Substrates for gallium nitride epitaxy

Materials Science and Engineering R 37 (2002) 61–127

Substrates for gallium nitride epitaxy L. Liu, J.H. Edgar* Department of Chemical Engineering, Kansas State University, Durland Hall, Manhattan, KS 66506-5102, USA

Abstract In this review, the structural, mechanical, thermal, and chemical properties of substrates used for gallium nitride (GaN) epitaxy are compiled, and the properties of GaN films deposited on these substrates are reviewed. Among semiconductors, GaN is unique; most of its applications uses thin GaN films deposited on foreign substrates (materials other than GaN); that is, heteroepitaxial thin films. As a consequence of heteroepitaxy, the quality of the GaN films is very dependent on the properties of the substrate—both the inherent properties such as lattice constants and thermal expansion coefficients, and process induced properties such as surface roughness, step height and terrace width, and wetting behavior. The consequences of heteroepitaxy are discussed, including the crystallographic orientation and polarity, surface morphology, and inherent and thermally induced stress in the GaN films. Defects such as threading dislocations, inversion domains, and the unintentional incorporation of impurities into the epitaxial GaN layer resulting from heteroepitaxy are presented along with their effect on device processing and performance. A summary of the structure and lattice constants for many semiconductors, metals, metal nitrides, and oxides used or considered for GaN epitaxy is presented. The properties, synthesis, advantages and disadvantages of the six most commonly employed substrates (sapphire, 6H-SiC, Si, GaAs, LiGaO2, and AlN) are presented. Useful substrate properties such as lattice constants, defect densities, elastic moduli, thermal expansion coefficients, thermal conductivities, etching characteristics, and reactivities under deposition conditions are presented. Efforts to reduce the defect densities and to optimize the electrical and optical properties of the GaN epitaxial film by substrate etching, nitridation, and slight misorientation from the (0 0 0 1) crystal plane are reviewed. The requirements, the obstacles, and the results to date to produce zincblende GaN on 3C-SiC/Si(0 0 1) and GaAs are discussed. Tables summarizing measures of the GaN quality such as XRD rocking curve FWHM, photoluminescence peak position and FWHM, and electron mobilities for GaN epitaxial layers produced by MOCVD, MBE, and HVPE for each substrate are given. The initial results using GaN substrates, prepared as bulk crystals and as free-standing epitaxial films, are reviewed. Finally, the promise and the directions of research on new potential substrates, such as compliant and porous substrates are described. # 2002 Elsevier Science B.V. All rights reserved. Keywords: Substrates; Gallium nitride (GaN); Epitaxy

1. Introduction The development of GaN-based devices has challenged the conventional thinking on the material requirements for successful device fabrication in several ways. With the exceptions of silicon on sapphire and GaN, no semiconductor has been commercialized exclusively using heteroepitaxial materials; all other semiconductors have employed bulk substrates. Before pn junction GaN-based light emitting diodes (LEDs) on sapphire substrates were demonstrated, good luminescence efficiency, regardless of the semiconductor, was thought to require very low dislocation densities, less than 106 cm
2 [1]. Researchers were surprised to measure excellent *

Corresponding autor. Tel.: þ1-785-532-4322; fax: þ1-785-532-7372. E-mail address: [email protected] (J.H. Edgar). 0927-796X/02/$ – see front matter # 2002 Elsevier Science B.V. All rights reserved. PII: S 0 9 2 7 - 7 9 6 X ( 0 2 ) 0 0 0 0 8 - 6

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luminescence efficiency from GaN LEDs despite dislocation densities four orders of magnitude higher. Unlike the majority of semiconductors, in GaN dislocations did not seem to degrade its optical and electrical properties. That sapphire was, and remains, the most common choice for GaNbased LEDs is particularly surprising, given its properties are seemingly unsuitable based on the usual assumptions made in choosing a substrate for epitaxy; it has large lattice constant and thermal expansion coefficient mismatches with GaN. As devices other than LEDs have been developed using GaN, it has become apparent that the conventional wisdom on the requirements for device fabrication and epitaxial growth still has merit. It is clear, in agreement with the conventional wisdom, that high dislocation densities are detrimental to the performance of devices with more sophisticated structures such as laser diodes, or for devices requiring large areas or operating at greater power densities [2]. Furthermore, although sapphire allows extremely versatile processing, its inherent properties do set limits on the final quality of the epitaxial film. In this review, the properties and processing of the most studied substrates for GaN epitaxy over the past 5 years are described. The influence of the substrate on the crystal orientation, crystal perfection, polarity, defect density, purity, strain, and surface morphology of the GaN epitaxial layer are reviewed. Summaries of the properties of GaN epitaxial films on the most common substrates are given. Useful substrate physical properties such as lattice constants, defect densities, mechanical properties, thermal expansion coefficients, thermal conductivities, and reactivities under reactive deposition conditions are presented. This review updates and expands the scope of several excellent reviews covering substrates for GaN epitaxy previously published [3–8]. 1.1. Properties, structure, and assessing the quality of GaN The physical properties of GaN make it an attractive semiconductor for many electronic [9] and optoelectronic devices [10]. Its wide, direct energy band gap makes it suitable for short wavelength emitters (LEDs and diodes lasers) and detectors. The wide energy band gap and good thermal stability of GaN is also advantageous for high-temperature (HT) and high power electronics. Gallium nitride forms solid solutions with AlN and InN, making a wide range (1.9–6.2 eV) of energy band gaps possible. This ability to form alloys is essential for producing specific wavelengths for emitters, and for creating heterojunctions with potential barriers into the device structures. Heat dissipation in devices is facilitated by GaN’s high thermal conductivity compared to silicon and gallium arsenide. Both n- and p-type conductivities are possible in GaN. Because the group III nitrides have a noncentrosymmetric structure and significantly ionic chemical bonding, they are strongly piezoelectric, and also undergo spontaneous polarization [11,12]. These effects can be advantageously employed, to increase the sheet carrier concentration in heterostructure transistor [13]. A summary of the physical properties of gallium nitride is given in Table 1. The thermal expansion coefficients of semiconductors including GaN depend on the temperature, and here only % changes in lattice constants for certain temperature range are listed in this paper. The temperature-dependent data of thermal expansion coefficient for various substrates including GaN, AlN, 6H-SiC, 3C-SiC, GaAs, Al2O3, ZnO, and MgO can be found in [19]. Gallium nitride normally has a wurtzite structure, with the space group of P63mc (no. 186). The wurtzite structure consists of alternating biatomic close-packed (0 0 0 1) planes of Ga and N pairs stacked in an ABABAB sequence. Atoms in the first and third layers are directly aligned with each other. Fig. 1 displays the perspective views of wurtzite GaN along [0 0 0 1], ½1 1 2 0 and ½1 0 1 0 directions, where the large circles represent gallium atoms and the small circles nitrogen. The closepacked planes are the (0 0 0 1) planes. The group III nitrides lack an inversion plane perpendicular to

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Table 1 Properties of GaN Property

Value

Reference

Energy band gap (eV) (300 K)

3.44

[14]

Maximum electron mobility (cm2/V s) 300 K 77 K

1350 19200

[15]

Maximum hole mobility (300 K) (cm2/V s)

13

[16]

Controlled doping range (cm ) n-type p-type

1016 to 4  1020 1016 to 6  1018

[17]

Melting point (K)

>2573 (at 60 kbar)

[18]

Lattice constants (300 K) a (nm) c (nm)

0.318843 0.518524

[19]

Percentage change in lattice constants (300–1400 K) Thermal conductivity (300 K) (W/cm K) Heat capacity (300 K) (J/mol K) Modulus of elasticity (GPa) Hardness (nanoindentation, 300 K) (GPa) Hardness (Knoop, 300 K) (GPa) Yield strength (1000 K) (MPa)

Da/a0 0.5749, Dc/c0 0.5032 2.1 35.3 210  23 15.5  0.9 10.8 100

[19] [20] [21] [22] [23] [23] [24]


3

the c-axis, thus, crystals surfaces have either a group III element (Al, Ga, or In) polarity (designated (0 0 0 1) or (0 0 0 1)A) or a N-polarity (designated (0 0 0 1) or (0 0 0 1)B). An excellent review on crystal polarity is given by Hellman [25]. The zincblende structure (space group F43m) of GaN can be stabilized in epitaxial films. The stacking sequence for the (1 1 1) close-packed planes in this structure is ABCABC. Perspective views of the zincblende structure are shown in Fig. 2. Many different techniques are employed to assess the quality of GaN heteroepitaxial films. A summary of some of the most important techniques with average, typical values for a thin (<5 mm) undoped GaN film on a sapphire (0 0 0 1) substrate is given in Table 2. As described in the subsequent sections of this review, it is no exaggeration to claim that the properties and quality of GaN film is primarily determined by the substrate on which it is deposited.

Fig. 1. Perspective views of wurtzite GaN along various directions: (a) [0 0 0 1]; (b) [1 1 2 0]; (c) [1 0 1 0].

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Fig. 2. Perspective views of zincblende GaN along various direction: (a) [1 0 0] (1  1  1 unit); (b) [1 1 0] (2  2  2 units); (c) [1 1 1] (2  2  2 units).

1.2. Potential substrates and the consequences of heteroepitaxy Because bulk gallium nitride crystals are not commercially available, most researchers have relied on heteroepitaxy, which is crystal growth on substrates of another material, for device fabrication. Most often, the lattice constant mismatch has been the primary criteria for determining the suitability of a material as a substrate for GaN epitaxy. A wide variety of materials have been studied for GaN epitaxy including insulating metal oxides, metals, metal nitrides, and other semiconductors. A summary of crystal structure and lattice constants of some of the candidate substrate materials for GaN epitaxy is given in Table 3. In practice, properties other than the lattice constants including the material’s crystal structure, surface finish, composition, reactivity, chemical, thermal, and electrical properties, are also important in determining its suitability as a substrate, as these greatly influence the resulting properties of the epitaxial layer, sometimes in unexpected ways. The substrate employed determines the crystal orientation, polarity, polytype, the surface morphology, strain, and the defect concentration of the GaN film. Thus, the substrate properties may ultimately determine whether the device achieves its optimal performance. The influence of the substrate on the polarity and polarization of the group III nitride epitaxial layer is particularly important. The chemical reactivity and the conditions required for good quality epitaxy depend on the polarity of the crystal. In many cases, the substrate controls the crystal polarity and the magnitude and sign (tensile or compressive) of the strain incorporated into the epitaxial layers, and thereby the extent of the polarization effect. Considerable variations are possible using a variety of epitaxial growth techniques, as evidenced from the case of sapphire, where epitaxial GaN films of either polarity can be controllably produced (see the subsection on sapphire). Nevertheless, the choice of substrate does provide limits on what can be done in subsequent processing. Thus far, the vast majority of substrates studied produce [0 0 0 1] oriented GaN, as this orientation is generally the most favorable for growing smooth films. However, interest in GaN epitaxial layers with other orientations is increasing to eliminate the polarization effects. Such effects can be deleterious for some optoelectronic devices, causing red shifts in emission. In addition, piezoelectric effects in quantum wells can cause a spatial separation of electrons and holes, thereby decreasing the recombination efficiency [72]. Problems arising from heteroepitaxy are summarized in Table 4. The misfit and threading dislocation densities in broad area epitaxially deposited GaN on foreign substrates such as sapphire

Table 2 Assessment techniques of the quality of epitaxial GaN films and typical measured values for a thin (<5 mm thick) GaN film on sapphire, except as noteda What is measured

LT (2–4 K); photoluminescence

Energy band gap; optical transitions Identity of the impurities between band edges and impurities present; defect densities (peak positions); peak widths

