal. Crys-tals were grown at source temperatures in the range of 1100-1200°C; however, the majority of experiments were conducted at 1200°C to achieve higher grovvth rates. The temperature of the BN surface where crystals nucleated was =1000°C. The grovvth time and pressure for the growth of the crystals shown in Figure 1 vvere 2.5 hrs and 760 Torr, respectively.
The direction of fastest grovvth and thus the crystal shape vvere observed to change vvith the changing Ga/NH? flux ratio and the growth temperature. These ob-servations are in contrast to ali previous reports which
Figure 1: GaN crystals grown by sublimation/recondensation on BN. Lines are spaced 1 mm
Slika 1: GaN kristali zrastli s sublimacijo/rekombinacijo na BN. Črte so oddaljene 1 mm
Figure 2: Secondary electron microscopy of a GaN crystal shovving well developed {1010} and {0001} crystallographic facets Slika 2: SEM posnetek GaN kristalov, ki kaže dobro razvite {1010} in {0001} kristalne ploskve
indicate that grovvth of bulk GaN from the vapor phase results primarily in long needles.
Figures 2 and 3 show SEM images of the same crys-tal. This crystal grew from a single isolated nucleation site and developed into a well faceted hexagonal shape terminated by flat {1010} and {0001} planeš. Figure 3 shows a higher magnification image in which the (0001) and (1010) facets are observed.
A spectral mass scan via SIMS indicated that ali im-purities with the exception of oxygen vvere at back-ground levels. Quantitative analysis revealed an oxygen concentration of 3x1018 atoms/cm3 vvhich is similar to that present in high quality GaN thin films.
A representative room temperature PL spectrum for bulk GaN taken at 300K is shovvn in Figure 4. Strong near band edge (bound exciton) emission vvith a peak po-sition at 365.0 nm (3.4 eV) and a FWHM of 9.0 nm (83 meV) was observed. The visible portion of the PL
Figure 3: Higher magnification of the crystal in Figure 2 Slika 3: Kristal na sliki 2 pri večji povečavi
		C IT U <1 1	-
(b) . (a)	x 500 ___	J	L :
	.... i .... i .. .		. i . . . .
i .... i .... i . i i ■ » - ■ ■ ■ - - • • 650 600 550 500 450 400 350 300
VVavelength [nm]
Figure 4: Photoluminescence spectrum of GaN taken at 300 K Slika 4: Fotoluminiscentni spekter GaN pri 300 K
spectrum is expanded 500 times in the inset in Figure 4. No yellow emission usually attributed to deep level tran-sitions vvas detected on this scale or to the naked eye. A peak position of 359.0 nm vvith a FWHM of 54 meV vvas detected for a PL spectrum obtained at 77 K.
The allovved Raman modes of the vvurtzite structure are presented in Figure 5. The inset in this figure shovvs that the E2<2) mode is at 567 cm"1 and has a FWHM = 3.5 cm1. These values are indicative of a material of the highest quality reported to date14.
The results of optical absorption studies are shovvn in Figure 6. The absorption band edge is distinct but is not as sharply deftned as observed in thin epitaxial films. This is due to the absorption tail belovv the band edge. The absorption edge is expected to shift to vvavelengths above the actual band gap (360 nm for GaN) as the material thickness increases. For example, in the absorption spectrum of a high quality 50 |im thick GaN film the absorption edge is observed at 369 nm, and 75% transmis-sion is observed at 379 nm 15. Since the GaN crystal
Figure 5: Raman spectrum of GaN Slika 5: Raman spekter GaN
VVavelength (nm)
Figure 6: Optical absorption spectrum of GaN crystal Slika 6: Optični absorpcijski spekter GaN kristala
from vvhich the above spectrum taken vvas = 300 |im thick, the absorption edge position of = 370 nm and 75% transmission indicate these crystals to be of high optical quality.
B. AIN Bulk Growth
Grovvth in the ?100-2250°C range
Single crystal platelets of AIN having thicknesses to one millimeter and covering the vvhole seed crystal vvere obtained at a source temperature of 2150°C and a 4 mm source-to-seed separation. The grovvth rate vvas estimated to be 0.5 mm/hr. The results of XRD and Laue back re-flection studies confirmed the monocrystallinity.
At higher grovvth temperatures betvveen 2150-2250°C several = 2x2 mm individual hexagonal crystals vvere obtained on the seeds, since at these temperatures, severe degradation of the SiC substrates resulted in isolated sta-ble nucleation sites. These crystals and the aforemen-tioned platelets ranged in color from green to blue. The coloration strongly indicated the incorporation of impu-rities vvhich vvas confirmed via SIMS analysis to be C and Si from the SiC substrates.
Upon cooling, the AIN crystals frequently delami-nated and cracked. This vvas most probably due to the mismatch in the coefficients of thermal expansion betvveen the two materials; hovvever, intrinsic stress in the deposited material and/or the extension of pre-existing cracks at the edges of the SiC substrates may also have contributed to these phenomena. Unfortunately, the thermal expansion coefficients data for these materials are not available for the entire temperature range of the ex-periments. Since AIN boules vvill ultimately be grovvn on obtained AIN crystals, this cracking problem should not be a significant barrier to the attainment of much larger crystals.
Grovvth in the 1950-2050°C range
Grovvth of AIN in the temperature range of 1950 to 2050°C vvas conducted because complete structural and chemical stability of the SiC seeds and a significant re-
Figure 7: Optical micrograph of a single crystal of A1N grovvn at 1950°C and a 4 mm source-to-seed separation (Scale: mm) Slika 7: Optični posnetek kristala A1N, ki je zrastel pri 1950°C in razdalji 4 mm med izvorom in kaljo (merilo: mm)
duction in the deterioration of the SiC coated crucibles were attained. Crystals grown in this temperature range were always colorless regardless of the morphology or growth site and contained almost two orders of magni-tude less Si and C than the crystals deposited at >1950°C; whereas, the O levels were similar in ali crys-tals. Typical growth rates were reduced to 30-50 ftm/hr. Grovvth runs of 10-15 hrs yielded 0.3-0.5 mm thick crys-tals on 1 cm2 SiC substrates. A 0.4 mm thick transparent A1N platelet grown at 1975°C can be seen in Figure 7. Unfortunately cracking occurred in these crystals as well and presumably for the same reasons stated in the previ-ous subsection. Crystals grown in both temperature ranges had very smooth surfaces (~ RMS = 6A) as determined by atomic force microscopy (AFM).
Ali crystals showed strong and well defined single crystalline XRD patterns. Only the (002) reflection posi-tioned at 36° was observed in symmetric E-2E scans for a crystal grown at 1950°C. This suggests that the resid-ual stress level in these crystals was low. Bright field, plan vievv TEM micrographs and associated selected area diffraction (SAD) patterns taken along the [0001] direc-tion shovved uniform contrast density throughout the specimen and spot patterns vvithout streaks or arcs, re-spectively, indicative of single-crystalline material vvithout high angle boundaries, stacking faults, misoriented grains or twinned regions.
A Raman spectrum acquired using back scattering geometry from the (0001) face of an transparent A1N crystal grovvn at 1950°C is presented in Figure 8. The spectrum exhibits the allovved modes for this geometry, namely, Ai(LO) = 893 cm1, E2<» = 250 cm"1, E2(2) = 660 cm"1 with no detectable contribution from the forbidden modes. The results support the aforementioned crystal-lographic and microstructural results in that a vvell defined vvurtzite structure exists vvithout significant con-centrations of structural defects or internal stress vvhich
Figure 8: Raman spectrum of transparent bulk A1N grown at 1950°C Slika 8: Raman spekter prosojnega A1N, kije bil zraščen pri 1950°C
could relax the selection rules. The inset shovvs the high resolution spectrum of the E2(2) mode. The FWHM of this peak was 7.0±0.5 cm1; the same spectrum taken from a blue crystal grovvn at 2200°C had a FWHM value of 9.5+0.5 cm1. This marked difference complements the results of the SIMS analysis in vvhich the colored crys-tals contained tvvo orders of magnitude more Si.
C. Metalorganic Vapor Phase Epitaxial Growth of GaN
and AlxGai-xN films
Previous research in our laboratories has shovvn that thin films of GaN deposited directly on 6H-SiC(0001) substrates at high and low temperatures had columnar-like grains, faceted surfaces and high net carrier concen-trations (nD-nA > 1 x 1019 cm 3)16. In contrast, in the pre-sent research monocrystalline thin films of both the HT-A1N buffer layers and the subsequently grovvn GaN films deposited on similar SiC substrates have been deposited vvith no misorientation or lovv-angle grain boundaries, as determined by selective are diffraction and microstructural analysis via transmission electron microscopy. Similar results have been achieved for AlxGai-xN vvithout the use of an A1N buffer layer. The stacking fault density vvas also very lovv. The dislocation density of the AlxGai_xN films at the SiC interface ap-peared similar to GaN films deposited on high temperature (HT) buffer layers8 9. The dislocation densities of the GaN and AUGa^N films decreased rapidly as a function of thickness; only threading dislocations vvhich result from misfit dislocations at the interface persisted through the film.