Van der Pauw–Hall measurement

What is inferred

Electron mobility; charge carrier concentration (n); donor (ND) and acceptor (NA) concentrations, activation energies Transmission electron microscopy Dislocation type and density; crystal polarity

Prevalence of scattering centers including impurities, point and line defects

X-ray rocking curves

Peak position: strain in the film and substrate; peak widths

Raman spectroscopy

Peak position; peak width

Biaxial stress in the film and substrate; dislocation type and densities Stress/strain; film quality; film polytype/orientation; electron concentration and mobility

Selective etching

Density of pits or whiskers remaining after etching Secondary ion mass spectrometry Elements present

Atomic force microscopy

Surface morphology; density of pits

Defect generation mechanisms

Dislocation densities

Typical values for GaN Donor-bound exciton (D8X): 3.467–3.479 eV; free-exciton (FX):b A 3.4799 eV, B 3.486 eV, C 3.5025 eV [26]; line-widths: MOCVD 1–2 meV, MBE 2–5 meV, HVPE: 1.4 meV mn ¼ 430 cm2/(V s) (300 K); n ¼ 5  1017 cm
3; ND ¼ 3.0  1017 cm
3, NA ¼ 1.8  1017 cm
3 Threading dislocations, primarily pure-edge dislocations with (1/3)h1 1 2 0i Burgers vectors; total defect densities (cm
2): MOCVD, MBE 108 to 1010, HVPE 400 mm; thick films 107; high densities of stacking faults and twins [27] Compressive stresses up to 1 GPa; full width at half maximal: (0 0 0 2) 30–270 arcsec, (1 1 0 2) 400–740 arcsec [28,29] E2(high) ¼ 567 cm
1 (unstrained) with stress-induced shifts of 2.9 cm
1 GPa (for biaxial stress); E2(high) linewidth of 3 cm
1 for good quality materials (300 K); from Raman modes and Raman selection rules; from A1(LO) phonon–plasmon coupled frequencies: undoped GaN 734 cm
1; n ¼ 5  1017 cm
3:749 cm
1; n ¼ 1  1018 cm
3:764 cm
1 [30] Etch pit densities (EPD) of 3  108 for MOCVD film [31]; EPD of 107 to 108 for thick HVPE material [32] NO  1017 to 1019 cm
3; NSi  1016 to 1017 cm
3 [33]

Identity of elements controlling the electrical and optical properties of the material; mechanisms and sources for the impurities Dislocation densities of 0.5 to 1.0  109 cm
2 [34] Growth mode type, density, and distribution of dislocations terminating at the film surface

L. Liu, J.H. Edgar / Materials Science and Engineering R 37 (2002) 61–127

Technique

a

Properties vary considerably with film thickness and impurity and doping concentrations. Examples of these figures of quality for exceptionally good material are given in the specific sections on substrates. b Values from a 400 mm thick nominally strain-free GaN layer.

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Table 3 Substrate candidates for GaN epitaxy Material

Structure

Space group

Lattice constants (nm) a

Semiconductors w-GaN zb-GaN r-GaN w-AlN zb-AlN r-AlN ZnO b-SiC SiC SiC BP GaAs GaP Si Oxides and sulphides Al2O3 (sapphire) MgAl2O4 (spinel) MgO LiGaO2 g-LiAlO2 NdGaO3 ScAlMgO4 Ca8La2(PO4)6O2 MoS2 LaAlO3 (Mn,Zn)Fe2O4

References c 0.5185

a

0.49795

a

0.52065

[35,36]

1.0053 1.51173

a

Wurtzite Zincblende Rock salt Wurtzite Zincblende Rock salt Wurtzite 3C (ZB) 4H (W) 6H (W) Zincblende Zincblende Zincblende Diamond

P63mc F43m Fm3m P63mc F43m Fm3m P63mc F43m

Rhombohedral

R3c Fd3m Fm3m Pna21 P41212 Pna21 R3m R3mR P63mc R3c Fd3m

0.4765 0.8083 0.421 0.5402 0.5169 0.5428 0.3236 0.9446 0.5364 0.85

1.311

P63mc P63mc Fm3m P63mc Fm3m Fm3m Fm3m

0.318 0.318 0.45776 0.3309 0.4502 0.4389 0.4241

0.519 0.519

Rock salt Orthorhombic Orthorhombic Apatite Rhombohedral Spinel

Metals and metal nitrides Hf HCP Zr HCP ZrN Rock salt Sc HCP ScN Rock salt NbN Rock salt TiN Rock salt a

b

P63mc F43m F43m F43m Fd3m

0.31885 0.4511 0.422 0.31106 0.438 0.404 0.32496 0.43596 0.3073 0.30806 0.4538 0.56533 0.54309 0.54310

a a

[37,38] a

[39,40] a

1.2982

a

[41–46] [47,48] 0.6372 0.5498

0.5007 0.6267 0.771 2.515 0.6922

0.54

a

[49–52] [53–56] [57,58] [59] [60,61] [62] [63] [64,65] [65] [66] [67] [67,68] [69] [70,71]

Detailed summaries of GaN epitaxial growth on these substrates are given in Sections 2–9.

and silicon carbide are typically between 108 and 1010 cm
2, compared to densities of essentially zero for silicon homoepitaxy, or 102 to 104 cm
2 for gallium arsenide homoepitaxy [73]. Other crystalline defects commonly observed in GaN heteroepitaxial layers include inversion domain boundaries, and stacking faults [74,75]. Such defects create non-radiative recombination centers, introduce energy states into the band gap, and reduce minority carrier lifetimes. Impurities diffuse more rapidly along threading dislocations than in bulk material, causing a non-uniform impurity distribution, and degrading pn junction abruptness [76]. Because of the high piezoelectric constants of GaN, the local strain surrounding threading dislocations causes submicron scale variations in the electrical potential and electric field in comparison to the bulk material [77]. Such defects are typically non-uniformly distributed, thus the electrical and luminescence properties of the material and devices made from the material are non-uniform as well. Defects increase the device threshold voltage and reverse bias leakage currents, deplete sheet charge carriers concentrations in heterostructure field effect

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Table 4 Problems commonly encountered with heteroepitaxy Substrate property

Consequence

Lateral (a-lattice constant) mismatch

High misfit (primarily edge) dislocation densities causing: high device leakage currents; short minority carrier lifetimes; reduced thermal conductivity; rapid impurity diffusion pathways Antiphase boundaries, inversion domain boundaries Double positioning boundaries (stacking mismatch boundaries) Thermally induced stress in the film and substrate; crack formation in the film and substrate Poor heat dissipation Contamination of the film by elements from the substrate; electronic interface states created by dangling bonds; poor wetting of the substrate by the growing film Mixed polarity in the epitaxial film; inversion domains

Vertical (c-lattice constant) mismatch Surface steps in non-isomorphic substrates Coefficient of thermal expansion mismatch Low thermal conductivity Different chemical composition than the epitaxial film Non-polar surface

transistors, and reduce the charge carrier mobility and thermal conductivity. These detrimental effects will prevent more complex or large area (necessary for high power) GaN devices from achieving their optimal performance. Regardless of the choice of substrate, many of its shortcomings such as its crystal quality or poor bonding characteristics with GaN can be ameliorated through an appropriate surface preparation such as nitridation, deposition of a low-temperature (LT) AlN or GaN buffer layer, multiple intervening LT buffer layers [78], epitaxial lateral overgrowth [79], pendeoepitaxy [80], and other techniques [81–83]. Through a variety of such techniques, GaN layers with dislocation densities as low as 107 cm
2 have been produced. Despite this success, even lower defect densities are necessary for more sophisticated devices operating at more extreme conditions of temperature, voltages, and current densities. Thus, substrates capable of supporting better quality GaN epitaxial layers will still be needed for realizing the full potential of GaN-based devices. This has driven research on methods of producing bulk AlN and GaN crystals.

2. Sapphire Sapphire, single crystal aluminum oxide, was the original substrate used in Maruskas and Tietjen’s pioneering study of GaN epitaxy by hydride vapor phase epitaxy (HVPE) in 1969 [84], and it remains the most commonly employed substrate for GaN epitaxy. The large lattice constant (15%) mismatch of sapphire (Al2O3) with GaN leads to high dislocation density (1010 cm
2) in the GaN epitaxial film [73]. These high defect densities reduce the charge carrier mobility, reduce the minority carrier lifetime, and decrease the thermal conductivity, all of which degrade device performance. Sapphire’s coefficient of thermal expansion is greater than GaN, thus, producing biaxial compressive stress in the layer as it cooled from the deposition temperature. For thick films, the stress can cause both the film and the substrate to crack [85]. The thermal conductivity of sapphire is low (about 0.25 W/cm K at 100 8C), thus, it is relatively poor at dissipating heat compared to other substrate materials. The cleavage planes of epitaxial GaN layer do not parallel to those of sapphire, making laser facet formation difficult. Sapphire is electrically insulating, thus, all electrical contacts must be made to the front side of the device, reducing the area available for devices and complicating device fabrication. In addition, there is evidence that

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Fig. 3. The unit cell of sapphire: (a) rhombohedral unit cell; (b) hexagonal unit cell.

oxygen from the sapphire causes unintentional doping in the GaN layer, raising its background electron concentration [86]. 2.1. Structure and properties of sapphire Sapphire has the space group of R3c (no. 167) and is mainly composed of ionic bonds. The ˚ 3 and single crystal can be described by both rhombohedral unit cells with volume 84.929 A 3 ˚ [87], which is displayed in Fig. 3. In the rhombohedral hexagonal unit cell with volume 254.792 A unit cell, there are 4Al3þ ions and 6O2
ions, and 10 ions in total. In the hexagonal unit cell, there are 12Al3þ ions and 18O2
ions, and 30 ions in total. The unit cell described in terms of hexagonal Miller-Bravais indices consists of six close packed (0 0 0 1) planes of O2
ions, sandwiching 12 planes of Al3þ ions that occupy two thirds of the available octahedral voids created by the O2
ions. The unreconstructed basal c-plane perspective views for both unit cells are shown in Fig. 4, where the polyhedra are the cell boxes. The faceting of sapphire crystal is shown in Fig. 5. All common surfaces employed for GaN epitaxy including the (0 0 0 1) and (1 1 0 0) are non-polar. Thus, the polarity

Fig. 4. Perspective views in (2  2  1) unit cells: (a) along the [0 0 0 1] direction in a rhombohedral unit cell; (b) along the (0 0 0 1) direction in hexagonal unit cell.