The surfaces of the GaN and AlxGai.xN films exhib-ited a slightly mottled appearance as a result of the step and terrace features on the grovvth surface of the 6H-SiC(0001) substrates. Random pinholes, caused by in-compiete coalescence of the tvvo dimensional islands vvhich occurred as an intermediate grovvth stage betvveen the initial nucleation and the final layer-by-layer grovvth stage representative of the majority of the film, vvere also observed. An increasing number of pinholes appeared on
the surface of AlxGai.xN compositions where x > 0.5. The pinhole density was decreased by increasing the growth temperature to enhanced surface mobility of the adatoms. The DCXRC measurements taken on GaN and AlxGa,.xN films revealed the FWHM of the (0002) re-flections to be as low as 58 and 186 arcsec, respectively.
The low-temperature (8K) PL spectra of the undoped GaN films revealed an intense near band-edge emission at 3.466 eV, vvhich has been attributed to an exciton bound to a neutral donor'718. The FWHM value of this peak vvas 4 me V. Also, a less intense peak vvas observed at higher energies (3.472 eV) vvhich is attributed to free excitonic recombination. The lovv-temperature (4.2K) CL spectra of the undoped AlxGai.xN films for compositions in the range of 0.05 < x < 0.96 revealed an intense near band-edge emission vvhich has been attributed to an exciton bound to a neutral donor (I2-line emission)17 '8. Broadening of this emission is attributed to both exciton scattering in the alloys as vvell as small variations in al-loy composition in the film. The lovvest FWHM value observed in the AlxGai_xN alloys vvas 31 me V. Strong de-fect peaks, previously ascribed to donor-acceptor pair recombination19, vvere observed at midgap energies. The broad peak centered at 545 nm (2.2 eV) for GaN, com-monly associated20 vvith deep-levels (DL) in the bandgap, vvas also observed; hovvever, these emissions shifted sublinearly vvith changing composition. The na-ture of this behavior is under investigation.
The compositions of the AlxGai.xN films vvere deter-mined using EDX, AES and RBS. Standards of A1N and GaN grovvn in the same reactor under similar conditions vvere used for the EDX and AES analyses. After care-fully consideration of the errors (±2 at.%) involved vvith each technique, compositions vvere assigned to each film. The data from EDX and AES measurements shovved excellent agreement. The RBS data did not agree as vvell vvith the other tvvo techniques due to small com-positional variations through the thickness of the film. Simulation of the composition determined by RBS vvas conducted only on the surface composition.
These compositions vvere compared vvith their re-spective CL emission peaks and bandgap as determined by SE. Using a parabolic model, the functional relation-ships betvveen I2-line emission energy of the CL and the Al mole fraction for 0 < x < 0.96 is shovvn in Figure 9 and expressed analytically as
E,2(x) = 3.47 + 0.64x + 1.78x2	(1)
Clearly, this shovvs a negative deviation from a linear fit. This is in general agreement vvith earlier research over a smaller range of x by other investigators1617.
D. Selective Growth and Lateral Growth
The prismatic morphology of the GaN and Alo.2Gao.8N stripes deposited vvithin the various vvindovv vvidths of the Si02 masks vvas observed using scanning electron microscopy. Micrographs of these results are
Figure 9: Low-temperature (4.2K) CL emissions of AlxGai_xN films as a function of aluminum mole fraction
Slika 9: Nizko temperaturni (4.2K) CL emisije AlxGai_xN filmov v odvisnosti od molarnega deleža A1N
shovvn in Figure 10. Both materials exhibited (1100) side facets and ridge lines (i.e., no truncation) vvhen deposited on the 3 |4m-wide Si02 vvindovvs. Truncated prismatic grovvth vvith (0001) top facets and (llOl) side facets vvas observed on stripe patterns vvith vvidths > 5 |im. Polycrystalline islands of AlxGai.xN nucleated on the Si02 mask because of the chemical interaction betvveen Al and Si02 2n. There vvas no significant difference in the final grovvth morphologies betvveen the GaN and the Alo.2Gao.8N patterns, except for a slight roughening of the (1101) facets of the latter. This roughening is believed due to the changes in the gas flovv dynamics caused by the formation of the polycrystalIine islands. No exces-sive grovvth vvas observed along the top edges of the truncated stripes, as shovvn in Figure 11; the (0001) top facets vvere very smooth and flat regardless of the vvidth of these stripes. This suggests that the lovv grovvth pres-sure reduced the lateral vapor phase diffusion of the re-active species over the mask to the vvindovv region rela-tive to that observed in related research conducted at one atmosphere21. The large ratio (=0.5) of the vvindovv-to-mask area may have also contributed to the perfection of the stripes, since lovver values of this ratio vvere observed22 to induce marked grovvth and rounding of the top edges at one atmosphere in GaAs, reportedly as a re-sult of inereased lateral vapor diffusion.