L. Liu, J.H. Edgar / Materials Science and Engineering R 37 (2002) 61–127

Fig. 5. Common facets of sapphire crystals: (a) view down c-axis; (b) surface planes (after [88]).

of the GaN epitaxial film is primarily controlled by the process conditions to which it is subjected. Table 5 shows the physical, chemical, thermal, mechanical and optical properties of sapphire. 2.2. Bulk crystal growth of sapphire Single crystal sapphire is manufactured by several methods, including the Verneuili technique (flame fusion growth), edge-defined film fed growth (EFG), Czochralski crystal growth, heat exchanger method (HEM), and gradient solidification method (GSM), also known as the temperature gradient technique (TGT). As a consequence of relatively HT gradients (>20 8C/cm) and fast growth rates, the Verneuli technique typically does not produce high quality crystals, hence it is not used to manufacture substrates for GaN. In the Czochralski technique, the crystal is pulled from a molten charge of Al2O3. A seed crystal is contacted with the surface of the melt, and is slowly withdrawn, causing crystallization to occur on the seed with its same crystal orientation [89,91]. Sapphire crystals of high crystalline perfection up to 15 cm in diameter and 25 cm in length can be grown along the c-axis and along the direction 608 to the c-axis. In the HEM, the Al2O3 charge is melted, and crystallization is initiated at the bottom center by cooling the center of the crucible with a flow of helium [92,93]. Crystallization proceeds as the solid/ liquid interface expands from the bottom to the top of the crystal. For the HEM, the temperature gradient is directed from the top to the bottom, resulting in less thermal convection. No moving parts are required in HEM growth of sapphire; thus, mechanical stresses are lower. The GSM technique is similar. Sapphire domes with different orientation are grown in double-walled molybdenum crucible, located at carefully designed thermal gradient. The growth is carried out via temperature lowering [94]. In EFG, the crystal is withdrawn from the melt as a ribbon with the cross-section controlled by a die [95]. This technique is ideal for producing thin rectangular cross-section, near-net sapphire crystals. Typically, single crystal sapphire contains several types of crystal defects including point defects, dislocations, and mosaic blocks. The most common dislocations are edge or mixed dislocations with Burgers vectors in the h11–20i direction, a consequence of the h11–20i (0 0 0 1) slip system [96]. Fig. 6 shows X-ray topography images for sapphire of the (0 0 0 1 2) surface (a) and in (0 3 3 0) transmission (b and c) from the same crystal produced by the HEM. A high-density

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Table 5 Properties of Sapphire from [89] except where noted Property

Value

Range

Lattice constant (nm) Melting point (8C) Density (g/cm3)

a ¼ 0.4765, c ¼ 1.2982 2030 3.98

20 8C


1


6

20 8C

Thermal expansion coefficients (K )

6.66  10 || c-axis 9.03  10
6 || c-axis 5.0  10
6 ? c-axis

20–50 8C 20–1000 8C 20 to 1000 8C

Percentages change in lattice constants

a/a0 ¼ 0.83, c/c0 ¼ 0.892

Heating from 293 to 1300 K

Thermal conductivity (W/cm K)

0.23 || c-axis 0.25 || a-axis

296 K 299 K

Heat capacity (J/K mol)

77.9

298 K

Young’s modulus (GPa)

452–460 in [0 0 0 1] direction 352–484 in the [1 1 2 0] direction

Tensile strength (MPa) Poisson’s ratio

190 0.25–0.30

300 K 300 K

Hardness Knoop Nanoindentation (GPa)

23.9  2.0 [90]

300 K

Dielectric constants

8.6 || c-axis 10.55 || a-axis

102 to 108 Hz 102 to 108 Hz

Refractive index

1.77 at the rate of l ¼ 577 nm 1.73 at the rate of l ¼ 2.33 mm

Energy band gap (eV) Resistivity (O cm)

8.1–8.6 >1011

300 K, experimental value 300 K

array of connected dislocations (indicated as D) is readily apparent. In the transmission images, the gray outside ring is due to absorption in the plastic used to mount the sample. As a consequence of lower temperature gradients, thermal induced stress, and mechanical stress, sapphire produced by the HEM is generally a higher quality than that produced by the Czochralski technique. Dislocation densities, as determined by etch pit densities [97] and X-ray topography [96], are approximately 1  103 cm
2 for sapphire produced by HEM compared to >104 cm
2 for Czochralski material.

Fig. 6. X-ray topography images for sapphire: (a) the (0 0 0 1 2) surface; (b and c) in (0 3 3 0) transmission.

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Table 6 Crystallographic relationship between GaN film and sapphire substrates Crystal plane

Miller indices (h j k l) or (h k l)

GaN plane || sapphire plane

c a

(0 0 0 1) or (0 0 1) (0 0 0 1) (1 1 2 0) or (1 1 0) (0 0 0 1); or (1 0 1 0)

r

(1 0 1 2) or (1 0 2)

(1 1 2 0), (1 2 1 6)

m

(1 0 1 0) or (1 0 0)

(1 0 1 3) or (1 2 1 2)

GaN direction || sapphire direction

Example references

[1 2 1 0] || [1 1 0 0] [1 1 2 0] || [1 1 0 0], [1 1 2 0] || [0 0 0 3]; [0 0 0 1] ^ [0 0 0 1] ¼ 308430 [0 0 0 1] || [1 1 0 1], [0 0 0 1] || [1 0 1 1], [1 1 0 0] || [1 1 2 0] [1 2 1 0] || [0 0 0 1], [1 0 1 0] || [1 2 1 0]

[99,100] [100–103] [100,104] [105]

To date, there have been no reports indicating that crystal defects present in the sapphire substrate are transferred into the GaN epitaxial film during growth. Such observations are probably obscured by the very high density of defects typically present at the interface between the GaN film and the sapphire substrate. The number of defects introduced into the GaN film due to similar defects in the sapphire is negligible compared to the high misfit dislocation density. 2.3. GaN film/sapphire substrate crystallographic relationships Epitaxial films have been grown on substrates of several orientations, and the resulting orientations are presented in Table 6. By far, the most common orientation of sapphire used for GaN is the c-plane sapphire. GaN epitaxy on c-plane sapphire results in c-plane oriented films, but with a 308 rotation of the in-plane GaN crystal directions with respect to the same directions in the sapphire. The 308 rotation of the (0 0 0 1) nitride plane with respect to the sapphire (0 0 0 1) occurs to reduce the lattice constant mismatch; the mismatch would be 30% without this rotation. With a proper nitridation and an optimized LT buffer layer of either AlN or GaN, very smooth GaN films can be obtained. The c-plane sapphire is very difficult to cleave, especially with a smooth cleave plane surface on the edge of the GaN layer. This is made worse by the 308 rotation of the AlN and GaN c-plane relative to the sapphire c-plane, such that the natural cleavage planes of the nitride and sapphire are not aligned. Stocker et al. [106] reported that c-plane sapphire cleaves along both the a- and m-planes, with roughly parallel cleaving of the m- and a-planes in the GaN film respectively. Large steps (90 nm) were produced on the surface of the m-plane of the GaN which were unsuitable for laser diode fabrication, hence, cleaves along the a-plane of sapphire unsuitable. Much smaller striations occurred on the a-plane of GaN, making it a more promising for fabricating lasers. Cleaving c-plane sapphire along its m-plane to expose the a-plane in the GaN remains an area of active research. GaN films grown on a-plane sapphire are also oriented in the [0 0 0 1] direction, and have the advantage that they are easily cleaved, along r-plane. This is potentially advantageous for forming edge-emitting lasers. Although the lattice mismatch for GaN films on a-plane sapphire is less (2%) than on c-plane sapphire (13.9%), no significant differences in GaN film quality have been reported between these two substrates. Nakamura et al. [107] has reported a pulsed MQW laser with cleaved facet mirrors fabricated in GaN on a-plane sapphire, which might be a promising for production of nitride-based lasers on a large scale. The r-plane of sapphire is unusual, as it is one of only a few substrates capable of producing GaN or AlN films with orientations other than the c-axis normal to the substrate. For either GaN or AlN epitaxy on r-plane (1 1 0 2) sapphire, the a-axis [1 1 2 0] orientation is produced. This is the

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preferred orientation for AlN-based surface acoustic wave devices, as it produces the electromechanical coupling coefficient. For on-axis sapphire substrates with this orientation, the AlN films are universally rough, forming a peak and valley morphology, bounded by (0 1 1 0) and (0 1 1 2) facets. The roughness generally increases and becomes more severe as the film thickness increases. By using r-plane sapphire substrates misoriented with a 48 tilt toward the [1–101] negative a-direction, Shibata et al. [108] improved the quality of AlN films by suppressing the formation of inverted twins. From the viewpoints of lattice mismatch and crystal symmetry, (1 0 1 0) sapphire (m-plane) seems the most suitable for GaN growth. However, the c-axis of GaN film grown on a (1 0 1 0) sapphire substrate has a non-zero inclination with respect to the c-axis of the sapphire substrate. Thus, twins may be generated, and the structure and surface morphology of GaN films becomes poor quality. This is a major disadvantage of a (1 0 1 0) plane compared to a (0 0 0 1) plane. Tripathy et al. [109] compared the quality of GaN layers with the same thickness (1.2 mm) grown on c-, a-, r-, and m-plane sapphire substrates by MBE. GaN layers were smooth and flat on cplane sapphire, somewhat rougher on a-plane, and rough on r-plane and m-plane sapphire. Only (0 0 0 2) and (0 0 0 4) XRD diffraction peaks were seen from GaN on c-plane and a-plane sapphire substrates due to the smooth surface morphology. The layers grown on r-plane and m-plane had mixed orientations, with both (0 0 0 2) and (1 1 2 0) diffraction peaks. From the smallest width of XRD rocking curve and narrow band PL spectrum, GaN on c-plane sapphire suggested the best crystalline quality. Although the PL intensities are comparable to that grown on c-plane sapphire, strong in-plane anisotropy of optical response was found in GaN grown on a-, r-, and m-plane sapphire substrates. Matsuoka and Higiwara [110] observed a significant improvement in the quality of GaN films deposited on m-plane sapphire substrates misoriented by 15–208. This misorientation suppressed twinning and enabled the growth of smooth, single crystal GaN films. Although the dislocation density was still 50% higher than on c-plane sapphire, this is still a notable improvement over previous efforts [110]. 2.4. Procedures for GaN epitaxy on sapphire substrates One of the reasons sapphire has been so successful as a substrate for GaN epitaxy is simply that much more research has gone into developing procedures for producing good quality films on it compared to other substrates. The large lattice constant mismatch between GaN and sapphire causes the film to be completely relaxed (not strained) essentially from the beginning of growth. Consequently, the defect density at the film/substrate interface is very high. The detrimental effect of sapphire’s large lattice constant mismatch must be ameliorated by a sophisticated processing scheme. First, the substrate surface is treated to remove contaminants, remnant polishing damage, and to produce a step and terrace surface structure. Second, the substrate is nitridated to alter the wetting characteristics of the deposited layers. Third, a thin buffer layer of either GaN or AlN (usually 10–100 nm thick) is deposited at a LT and annealed to produce a surface ready for the final epitaxial growth. There can be many variations in process conditions of each of these steps, and the best procedure is continuing to be defined. Summaries of the research done on each of these steps follows. 2.4.1. Surface preparation Sapphire substrates as received from the manufacturer typically contain randomly distributed scratches due to polishing. Substrates typically have root-mean square (RMS) roughness values

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between 0.8 and 2.1 nm over 1 mm2 areas. Heffelfinger and Carter [111] reported that annealing under flowing air at 1400 8C for 5 min reduced the RMS roughness of (0 0 0 1) surfaces from 0.8 to 0.3 nm. Similarly, Cui et al. [112] reported the combined removal of polishing damage by wet etching and HT annealing in air produced atomically flat sapphire surfaces. GaN films grown on such substrates had reduced XRD FWHM (180 arcsec for (0 0 0 2) reflections and 300 arcsec for (1 1 0 4) reflections respectively) compared to films deposited on as-received substrates. Many wet and in situ methods of etching sapphire have been reported in the literature, including phosphoric acid (H3PO4) [113], sulfuric–phosphoric acid mixtures (H2SO4–H3PO4) [114], fluorinated and chloro-fluorinated hydrocarbons [115], tetrafluoro sulfur (SF4) and sulfur hexafluoride (SF6) [116], but such aggressive etches have rarely been used. Instead, the most common substrate preparation procedure prior to the MOCVD growth of GaN and AlN is to simply heat the sapphire under flowing hydrogen at temperatures between 1000 and 1100 8C. Extrapolating measurements taken by Marasina et al. [117], the etch rates of sapphire in flowing hydrogen reported below 1100 8C should be negligible. Sapphire etching by hydrogen does occur at this temperature if gallium is present, and can cause hexagonal pits [118]. Surprisingly, the crystal quality of subsequently deposited GaN films was insensitive to the presence of these pits. 2.4.2. Sapphire nitridation The nitridation of sapphire before the growth of a LT buffer AlN or GaN is an important step for reducing the defect density, enhancing the electron mobility, and reducing the yellow luminescence in subsequently deposited films. Sapphire is nitridated by exposure to nitrogen plasmas or thermally cracked ammonia in MBE systems or by exposure to ammonia/hydrogen gas mixtures in MOCVD systems. The consensus is that the benefits of the nitridation layer are not due to a gradual transition in the physical properties from one material type to another (as is done with buffer layers), but are instead due to a change in the surface energy in GaN or AlN buffer layer deposited at LT. A number of researchers have studied the nitridation mechanism for c-plane sapphire. Under the conditions of temperature used for MBE and MOCVD growth, AlOxN1
x should be unstable and nitridation of sapphire should instead form AlN [119]. An abrupt interface between the unreacted sapphire and the nitridated AlN film implied that the rate of nitridation reaction is diffusion limited [120]. The authors estimated the diffusion coefficient for volume diffusion of oxygen anions is 10–20 cm2/s for polycrystalline sapphire at 950 8C. Hashimoto et al. [120] modeled the sapphire nitridation process to estimate the nitrogen species diffusion coefficient of 4  10
18 cm2/s at 1000 8C, a very low value in comparison to the diffusion of elements in silicon at the same temperature, but characteristic of diffusion coefficients in refractory nitrides in general. From their model, they estimated a nitride layer thickness of 0.6 nm is optimal to obtain high quality GaN by MOCVD. Such an AlN layer promotes GaN nucleation and increases the wetting of the GaN overlayer from 550 to 820 8C in MBE growth [121]. Analyzing the nitridation layer produced under MOCVD conditions, Uchida et al. [122] concluded that nitridation occurs very rapidly for times less than 3 min, then much more slowly. For short nitridation times (<3 min) the surface is relatively smooth, but stress-induced protrusions develop for longer nitridation times. The density of these protrusions increases with time, making the surface progressively rougher. In their study [122], the layer produced was an amorphous Al–O–N mixture. Interestingly, this amorphous layer was not seen in subsequent TEM micrographs after the deposition of a 4 mm thick GaN layer, which they suggest is due to diffusion of the N and O atoms into the crystalline layer. The duration of the nitridation is an important factor as to whether it is effective in improving the GaN layer crystal quality, as usually too long of a nitridation time leads to a reduction in the

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Fig. 7. AFM images of the sapphire surface with different nitridation duration: (a) 1 min; (b) 2 min; and (c) 30 min. The image area and height ranges are 1 mm  2 mm and 2 nm, respectively (after [124]).

film’s quality. Nitridating sapphire at 1050 8C, Keller et al. [29] observed in MOCVD GaN layers a reduction in the electron mobility (from 592 to 149 cm2/V s) and an increase in defect-related yellow luminescence as the nitridation time was increased from 60 to 400 s. This changed the morphology of the subsequently deposited LT GaN buffer layer after its temperature was raised to 1080 8C, from rough, highly faceted, and of the zincblende structure, to smooth films of the wurtzite structure. Kim et al. [123] also confirmed that a short nitridation time of 30 s produced a better surface morphology and a higher electron mobility than that of 180 s in MOCVD growth of GaN on c-plane sapphire. Fig. 7 shows the AFM image of sapphire surface after nitridation at 1000 8C and 760 Torr in MOCVD growth with different duration [124]. On the other hand, the temperature of the nitridation determines the limiting thickness for the nitridated layer. Excessive nitridation beyond the limiting thickness caused strain relaxation by a reduction of lateral domain size and an increase in lateral mosaicity [125]. Paek et al. [126] found that the protrusion density on the nitridated sapphire surface was critically dependent on the nitridation temperature. By controlling the RF power and nitridation temperature even at LT such as 400 8C in remote plasma enhanced ultrahigh vacuum MOCVD, they found the best crystallinity of GaN could be grown at temperature as low as 450 8C. In their case, the film crystallinity was improved as the nitridation time increased, possibly due to the low nitridation temperature. Widmann et al. [127] confirmed that lowering the nitridation temperature to as low as 200 8C for 90 min drastically improved the GaN structural and optical properties in MBE growth. The low nitridation temperature produced an ordered array of pure edge dislocations at the interface between nitridated sapphire and the AlN buffer layer, while a high nitridation temperature produced a perturbed interface with occurrence of cubic crystallites in the AlN buffer. Fig. 8 shows the TEM images of the AlN buffer layers grown on sapphire nitridated with nitrogen plasma at different temperatures for 90 min. The chemical alternation of surfaces of sapphire substrates using particle beams has been studied as an alternative process to nitridation. Advantages of simplicity and room temperature operation were claimed over nitridation. The reactive ion (N2þ) beam has been used in the

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Fig. 8. TEM images of AlN buffer layers grown on nitridated c-plane sapphire with nitrogen plasma for 90 min at different temperature: (a) 750 8C; (b) 200 8C (after [127]).

pretreatment of sapphire substrates and a very thin amorphous-like disordered AlON layer was formed. The optical properties of GaN epilayers subsequently deposited by MOCVD were improved by this method [128,129]. 2.4.3. Buffer layer In addition to nitridation, a LT GaN or AlN buffer (usually about 500–550 8C for MOCVD, 400 8C for MBE) dramatically improves the surface morphology and crystalline quality of GaN sequentially deposited at HT (usually about 1100 8C for MOCVD, 700 8C for MBE) on sapphire. Annealing at a high growth temperature crystallizes the amorphous buffer layers into films with preferential orientation ([0 0 0 1] for c-sapphire) and reduces the density of low-angle grain boundaries. Variations in process conditions in the growth of the buffer layer on the quality of subsequent GaN layers have been explored, such as growth temperature [130], buffer layer thickness [130,131], thermal treatment [132], and so on. Due to the large lattice constant mismatch between the sapphire and the GaN epilayer, a high density of threading defects (about 1010 cm
2) propagates directly through the as-grown layer. Fig. 9 shows conventional TEM image illustrating the high density of threading defects within CVD-grown GaN/c-sapphire [133]. Qian et al. [134] proposed that the main sources of threading dislocations are the low angle grain boundaries, formed during coalescence of islands at the initial stages of GaN growth. Clearly, the non-isomorphic properties of sapphire will result in the threading defects in GaN film [135,136]. The generation of dislocations with screw components appears to be more complex [137]. Edge dislocations are highly immobile within GaN, and this suggests that a higher proportion of screw dislocations be implicated in the relaxation process, presuming they are surface nucleated [133]. 2.4.4. Stress in GaN epilayers Because of the large lattice constant mismatch between GaN and sapphire (0 0 0 1), the GaN film can not be strained to lattice match to the sapphire; instead the film deposits relaxed to its unconstrained values. Nevertheless, a large biaxial compressive stress (1 GPa) is generated in the

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Fig. 9. Conventional TEM image illustrating the high density of threading defects within CVD-grown GaN/c-sapphire (after [133]).

GaN epilayer grown on c-plane sapphire upon cooling from the deposition temperature since the thermal expansion coefficient of sapphire is greater than that of GaN. Consequently, the GaN epilayer has a smaller in-plane (a-plane) lattice constant and a larger out-of-plane (c-plane) lattice constant than that of bulk monocrystal [138]. This stress causes changes in the photoluminescence spectra, the Raman peak positions, and cracking of both the GaN film and the sapphire substrate. Investigating the cracking of thick GaN films grown on sapphire, Etzkorn and Clarke [85] found that cracking is motivated by tensile growth stresses once a critical thickness was reached, and continued film growth buried the cracks. Kisielowski [139] gave a detailed review on strain in GaN thin films and heterostructures. Implicitly, the stress and strain can be utilized to grow high-quality GaN thin films. The actual magnitude of the stress is dependent on the process conditions, film thickness, and the film doping. Hiramatsu et al. [140] measured the thermal strain in a series of GaN films varying in thickness from 0.6 to 1200 mm on Al2O3(0 0 0 1). The compressive strain in the films increases with film thickness up to approximately 4.0 mm; cracks form in the film and substrate at greater film thicknesses. The film is completely relaxed only when the thickness is greater than 100 mm. In situ measurements of the stress in pure GaN films during deposition by MOCVD by Hearne et al. [141] showed that the film actually deposits in tension, regardless of whether an AlN or GaN buffer layer is employed. This observation helps to explain why the measured final stress in the films at room temperature was consistently less (
0:66  0:1 GPa) than that calculated from thermal expansion data (
1:4  0:7 GPa) for films 1–3 mm thick. The stress in the GaN epitaxial film is reduced as the thickness of either an AlN [141,142] or GaN [143] buffer layer is increased. Kim et al. [144,145] demonstrated that a Ga-rich GaN buffer can reduce tensile stress during HT growth, and therefore increase electron Hall mobility and reduce threading dislocation density. Romano et al. [146] studied the effect of silicon doping on strain in the GaN film, and showed that the stain changes from compressive to tensile as the silicon concentration is increased above approximately 5  1018 cm
3. Relationships have been established between the stress and the optical properties of the GaN film. Both the Raman phonon peak positions and the photoluminescence band-edge emission shift to higher frequencies as the compressive stress in the GaN is increased [142]. Kozawa et al. [147] found shifts and broadening in the E2 phonon peak of GaN with the stress in Raman measurements. They also found frequency shift at first time with the relation Do ¼ 6:2s, where the biaxial stress s is expressed in GPa. Wang et al. [131] analyzed their GaN epilayer grown by MOCVD with PL. They proposed that the scattered values of the near band gap exciton energy in literature between various authors due to different strain introduced by the different growth temperature.

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2.4.5. Substrate misorientation The effect of substrate misorientation for both c-plane and a-plane sapphire on the GaN epitaxial layer has been investigated, motivated by the improvements in crystal quality and film properties seen in the homoepitaxy of other semiconductors such as GaAs, Si, and SiC, and for heteroepitaxy of GaAs on Si. Both positive and negative results have been reported. For c-plane sapphire, Grudowski et al. [148] observed an increase in electron mobility (from 250 to 400 cm2/V s as the misorientation was changed from 0 to 108 for unintentionally doped GaN. Both Grudowski et al. [149] and Trager-Cowan et al. [150] (with Si-doped films in this case) reported that the yellow photoluminescence peak at 2.2 eV decreased with misorientation angle, possibly due to changes in the extent of impurity incorporation and electrical activation. On the downside, the structural quality worsens with misorientation from the c-plane, as evident from broader FWHM of the GaN X-ray (0 0 0 2) rocking curves (from 8.6 to 20.4 arcmin), rougher surface morphologies, a higher density of hexagonal pits, and spatially non-uniform cathodoluminescence. Pe´ cz et al. [151] also reported that the FWHM of both symmetric (0 0 0 2) and asymmetric (1 1 2 4) X-ray rocking curves for 3.58 misoriented substrates was broader compared to exactly oriented substrates, but claimed the misoriented substrates had a wider range of temperature under which smooth films could be deposited. Parillaud et al. [152] deposited GaN by HVPE on c-sapphire with 2, 4 and 68 misorientation towards m- and a-directions. There was no significant improvement in the crystalline quality or surface morphology, only slight improvement of the FWHM of the X-ray rocking curve for samples on 28 toward a-plane misoriented sapphire (from 746 to 496 arcsec). Yuasa et al. [153] also found no change in the electroluminescence intensity and peak wavelength for 4 mm thick GaN deposited on 0.03, 0.09, 0.17 and 0.258 misoriented c-sapphire by MOCVD. For a-plane sapphire, Fatemi et al. [154] determined that the X-ray rocking curve FWHM from the (0 0 0 2) peak decreased by half (from 425 to 258 arcsec), and the electron mobility increased by 2 (from 250 to 552 cm2/V s) as the misorientation was increase from 0 to 1.58. They attributed this to a reduction in the total dislocation density by half due primarily to a reduction in the edge dislocation density. The screw dislocation density was much less affected by the misorientation. Similar to the results on c-plane substrates, the yellow photoluminescence intensity was also reduced with the substrate misorientation. Someya et al. [155] showed that the roughness of the GaN film is a very sensitive to small changes in the misorientation of a-plane sapphire. The roughness of GaN films on substrates misoriented 0.258 off the a-plane was less than half that of films deposited on exactly oriented substrates. A maximum in photoluminescence intensity and a minimum in the X-ray rocking curve FWHM for the (0 0 0 2) peak (269 arcsec) was seen between substrate misorientations of 0.25 and 0.58 [155]. Clearly, definitive conclusions are difficult to make in these studies since changes in the misorientation can not be made without unintentionally changing other factors such the original roughness of the substrate, the density of steps and the widths of terraces on the substrate, and the reactivity of sapphire during nitridation. A substrate preparation procedure optimized for an exactly oriented sapphire substrate may not be the best procedure for a misoriented substrate. The consequences of the misorientation have been reported, but the root cause of the observed behavior has generally only been speculated upon. 2.4.6. Polarity Control of the polarity of GaN films is critical in epitaxy, as controls the surface morphology [156,157], doping character [158,159], and most important of all, the direction of piezoelectric field. The latter is crucial to the device performance. Generally for GaN on c-plane sapphire, smooth high quality films grown by MOCVD and HVPE are Ga-polarity [160–162]. In contrast to most