An inerease in the flovv rate of TEG resulted in a de-crease in the area of the (0001) top facets and the deveiopment of (1101) side facets. This behavior supports the model23 that the resulting morphology of the selectively grovvn GaN depends on the balance betvveen the incom-ing vapor flux on the (0001) top facets and the rate of diffusion on the (0001) surface to the (1101) side facets. Inereased exposure of the former to the TEG via an inerease in the flovv rate causes the grovvth rate of these facets to become faster than the latter.
The optimum conditions for the selective grovvth of the GaN pyramids on the circular patterns vvere based on the selective grovvth conditions on the stripe patterns. Each pyramid contained six (1101) side facets. The grovvth rate of these pyramids vvas strongly dependent
Figure 10: Seeondary electron microscopy micrographs of GaN(left) and Alo.2Gao.8N(right) stripes selectively grown at 1050°C and having initial SiC>2 widths of (a and b) 3 pm and (c and d) 5 pm
Slika 10: SEM posnetki GaN (levo) in Alo.2Gay.8N (desno) trakov selektivno zrastli pri 1050°C z začetno širino SiC>2 (a in b) 3 pm in (c in d) 5 pm
Figure 11: Secondary electron microscopy micrographs of 10 pm wide GaN and Alo.2Gao.8N stripes selectively grown at 1050°C with 10 pm wide Si02 windows
Slika 11: SEM posnetki 10 pm širokih GaN in Al0.2Ga0.sN lamel selektivno zraščeni pri 1050°C z 10 pm širokimi SiCh okni
Figure 12: (a) SEM micrograph of a Si-doped GaN hexagonaI pyramid array grown at 1000°C. (b) High magnification SEM image of the apex of a hexagonal pyramid having a tip radius of 100 nm Slika 12: SEM posnetek heksagonalne združbe piramid GaN dopiranega s Si zrastlih pri 1000°C. (b) SEM posnetek vrha heksagonalne piramide z radijem konice 100 nm
turn-on voltage corresponds to a turn-on field of 25 V/pm. Using the same system, a polycrystalline p-type diamond film (p = 2.5E17cm~3) grown on Si(100) exhib-ited a turn-on field intensity of 27 V/pm. The Fowler-Nordheim (F-N) plot obtained from the I-V data was lin-ear and, therefore, indicates that the emission occurred via electron tunneling through the GaN.
The morphologies of the GaN layers selectively grown on the stripe openings were a strong function of the growth temperature, the flow rates of TEG and the stripe orientation. Continuous 2 pm thick GaN layers were obtained using 3 pm wide stripe openings spaced 7 pm apart and oriented along <1100> (Figure 14 (a)). The growth parameters were 1100°C and a TEG flow rate of 26 mmol/min. Plan view SEM of the coalesced GaN layer revealed a microscopically flat and pit-free surface, as shown in Figure 14 (b). Atomic force micros-copy showed the surfaces of the laterally grown GaN layers to possess a terrace structure having an average step height of 0.32 nm. The average RMS roughness val-ues of the regrown and overgrown layers were 0.23 nm and 0.29 nm, respectively.
The cross-sectional TEM micrograph presented in Figure 15 shows a typical laterally overgrown GaN. Threading dislocations, originating from the GaN/AlN buffer layer interface, propagate to the top surface of the regrown GaN layer within the window regions of the mask. The dislocation density within these regions, cal-culated from plan view TEM micrographs is approxi-mately 109 cm 2. By contrast, there were no observable threading dislocations in the overgrown layer. Additional microstructural studies of the areas of lateral growth obtained using various growth conditions have shown that the overgrovvn GaN layers contain only a few dislocations.
upon the ratio of the window-to-mask area in the pat-terned region as well as the selective growth conditions. The average diagonal width of the pyramids was 7.7 pm using a ratio of 0.1. Hovvever, increasing the ratio to 0.23 resulted in an average diagonal width of 5.7 pm for the same growth conditions. These results indicate that the lateral diffusion of the reactive species from the mask to the window area is also an important factor for the suc-cessful fabrication of the GaN pyramids. As shown in Figure 12(a), the grovvth of an uniform array of Si-doped GaN pyramids in a 0.5x0.5 mm2 area was achieved. The high magnification SEM image shown in Figure 12(b) reveals that the tip radius of the pyramids was less than 100 nm.