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researchers, Sumiya et al. [163] was able to produce N-polarity films by MOCVD. The polarity of GaN films on c-plane sapphire grown by MBE can be controlled by the buffer layer employed. A HT (>770 8C) GaN buffer layer produces N-polarity films [17,160,164,165], while a HT (>900 8C) AlN buffer layer produces Ga-polarity films [160,165]. Kung et al. [166] showed both c- and r-plane sapphires are oxygen-terminated and the HCP structure of the O2
ions in the sapphire continues by the N atoms in the GaN, hence Ga-polarity is generated in epitaxy. Generally, Ga-polarity GaN is preferable for obtaining smooth GaN films hence abrupt interfaces and heterojunctions. Sapphire is a non-polar substrate. Inversion domain defects (antiphase boundaries) will be produced due to the step-related inhomogeneities of the GaN/sapphire interface [167]. These polarity-inverted domains might result in the suppressed luminescence characteristics in these inhomogeneous areas [168]. Potin et al. [169] found that the inversion boundaries are nanometric (5–20 nm) limited by {1 0 1 0} planes crossing the whole epitaxial layer. Therefore, in general the inversion domain has a hexagonal columnar structure surrounded by (1 0 1 0) planes extending through the whole layer since the termination of inversion domain introducing the domain boundaries in the (0 0 0 1) plane is energetically unfavorable [170].

3. SiC Silicon carbide (both the 4H- and 6H-polytype) has several advantages over sapphire for GaN epitaxy, including a smaller lattice constant mismatch (3.1%) for [0 0 0 1] oriented films, and a much higher thermal conductivity (3.8 W/cm K) [4]. Conductive substrates are available, making electrical contacts to the backside of the substrate possible, thereby simplifying the device structure compared to sapphire substrates. The crystal planes in epitaxial GaN parallel those of the SiC substrate, making facets formation by cleaving easier. It is available with both carbon and silicon polarities, potentially making control of the GaN film polarity easier. However, most studies have been on Si-polarity substrates. All studies cited in this section are for Si-polarity substrates except when noted otherwise. High gain heterojunction bipolar transistors taking advantage of the discontinuity created at the GaN/SiC interface are possible [171]. However, SiC does have its disadvantages. Gallium nitride epitaxy directly on SiC is problematic, due to poor wetting between these materials [172]. This can be remedied by using an AlN or AlxGa1
xN buffer layer, but such layers increase the resistance between the device and the substrate. Even though the lattice constant mismatch for SiC is smaller than that for sapphire, it is still sufficiently large to cause a large density of defects to form in the GaN layers. Preparing smooth silicon carbide surfaces is difficult, thus, its surface roughness as received from the vendors is an order of magnitude (1 nm RMS) higher than that for sapphire (0.1 nm RMS). This roughness and also remnant subsurface polishing damage are sources of defects in the GaN epitaxial layer. The screw dislocation density in SiC is 103 to 104 cm
2 [173], and these defects may propagate into the GaN epitaxial layer and degrade device performance. SiC’s thermal expansion coefficient is less than that of AlN or GaN, thus, the films are typically under biaxial tension at room temperature. Lastly, the cost of silicon carbide substrates is high, and currently single crystal SiC is produced by relatively few manufacturers. 3.1. Polytypes and properties of SiC SiC exists in more than 250 polytypes—one-dimensional variations of the stacking sequence of close packed biatomic planes [174]. The basic unit of structure for all polytypes is a covalently

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Fig. 10. The tetragonal bonding of a carbon atom with four nearest silicon neighbors. The distance a and C–Si bond are ˚ , respectively. The right side is the three-dimensional structure of 2H-SiC structure. The approximately 3.08 and 1.89 A vertical direction in this figure corresponds to the [0 0 0 1] direction in the crystal.

bonded tetrahedron of C atoms with a Si atom at its center or vice versa, i.e. either SiC4 or CSi4, as illustrated in Fig. 10. The three-dimensional periodical repetition (stacking) of this unit, satisfies four-fold coordination at any structural point, therefore it may be used to generate a crystalline structure. The resulting crystalline structure can exhibit different crystallographic symmetries depending upon the stacking sequence along the c-axis. The various polytypes are distinguished by the stacking order between succeeding double layers of carbon and silicon atoms shown in Fig. 11. By observing the SiC crystal from the side, the stacking sequence can be projected as in Fig. 12. The ˚ for all polytypes. distance a between neighboring silicon or carbon atoms is approximately 3.08 A The height of the unit cell, c, varies with the different polytypes hence the ratio c/a also differs from polytype to polytype. The measured c/a ratios are approximately 1.641, 3.271 2H-, pffiffiffiffiffiffiffiffiandp4.908 ffiffiffiffiffiffiffiffi for the pffiffiffiffiffiffiffi ffi 4H- and 6H-SiC polytypes, respectively, very close to the ideal values ( 8=3, 2 8=3 and 3 8=3) [175]. Each polytype has a unique set of electronic and optical properties. The bandgaps at liquid helium temperature of the different polytypes range between 2.39 (for 3C-SiC polytype) and 3.33 eV (for 2H-SiC polytype). The two most important polytypes as substrates for GaN epitaxy, 6H- and 4H-SiC, have bandgaps at liquid helium temperature of 3.02 and 3.27 eV, respectively. The hexagonal polytypes of SiC such as 4H-, 6H-SiC belongs to the same space group P63mc (no. 186) as wurtzite GaN. The most studied substrates for GaN epitaxy are the 3C-SiC/Si(1 0 0) and

Fig. 11. The stacking sequence of double layers of the three most common SiC polytypes (after [176]).

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Fig. 12. The [1 1 2 0] plane of the 6H-, 4H-, 3C-, 2H-SiC polytypes (after [176]).

6H-SiC, as these polytypes have been the most readily prepared or commercially available for the longest time. With 4H-SiC now commercially available, its use as substrate for GaN epitaxy will become more common. Tables 7 and 8 show the physical, chemical, thermal, mechanical, electrical and optical properties of SiC at room temperature. Table 7 Structural, mechanical, and thermal properties of SiC at room temperature Properties

Polytype

Value

Lattice constant (nm)

3C 2H 4H 6H

a a a a

Density (g/cm3)

3C 2H 6H

3.166 3.214 3.211

[178]

Melting point (8C) Heat capacity (J/g K)

3C 6H

2793 0.71

[179] [180]

Thermal conductivity (W/cm K)

3C 4H 6H

Linear thermal expansion coefficient (10
6 K
1)

3C 6H

3.9 4.46 a-axis 4.16 c-axis

[97]

Percentage change in lattice constants (300–1400 K)

6H 3C

Da/a0: 0.4781, Dc/c0: 0.4976 Da/a0: 0.5140

[19]

Young’s modulus (GPa) Poisson’s ratio Refractive index (ordinary ray)

3C Ceramic 3C 2H 4H 6H

440 0.183–0.192 2.6916 at the 2.6686 at the 2.6980 at the 2.6894 at the

[182] [183] [184]

¼ ¼ ¼ ¼

Reference

0.43596 0.30753, c ¼ 0.50480 0.30730, c ¼ 1.0053 0.30806, c ¼ 1.51173

[177]

[181]

rate rate rate rate

of of of of

l l l l

¼ ¼ ¼ ¼

498 nm 500 nm 498 nm 498 nm

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81

Table 8 Electrical and optical properties of SiC ([185] except as noted) Energy band gap (eV) (T < 5 K) Charge carrier mobilities (cm2/V s) (300 K) Electrons Holes High breakdown electric field (V/cm) High saturation drift velocity (cm/s) Static dielectric constant Electrical resistivity (undoped) (O cm)

3C-SiC

6H-SiC

4H-SiC

2.4

3.02

3.26

800 40

400 101

1000 115

2.12  106

2.4  106 2.0  107 9.66

2.2  106 2.0  107

9.72

102 to 103 [186]

3.2. Bulk crystal growth of SiC Currently, good quality and large size wafers of two polytypes, 4H- and 6H-SiC, are available commercially. Bulk SiC crystals are produced by sublimation in the modified Lely process developed by Tairov and Tsvetkov [187] in 1978. The modified Lely process employs a SiC seed crystal to control the polytype and the crystal orientation produced. Growth is performed at 2200 8C in an argon atmosphere at pressures between 20 and 500 Torr in a graphite, tantalum, or tantalum carbide crucible. SiC crystals grow by the vapor transport of Si, Si2C, and SiC2 driven by a temperature difference between a SiC source zone and a growth zone. 6H-SiC substrates made by this technique became commercially available in 1991. Now 4H- and 6H-SiC wafers up to 100 mm in diameter are available, on and off-axis (typically 3.58 for 6H-SiC and 88 for 4H-SiC) from the (0 0 0 1) plane, in either silicon and carbon polarities. The longstanding problem with bulk SiC has been high screw dislocation densities (see Fig. 13) [188]. Depending on the magnitude of its Burger’s vector, the core of a screw dislocation in SiC can be hollow or closed. Hollow core screw dislocations (also known as nanopipes or micropipes) occur when the Burger’s vector (b) is two or more times the c-lattice constant for 6H-SiC, or three times the c-lattice constant for 4H-SiC. These holes are oriented along the c-axis and can run through the

Fig. 13. SWBXT back-reflection images of closed-core (smaller white spots) and hollow-core (large white spots) screw dislocation in a (0 0 0 1)6H-SiC wafer. The faint lines connecting these screw dislocation images are basal planes (after [188]).

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Fig. 14. (a) Synchrotron transmission topography showing the double contrast images of micropipes in a longitudinally cut SiC wafer. (b) Simulated image of a right-handed micropipe in the transmission geometry (after [188]).

entire length of the crystal (see Fig. 14) [188]. As the micropipe density in ‘best wafers’ has been reduced to less than 1 cm
2 [189], SiC bulk crystal growth research has focused on the effects of closed-core screw dislocations, which are present in high densities, approximately 103 to 104 cm
2. The detrimental effects of micropipes and screw dislocations have been clearly demonstrated for SiC-based devices [173,190]. Presumably, these defects also compromise the performance of GaN devices on SiC substrates, but such verifying studies have not yet been reported. Both 4H- and 6H-SiC wafers are available in low resistivity n- and p-type forms. Recently, interest in semi-insulating SiC has developed, as it is the preferred substrate for both SiC- and AlGaN-based power microwave devices. Typical electrical properties of commercial SiC are summarized in Table 9. Similar values of these properties are available for both polytypes. 3.3. GaN film/SiC substrate crystallographic relationships Both AlN and GaN deposit on hexagonal SiC substrates in a simple epitaxial relationship, that is [0 0 0 2]III-N || [0 0 0 6]SiC and [1 1 2 0]III-N || [1 1 2 0]SiC. Using (1 1 0 0)6H-SiC substrates (not available commercially), Horino et al. [191] established the epitaxial relationship as (1 1 0 0)III-N || (1 1 0 0)SiC and [1 1 2 0]III-N || [1 1 2 0]SiC. 3.4. Procedures for GaN epitaxy on SiC substrates The many studies of GaN epitaxy on SiC substrates can be divided into three areas: substrate preparation, nucleation, and growth. Substrate preparation studies have focused on methods of producing cleaner, more ordered SiC surface. This is done to enhance nucleation, and to minimize Table 9 Conductivity types, impurities and resistivities of commercial 4H- and 6H-SiC wafers Conductivity type N P Semi-insulating

Impurity N Al

Carrier concentration range (cm
3) 15

19

10 to 10 1015 to 1019

Resistivity range (O cm) 0.01–0.10 1–10 105

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defects in the final epitaxial layer. Nucleation studies have considered the effects of temperature, substrate polarity, and the effects of process conditions (temperature, V/III ratio, etc.) on the growth mode and deposit morphology of the initially deposited AlN or GaN. Growth studies look at the cumulative effects of surface preparation and nucleation on the film quality and stress in the film. Studies prior to 1997 were summarized in [4]. The results from more recent studies are displayed in Table 10. Conclusions from these three types of studies are presented below. 3.4.1. Surface preparation Although the smaller lattice constant mismatch suggests SiC should produce better GaN films than sapphire, the best GaN films on both substrates are similar in quality. Early, it was recognized that one factor limiting the quality of GaN films on SiC substrates may be the poor surface finish (roughness) of SiC substrates as-received from vendors. Sapphire can be prepared with surfaces having one-tenth the roughness of SiC. Producing SiC substrates with comparable surface smoothness has proved difficult due to its extreme hardness and high chemical stability. Typically, as-received SiC substrates have surface scratches and subsurface crystal damage, remnants of the polishing procedure. An example of the surface scratches on a 6H-SiC substrate is shown in Fig. 15. In addition to the scratches, there is very little order to the surface structure of as-received substrates; the surface may terminate on any step position A, B, or C (see Figs. 11 and 12). Defects generated at the surface may propagate into the epitaxial structure and, hence, degrade the nitride film quality. To remove this damage from the 6H-SiC surface, and to produce a more uniform step and terrace structure, several methods have been examined including oxidation of the SiC to form SiO2 which is subsequently removed by hydrofluoric acid [204]; etching in HCl þ H2 [205–208]; reactive ion etching [209–211]; or most commonly, HT (>1500 8C) etching in hydrogen [195,212–216]. Hydrogen etching removes all scratches, and on on-axis 6H-SiC substrates produces an ordered surface of steps one full (1.5 nm) unit cell in height, which presumably all terminate on the same stacking position [217]. 3.4.2. Nucleation on 6H-SiC substrates The nucleation behavior of GaN and AlN are significantly different. Aluminum nitride nucleates at both HT and LT. In contrast, GaN only nucleates on 6H-SiC at LTs (<800 8C), in a random, three-dimensional island growth.

Fig. 15. AFM image of as-received on-axis 6H-SiC with RMS 0.85 nm.

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Method

Substrate (6H-SiC)

Buffer layer/temperature or nucleation/temperature (8C)

Thickness (nm)

Growth temperature (8C)

V/III

Thick (mm)/ rate (mm/h)

FWHM of DCX (arcsec)

Reference

MBE MBE MBE MBE MOCVD MOCVD MOCVD MOCVD MOCVD HVPE CBE

Si(0 0 0 1) Si(0 0 0 1) Si(0 0 0 1) 3.58 off 4H Si(0 0 0 1) 3.58 off (0 0 0 1) Si(0 0 0 1) Si(0 0 0 1) Si(0 0 0 1) Si(0 0 0 1) p-type Si(0 0 0 1)

GaN/660 AlN/660 AlN/550–600 None Al0.08Ga0.92N/525 None GaN/550 AlN/1100 GaN/980–1080 None AlN/900

10 N/A 26

710 720 770 650 1025 980 1100 1100 1080 950 850

0.8 N/A 1.0 N/A 10000 70 4000 4000 N/A N/A N/A

1/0.7 1/0.20 1.0 0.5–1 1.3 /1.0 1–2 3.0 0.6 0.5–1.5 /0.15

300 500 228 130 138 700 378 245 107 90–180 550

[192,193] [194] [195] [196] [197] [198,199] [200] [201] [172] [202] [203]

a

50 None 100 100 100 20

MBE, molecular beam epitaxy; MOCVD, low pressure chemical vapor deposition; CBE, chemical beam epitaxy; HVPE, hydride vapor phase epitaxy; FWHM, full width half maximum (for (0 0 0 2) or (0 0 0 4) peak o scan); N/A, not available.

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Table 10 Most recent works of GaN expitaxy on SiC since 1997a

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Fig. 16. (a) Schematic of GaN/SiC and (b) GaN/AlN/SiC growth. Note that the strain is fully relieved in (a), whereas it is only partially relaxed in (b) (after [194]).

A high deposition temperatures (>1100 8C), AlN nucleates in a two-dimensional manner, resulting in complete surface coverage after the deposition of very little material. The AlN is truly epitaxial, without the high density of defects at the film/substrate interface seen with AlN/sapphire. It helps reduce the direct mismatch of 3.4% between GaN and 6H-SiC, and promotes wetting of the substrate surface. Fig. 16 illustrates the effect of HT AlN buffer deposited at 660 8C for GaN growth by MBE at 720 8C [194]. The postgrowth thermal processing of thin (200 nm) AlN films grown by MBE at HT in the range of 1200–1400 8C has shown to markedly improve the quality [218]. This may be used to improve further the quality of HT AlN buffer layer. At low deposition temperatures, AlN deposited on SiC substrates has a poor crystallinity and a rough surface morphology. The interface between the SiC substrate and AlN buffer layer is usually atomically abrupt and follows SiC(0 0 0 6) basal planes for MOCVD growth, as shown in Fig. 17 [219]. Meanwhile, the initial buffer layer growth is three-dimensional, indicated by the existence of small crystallites. The diameter of the crystallites has a value comparable to their height, suggesting high nucleating rates and slow growth kinetics, which are different from those of films grown on sapphire substrate. An AlN buffer layer deposited at LT will form large grains during an annealing temperature ramp, improving the structural quality of the layer and hence improving the structural quality of any subsequently deposited layer. Nevertheless, the subsequently deposited GaN film quality is inferior compared to GaN films deposited on HT AlN buffer layers. This is in contract to the case of AlN buffer layer on sapphire, where LT buffer layer results higher GaN quality than that of HT buffer layer [220].

Fig. 17. Lattice image of the buffer layer region. The high defect density at the buffer layer resulted from threedimensional growth of the AlN layer (after [219]).

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As an alternative to pure AlN, AlxGa1
xN buffer layers have also been considered as the later has a smaller lattice mismatch to GaN and can be electrically conductive. The AlxGa1
xN buffer layer provides a mean of changing the relative lattice match between the substrate and upper GaN layer, it is another control on the GaN film quality. Lin and Cheng [197] found the GaN quality improved with an Al0.08Ga0.92N buffer layer grown by low-pressure MOCVD. The deposited GaN epitaxial layer had a mobility and carrier concentration of 612 cm2/V s and 1:3  1017 cm
3 (at 300 K), respectively, and the FWHM of the X-ray (0 0 0 2) rocking curve was 145 arcsec for 1.3 mm GaN. For heterojunction GaN/SiC devices, direct growth of GaN on SiC without an AlN buffer layer is necessary [193]. Direct nucleation of GaN has been applied for the growth on SiC substrate in MBE [192,193,213,221] and MOCVD [172,222,223]. Both HT and LT GaN nucleation accompanied with a sequential HT annealing have been examined. Ploog et al. [224] pointed out that a low nucleation temperature is required to achieve GaN wetting in MBE growth, although higher nucleation temperatures are expected to improve the structural perfection of the individual GaN nuclei. Due to the low thermal stability of GaN, the temperature for GaN nucleation is usually performed at the same or lower temperature as film growth. Although there is no direct comparison for the film qualities between HT nucleation and LT nucleation, comparable GaN quality has been achieved with deliberately designed growth process by both methods. However, the effect of buffer layer or nucleation on the GaN film quality is so complicated that the optimal growth or nucleation condition and annealing process may depend on every growth system such as III/V ratio, bulk film growth temperature and so on. 3.4.3. Growth of GaN and AlN on 6H-SiC substrates Torres et al. [212] showed that a HT (1600 8C) H2 etch of the 6H-SiC substrate produced a more ordered step and terrace structure and subsequently higher quality AlN epitaxial films in comparison to similar films deposited on as-received 6H-SiC substrates. Xie and coworkers [205,206] demonstrated that both the surface morphology and crystal quality of epitaxial GaN films were enhanced by first etching the 6H-SiC substrates. Smith et al. [213] also found that the threading dislocations originated from substrate/buffer and/or buffer/film interfaces were related to the smoothness of the substrate. So surface treatment of SiC is required to produce high quality GaN films. Fig. 18 shows different morphologies of GaN deposited on 6H-SiC(0 0 0 1) with different surface treatments, where sample (a) and (c) were deposited on 50 nm AlN buffer layer at the same temperature (800 8C) as that of GaN growth by MBE [214]. Torvik et al. [225] found if the native oxide is not completely removed from the SiC surface, holes will be generated in the deposited GaN. The critical step for substrate cleaning was the removal of any thermal oxide on the SiC prior to GaN growth. Meanwhile, nitridation can also improve the smoothness of SiC substrate due to a combination of nitrogen chemisorption and etching at 1050 8C in NH3 flow for 30 min [225]. The SIMS measurements showed no evidence of excess H, O, or N in the nitridated SiC. 3.4.4. Stress in GaN epilayers Due to the smaller thermal expansion coefficient for SiC than for the wurtzite GaN, the final GaN films on SiC typically are in stress, but both the magnitude and sign of the strain vary considerably with process conditions. This accounts for some of the scatter in the recent data on the optical properties of group III nitrides [226]. Typically, GaN epitaxial layers on SiC are highly strained, with the average value of biaxial stress ranging up to 1 GPa. The large lattice mismatch together with the considerable differences in the thermal expansion coefficient between GaN film

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Fig. 18. Different TEM morphologies of GaN deposited on (0 0 0 1)6H-SiC with different surface treatments: (a) chemical cleaning; (b) chemical cleaning þ hydrogen plasma treatment; (c) chemical cleaning þ hydrogen plasma treatment þ 1300 8C hydrogen heat treatment for 30 min (after [214]).