The field emission current from the Si-doped GaN pyramid arrays was measured as a function of the anode voltage. The I-V curve in Figure 13 shovvs a turn-on voltage of =680V for a current of 10.8 nA at a distance of 27 pm betvveen the pyramid array and the anode. This
0.0E+00
400 500 600 700 800 900 1000
Figure 13: Emission current and anode voltage characteristics of Si-doped GaN hexagonal pyramid array shown in Figure 12 Slika 13: Karakteristike emisijskega toka in anodne napetosti heksagonalnih piramid s Si dopiranega GaN s slike 12
s.**« rt
Figure 14: (a) Cross-section and (b) surface SEM micrographs of coalesced GaN layers grown on 3 ^im wide and 7 |im spaced stripe openings, respectively, oriented along <ll00>
Slika 14: Posnetki (a) prereza in (b) površina kolesciranih slojev GaN. ki so zrastli na 3 um širokih 7 |im med seboj oddaljenih lamelnih odprtinah orientiranih vzdolž <ll00>
4 CONCLUSIONS
Growth of bulk, vvurtistic GaN crystals to 3 mm
length vvas achieved at a substrate temperature and pres-
sure of 1100°C and 760 Torr, respectively, via sublima-
tion of GaN pellets produced by uniaxial cold pressing of GaN povvder. The concentrations of H, C and Si vvere < 1016 atoms/cm3 and the concentration of O vvas = 3xl018 atoms/cm3. Strong, sharp near band edge (bound exciton) emission vvas observed in the PL spectra of these crystals. No yellow emission vvas observed. An op-tical absorption edge of = 370 nm and 75% transmission vvas determined. The Raman spectrum shovved narrovv and well positioned peaks.
Seeded bulk grovvth of single crystalline A1N (001) platelets vvas achieved on 6H-SiC (0001) substrates in the temperature range of 1950-2250°C at 500 torr of N2. Color variations vvere observed above 2150°C and linked to the incorporation of Si and C from the substrate and the grovvth crucible. The results of TEM and XRD analy-ses revealed low densities of line and planar defects and
Figure 15: Cross-section TEM micrograph of a section of a laterally overgrown GaN layer on an SiOj mask region Slika 15: TEM posnetek prereza lateralno zrastlega GaN sloja na področju Si02 maske
the absence of residual stress in the grovvn crystals. No misoriented grains or tvvinned regions vvere observed. Very smooth surfaces (RMS=6 A) vvere observed via AFM.
Monocrystalline GaN and AlxGa,.xN(0001) (0.05 < x < 0.96) thin films, void of oriented domain structures and associated lovv-angle grain boundaries, vvere ob-tained via MOVPE on a(6H)-SiC(0001) vvafers. A 1000A high temperature (HT) A1N buffer layer vvas em-ployed for the GaN deposition vvhile AlxGai_xN vvas deposited directly on 6H-SiC. Double-crystal XRC meas-urements shovved FWHM values as lovv as 58 and 186 are sec for the GaN(0002) and AlxGai_xN(0002) reflec-tions. Photoluminescence spectra of GaN shovved bound and free excitonic recombination. Spectra obtained via CL of AlxGai_xN shovved strong near band-edge emis-sions vvith FWHM values as lovv as 31 me V.
The selective grovvth of GaN and Al0.2Ga0.sN has been achieved on striped and circular patterned GaN/AlN/6H-SiC(0001) multilayer substrates. Prismatic morphology vvith vvell-defined (1101) side facets vvas observed on 3 |im vvide stripes for both materials. Trun-cated prismatic grovvth vvith smooth, flat (0001) top facets and (1101) side facets vvere obtained on stripe pat-terns vvith vvidths > 5 pm. Uniform hexagonal pyramid arrays of Si-doped GaN vvere successfully grovvn on circular patterns having diameters of 5 |im. Field emission measurements of these arrays shovved a turn-on field of 25 V/pm and an associated emission current of 10.8 nA at an anode-to-pyramid array distance of 27 Lateral grovvth and coalescence over the Si02 masks have been achieved using stripes oriented along <ll00>. A density of =109 cm 2 threading dislocations, originating from the
underlying GaN/AlN interface, were contained in the GaN grown in the window regions. The overgrowth re-gions contained a very low density of dislocations.
ACKNOWLEDGEMENTS
The authors express their appreciation to Cree Research, Inc. for the SiC wafers. This work vvas supported by the Office of Naval Research under research contracts N00014-96-1-0765 and N00014-92-J-1477. R. Daviš vvas supported in part by the Kobe Steel, Ltd. Professor-ship.
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