and SiC, lead to a large extended defects density of 109 to 1010 cm
2. Optical and Raman spectroscopy studies of MOCVD-grown and MBE-grown GaN on SiC confirm they are under tensile stress, while those on sapphire substrates are under compression [227,228]. This suggests that the mismatch of the thermal expansion coefficient is the dominant factor determining the stress in the GaN films on SiC substrate. The residual built-in tensional biaxial strain, introduced by the SiC substrate, causes a downshift of the excitonic band gap and reduced valence band splitting in comparison with bulk strain-free GaN samples [229]. The quantitative relationship between the band gap and strain in GaN were investigated by Raman spectroscopy and photoluminescence [230], and the biaxial stress of 1 GPa results in a shift of the excitonic photoluminescence lines of 20  3 meV. The compressive lattice mismatch stresses were assumed to change into tensile thermal mismatch stress after a few nanometers of growth, but actually the situation is more complicated [231]. Analysis of samples with 100 nm of AlN buffer layers grown at 1000 8C by MOCVD, revealed that the films were mostly compressive for thicknesses less than about 0.7 mm, were tensile up to about 2 mm, then abruptly become less tensile with stress values near 100 MPa thereafter. Apparently, the 6H-SiC substrate/HT AlN buffer layer combination enables a wide variety of options for tailoring stress states in GaN layers. This may be potentially useful in valence band engineering and device processing applications. Some researchers have investigated the various properties of GaN films on SiC, most of them focusing on optical properties [222,229,232]. Obviously, the built-in stress due to the mismatch of thermal expansion coefficients and lattice parameters will change the band gap energy, and hence the luminescent and other properties. Thus, different research groups have measured different values for material properties due to the difference in growth method and growth conditions such as temperature, layer thickness and layer sequence. The conflicting strain results for GaN/6H-SiC may be explained by the influence of the AlN nucleation layers on the growth mode and strain relief. Waltereit and Brandt [194] found that the strain state of the GaN layer was determined by its growth mode, which in turn was governed by the degree of wetting of the substrate rather than by lattice mismatch. They predicted that the tensile strain which favors the formation of cracks in the GaN layer can be avoided by growing GaN epitaxial layers on thin, coherently strained AlN nucleation layers.

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Fig. 19. (a) High-resolution electron micrograph of wurtzite GaN grown by plasma-enhanced MBE on 6H-SiC showing substrate steps and an associated SMB. (b) Atomic model of wurtzite GaN on 6H-SiC. Steps on SiC surface such as S1 are likely to create stacking mismatch boundaries whereas others such as S2 will not (after [235]).

3.4.5. Substrate misorientation Misorientation of the SiC from the (0 0 0 1) plane can enhance two dimensional growth of epitaxial GaN by reducing the distance which atoms must diffuse on the substrate surface to incorporate into the crystal [233]. Smith et al.’s work shows that the misorientation (miscut angle of 48) can relieve the stress at the interface between the SiC substrate and GaN films [213]. But many stacking mismatch boundaries (SMB, a special kind of double-positioning boundaries, DPB) are inevitable because of the steps on the substrate surface. SMBs would always be created for two- and four-bilayer steps, and with a 2/3 possibility for one-, three- and five-bilayer steps, but never for sixbilayer steps [234]. The SMBs are results of individual island growth and coalescence. SMBs must be considered as inevitable during growth of GaN on non-isomorphic substrates such as SiC [235]. But, it also provides a consequential result that the LT buffer layer and proper etching of substrate to get six-bilayer height, such as HT (1300–1500 8C) HCl etching [236], will reduce SMB density and hence ensure good film quality. Fig. 19 shows the SMB produced at GaN/SiC interface [235]. Davis et al. [237] investigated the structure and crystal quality of AlN layers grown on on-axis and vicinal 6H-SiC(0 0 0 1) by gas-source MBE. They reported that the on-axis layer has a smooth surface and excellent thickness uniformity indicating two-dimensional growth. In contrast, a slightly rougher surface caused by island-like areas was seen in the AlN film grown on the vicinal surface. This results from the mismatch of c-axis lattice parameters between SiC and AlN since larger density of AlN islands can coalesce on the on-axis substrate. 3.4.6. Polarity Polarity control could be a key advantage of SiC compared to sapphire if a single polarity is produced on a specific substrate. For sapphire without the proper nitridation and buffer layer, films of mixed polarity may be deposited. The electrical polarity of the GaN/SiC interface plays a central role in determining the quality of the GaN optoelectronic material. The polarity of the SiC substrate strongly influences the surface morphology and crystal quality of the epitaxial GaN films. Electronic structure calculations for the (0 0 0 1) interface indicates that the stronger bonds are at the Si/N interface and at the C/Ga interface [238] and therefore

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Fig. 20. (a) High-resolution electron micrograph from region of SiC/GaN interface. (b) Possible atomic arrangements for GaN/SiC interface with (0 0 0 1)-oriented substrate and epilayer, with allowance for some intermixing to maintain charge balance (after [241]).

Si-terminated SiC substrate should produce a Ga-terminated (0 0 0 1) GaN epilayer. Ren and Dow [239,240] analyzed the lattice-matching of SiC with GaN using a tight-binding model of the electronic structure. They argued that the GaN grown on a C-terminated SiC(0 0 0 1) substrate has a local microscopic lattice mismatch of 6%, while in contrast, GaN grown on the Si-terminated surface has local lattice mismatch <3% by theoretical calculation. They predicted that the tri-bonded N/Si interface of GaN/SiC should produce excellent GaN and AlN films. Stirman et al. [241] compared the high-solution electron microscope image and simulation image to examine the atomic arrangements in a GaN/SiC heterostructure concluding that the atomic arrangement at the GaN/SiC interface most likely consisted of N bonded with Si, but with some Ga bonded to C to maintain the charge balance. Fig. 20 shows the HREM of GaN/SiC and proposed possible atomic arrangements for GaN/SiC interface. Although the polarity of GaN on SiC has been studied by a number of techniques, many experimental results in the literature are in conflict. The above theoretical argument is considered as ‘‘standard framework’’, where Ga-face GaN grows on Si-face SiC and N-face GaN grows on C-face SiC [25]. With a few exemptions such as the work of Sasaki and Matsuoka [242], where they concluded that the epitaxial GaN layers on (0 0 0 1)Si and (0 0 0 1)C SiC were terminated with N and Ga, respectively, most results are consistent with ‘‘standard framework’’. 3.4.7. Zincblende GaN Growth on 3C-SiC/Si(0 0 1) 3C-SiC(0 0 1) films deposited on Si(0 0 1) have been favored for producing zincblende GaN films due to the relatively small lattice constant mismatch (3.4%) and its excellent thermal stability of the 3C-SiC. However, 3C-SiC/Si(0 0 1) substrates generally suffer from the very poor crystal quality and rough surfaces, both of which tend to promote the nucleation of wurtzite GaN. Furthermore, zincblende AlN is very difficult to stabilize—wurtzite AlN is strongly favored regardless of the substrate or its orientation, thus, it is not effective as a buffer for growth of zincblende GaN. Several groups have studied zincblende GaN epitaxy on 3C-SiC/Si(0 0 1) [243]. The crystal quality and polytype purity of the GaN in MOCVD growth is greatly affected by the V/III ratio as well as the temperature [244]. More works have been accomplished by MBE [243,245–247] due to its low growth temperature, which tends to minimize the amount of thermodynamically stable wurtzite GaN [248]. Daudin et al. [248] found that the V/III ratio affect the wurtzite/zincblende ratio in MBE deposited films, while a N-rich condition was detrimental for the growth of zincblende phase. The experimental results can be explained by the assumption that MBE growth of zincblende

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GaN is mainly governed by the impinging active N flux, which directly determines the mean-free path of gallium adatoms. It was also found that flatter substrate surface improves film quality, and LPCVD produced 3C-SiC/Si(0 0 1) substrate will result in better GaN film than CVD produced one [243,248]. Recently, directly carbonization of Si(0 0 1) surface has been utilized to produce thin (thickness of several nanometers) 3C-SiC film with flat surface (RMS only 0.3 nm) in MBE growth [249,250]. High quality zincblende GaN films were grown on these substrates and it was also found that Ga-rich growth condition is beneficial for growing high quality film with atomically flat surface [249]. The quality of zincblende GaN films is normally worse than wurtzite ones, as the FWHM of rocking curve for cubic GaN films is usually several tens of minutes, compared to as little as 2–3 arcmin for wurtzite ones. More discussion of zincblende GaN growth is presented in the section on GaAs substrates.

4. Si The favorable physical properties, high quality, and low cost of silicon (Si) make it a very attractive substrate for GaN-based devices. Silicon wafers are very low price and are available in very large size due to its mature development and large-scale production. Silicon has good thermal stability under GaN epitaxial growth conditions. The crystal perfection of silicon is better than any other substrate material used for GaN epitaxy and its surfaces can be prepared with extremely smooth finishes. The possibility of integrating optoelectronic GaN devices with Si electronic devices is also very attractive [251]. To date, the quality of GaN epitaxial layers on silicon has been much poorer than that on sapphire or silicon carbide, due to large lattice constant and thermal expansion coefficient mismatch, and the tendency of silicon to form an amorphous silicon nitride layer when exposed to reactive nitrogen sources. GaN and AlN grown on Si(1 1 1) are highly defective and nonradiative carrier recombination channels reduce the luminescence efficiency of many optoelectronic devices. Nevertheless, the incentives for using Si substrates remain high, and good progress in reducing the defect density by using epitaxial lateral overgrowth or pendeoepitaxy has been reported. Many GaN devices have been demonstrated on Si substrates including LEDs and HEMTs. 4.1. Structure and properties of Si Si has a diamond-lattice structure with the space group of Fd3m (no. 227), which belongs to the cubic-crystal family. It can be represented as two interpenetrating fcc sublattices with one sublattice displaced from thepother ffiffiffi by one quarter of the distance along a body diagonal of the cube (i.e. the displacement of a 3=4, where a ¼ 0:543 nm). Each atom in the lattice is surrounded by four equidistant nearest neighbors that lie at the corners of a tetrahedron. Fig. 21 illustrates the perspective views along the [0 0 1], [0 1 1] and [1 1 1] directions of a Si unit cell. The properties of undoped and doped Si have been well characterized. Table 11 lists physical, chemical, thermal, mechanical and optical properties of Si at room temperature. 4.2. Bulk crystal growth of Si The bulk crystal growth of silicon is quite mature and the quality of its wafers is very high. Silicon single crystal ingots are usually produced by either the Czochralski (CZ) method, or the float zone (FZ) method. These ingots eventually become thin Si wafers through the processes of shaping, slicing, lapping, etching, polishing, cleaning and inspection.

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91

Fig. 21. Perspective views of Si along various directions: (a) [0 0 1]; (b) [0 1 1]; (c) [1 1 1].

The CZ growth method is much more common due to its capability of easily producing crystals with large diameters of several inches (up to 12 in. today), and over 85% of silicon crystals are grown by this method today. For growing doped crystal, impurities can be added directly to the melt to create p- and n-type silicons. The only significant drawback of the CZ method is that the impurities, mainly oxygen (typical at a level of 1018 cm
3) and carbon (typical at a level of 1016 cm
3), can be incorporated into the growing ingot from the erosion of the quartz crucible and from graphite furnace fixtures. However, the incorporation of these impurities can be beneficial as oxygen increases the yield strength and can act as internal getter to tie up metallic contaminants [263]. The FZ method makes smaller-diameter wafers than the CZ method. It does not use any crucible, therefore, the residual impurity concentration in the grown ingot is markedly reduced, making it easier to grow high resistivity material. Doping of the crystal can be accomplished either by starting with a doped poly-silicon rod, a doped seed, or by maintaining a gas ambient during the FZ process that contains a dilute concentration of the desired dopant. FZ material is primarily used in applications which requiring high resistivities, low oxygen content, or both, such as some detectors or power devices. These applications represent a small fraction of the silicon market [263]. Table 11 Properties of Si at room temperature Properties

Value

Reference

Lattice constant (nm) Density (g/cm3) Melting point ( 8C) Heat capacity (J/g K) Thermal conductivity (W/cm K) Thermal diffusivity (cm2/s) Thermal expansion(linear)  10
6 K
1 Percentage change in lattice constant: Da/a0 (298–1311 K) Shear modulus (GPa) Bulk modulus (GPa) Young’s modulus (GPa) Poisson’s ratio Refractive index Static dielectric constant Electrical resistivity (undoped) (kO cm) Energy band gap (eV) Electron mobility (cm2/V s) Hole mobility (cm2/V s)

0.543102 2.3290 1414 0.70 1.56 0.86 2.616 0.3995 680 97.74 165.6 0.218 3.42 11.8 Up to 50 1.124 1430–1448 473–495

[252] [252] [253] [254] [255] [256] [257] [258] [259] [259] [259] [259] [254] [254] [254] [260] [261] [262]

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4.3. GaN epitaxy on Si(0 0 1) substrates Both zincblende and wurtzite GaN epilayers have been grown on the Si(0 0 1) by MBE [264–267], MOCVD [268–271] or HVPE [272]. The GaN films grown directly on Si(0 0 1) are phase mixed [265], or wurtzite phase only [264,267,271]. Due to the large lattice constant mismatch, it is difficult to grow pure zincblende phase GaN directly on the Si(0 0 1). Furthermore, an amorphous SixNy layer may form at the GaN/Si interface introducing a phase mixture [273]. Buffer layers are typically deposited first on the Si substrate to enhance wetting and reduce the reactivity of the Si substrate, and to provide a better lattice constant match between the film and substrate—a transition to reduce the lattice constant mismatch. A thick layer of 3C-SiC produced by CVD or a thin layer by direct carbonization is commonly used to eliminate this amorphous layer and promote zincblende GaN growth. GaN epitaxy on 3C-SiC/Si is addressed in zincblende phase growth of the SiC substrate section. The utilization of other buffer layers, such as g-Al2O3 [269], AlN or AlxGa1
xN [270,272] resulted in the wurtzite phase growth. The orientation relationship was GaN(0 0 0 1) || Si(0 0 1) and GaN[1 1 2 0] || Si[1 1 0] [269,273] or AlN[1 1 0 0] || Si[1 1 0] for another domain with a 308 rotation for AlN epilayer on Si(1 0 0) by MBE [273]. Fig. 22 shows the orientational relationship of wurtzite AlN to the Si(0 0 1) substrate. Although the GaN films grown on GaAs/Si(0 0 1) were predominantly zincblende, there were additional structural defects as compared with the films on GaAs(0 0 1) indicated by the low efficiency of PL [265]. The growth of pure zincblende GaN film on Si(0 0 1) remains a real challenge. 4.4. GaN epitaxy on Si(1 1 1) substrates Si(1 1 1) is generally the preferred substrate for GaN epitaxy. The quality of GaN on these substrates has been significantly improved and the devices such as field-effect transistors [274],

Fig. 22. Atomic arrangement for the heteroepitaxial nucleation of 2H-AlN on the Si(0 0 1) surface. The two AlN domains with a 308 rotation are formed on neighboring terraces (1) and (2) separated by single atomic step boundary according to the Si dangling bond-directions (after [273]).

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Schottky-barrier ultraviolet detectors [275], and blue LEDs [276] have been achieved in the past few years. Table 12 lists the recent GaN epitaxial growth on Si(1 1 1) by MBE, MOCVD and HVPE since 1997. Since Si is a non-isomorphic substrate, surface steps will possibly lead to an SMB, as discussed in the section of the SiC substrate. Martin et al. [288] showed surface preparations reducing irregularities on the substrate surface greatly improved the crystalline and luminescent quality of a GaN epilayer grown by MBE. The X-ray rocking curve FWHM decreased from 18 to 20 arcmin. To overcome the problem of a large difference in lattice parameters, and the strength of the Si–N bond, different buffer layers such as AlAs [284,285], oxidized AlAs [282], GaAs [277], AlN [277–281,283,289], LT GaN [287,290], ZnO [287], and 3C-SiC(1 1 1) [249,291,292] have been tested for Si(1 1 1) substrates. The choice of buffer layer is still an open problem, and the quality of consequently deposited GaN epilayers may strongly depend on the buffer layer properties. In addition, to avoid the formation of silicon nitride, it is a common practice in MBE growth to initiate the AlN buffer layer growth by exposing the Si substrate to the flux of aluminum and ammonia alternatively for a short time [277,279]. These techniques help suppress the formation of SiNx at the interface and result in a 2D growth mode of GaN. Fig. 23 shows the amorphous SixNy layer at the GaN/Si(1 1 1) interface when a GaN layer is grown directly on the cleaned Si(111) surface by MBE and the resulting GaN epilayer was polycrystalline [281]. An AlN buffer layer can prevent the formation of SixNy. Furthermore it also promotes the 2D growth of GaN, and is more used for wurtzite GaN epitaxy on Si(1 1 1). The TEM and HREM images of AlN films deposited at 640–720 8C on Si(1 1 1) by MBE shown in Fig. 24 display an abrupt interface and a decrease in defects along the thickness [293]. Hellman et al. [294] investigated the nucleation behavior of AlN on Si(1 1 1) surface by MBE. In spite of the 23.4% misfit in the AlN/Si system, optimized AlN films grown on a (7  7) reconstructed Si surface made excellent substrates for GaN heteroepitaxy. By deliberately choosing the growth condition in MBE, the AlN film grows on Si in a two dimensional mode with an atomically smooth surface [278,295]. In MOCVD growth, HTs grown AlN buffer layer is much better than LT one and it is necessary to enhance 2D growth for GaN epitaxy [286,296]. The greater thermal expansion coefficient of GaN as compared to that of Si (5:6  10
6 K versus 2:6  10
6 K) puts the film in tension, hence cracks may be generated in the GaN during the cool-down cycle. Follstaedt et al. [297] deposited 2.2 mm thick GaN layer at 1060 8C on a 1080 8C AlN buffer layer by MOCVD. While the GaN surface was smooth immediately after growth, cracking occurred along the equivalent {1 1 0 0} GaN planes upon cooling to room temperature.

˚ ) SixNy layer at the GaN/Si(1 1 1) Fig. 23. HRTEM micrograph showing the development of an amorphous, thin (20–30 A interface (after [281]).

93

94

Method

Buffer layer

T (8C)

Thickness or rate

Characterization

Reference

MBE

AlN, at 890  30 8C

730

0.4–1.1 mm/h

[277]

MBE MBE

AlN, 125 nm at 650 8C AlN, 0.1 mm/h at 850 8C

770 800

880 nm 1.5 mm at 1.0 mm/h

MBE

AlN, different thickness with maximum as 125 nm AlN, 0.2–0.5 mm/h at 850 8C near stoichiometry condition AlOx, oxidized from AlAs deposited at 460 and 760 8C AlN, 80 nm at 1100 8C; AlGaN, 250 nm at 1080 8C AlAs 20–30 nm at 425 8C then 750 8C, and LT GaN at 500 8C AlAs 40 nm at 450 8C then 720 8C, and AlN at 950 8C AlN at 1060 8C

850

PL peak at 3.408 eV with FWHM < 40 meV at room temperature RHEED, PL peak at 3.453 eV at room temperature RHEED, XRD o scan with FWHM 17 arcmin, PL peak at 3.471 eV with FWHM 5 meV at 10 K AFM, STM, PL peak at 3.455 eV at 10 K

MBE MOCVD MOCVD MOCVD MOCVD MOCVD HVPE

ZnO 500 nm at 400 8C, LT GaN 500 nm at 600 8C

660–770

0.5 mm/h

Multistep

[278] [279] [280]

RHEED, XRD o scan with FWHM 8 arcmin, PL peak at 3.465 eV with FWHM 15 meV at 10 K XRD, CL, comparable with grown on sapphire

[282]

[281]

1080

1.0 mm

RHEED, XRD o scan with FWHM 10 arcmin

[283]

1050

1.0 mm

AFM, SIMS, PL, XRD o scan with FWHM 45 arcmin

[284]

1010–1040

1.1 mm

AFM, PL, XRD o scan with FWHM 610 arcsec

[285]

1060 8C

0.5 mm

[286]

880–1020

35 mm/h for 10 min

AFM, XRD o scan with FWHM 656 arcsec, RMS 0.3 nm on 1 mm2 AFM, XRD, PL with peak at 3.45 eV

[287]

L. Liu, J.H. Edgar / Materials Science and Engineering R 37 (2002) 61–127

Table 12 Recent GaN epitaxy on Si(1 1 1)

L. Liu, J.H. Edgar / Materials Science and Engineering R 37 (2002) 61–127

Fig. 24. (a) TEM image of an AlN film in cross-section, under multibeam conditions along the h0 1 1i axis. Note the ˚ in the AlN film. (b) HREM image of the Si(1 1 1)/AlN interface observed decrease in defect density occurring at 250 A along the Sih0 1 1i axis. A step in the silicon substrate inducing a rotation of the AlN planes on the top occurs in the center (after [293]).

The sample cracked into small areas, with separation between cracks from 20 to 100 mm. The HT AlN buffer layer did not relieve the stress between the substrate and the GaN epilayer. With an AlN buffer layer deposited at low temperature such as 700 8C, or an intermediate layer consisting of AlN and AlGaN at high temperature (1100 8C) [283], or a thick AlN-to GaN graded buffer layer [298], or using AlN/GaN supperlattices [299], the crack problem was eliminated. Nikishin et al. [277] reported a biaxial stress of 160 MPa and an in-plane stress þ0.09% for GaN epilayers deposited by MBE varied from between 0.4 and 2.2 mm. Kim et al. [313] demonstrated that step-graded Alx Ga1
x N buffer layer could substantially reduce the crack density, and improve the structural and optical properties of GaN epilayers grown by MBE. Marchand et al. [298] also confirmed that the thickness of AlN buffer layer will change the sign of stress (tensile or compressive) in AlN buffer layer and stress magnitude in GaN epilayer. GaN films grown by MOCVD on a thick (800 nm), AlN-to-GaN graded buffer layer were under compressive stress and crack free. XRD measurements consistently verified the orientation relationship of GaN[0 0 0 1] || Si[1 1 1] [278,279,281–284]. The diffraction pattern taken on the GaN/AlGaN/Si(1 1 1) interface by MBE at 750 8C shows the relationship between Si and GaN to be GaNh0 0 0 1i || Sih1 1 1i and GaNh2 1 1 0i || Sih0 1 1i [276]. Meanwhile, the growth of AlN on the Si(111) by MBE showed two epitaxial relationships [293]: (1) AlN(0 0 0 1)[2 1 1 0] || Si(1 1 1)[0 2 2], for deposition temperature >650 8C; (2) AlN(0 0 0 1)[1 0 1 0] || Si(1 1 1)[0 2 2]. These two relationships corresponded to the (4  4) and (7  7) lattice, respectively. The LT grown AlN by MBE at 400–600 8C showed the relationship of AlN(0 0 0 1)[0 1 1 0] || Si(1 1 1)[1 1 2] [300]. This result is similar to previous works [301,302]. Zhao et al. [303] determined that GaN deposited on Si(1 1 1) with an AlN buffer layer by MOCVD had the Ga-polarity. Hellman [25] concluded that GaN grown on an AlN buffer layer on Si(1 1 1) was Ga faced and GaN grown directly on Si(1 1 1) was N faced according to the standard framework. The use of compliant silicon-on-insulator (SOI) substrates or porous silicon (PS) has been applied for GaN growth both by MBE [304–306] and MOCVD [307–309]. This substrate improves the quality of epitaxial GaN layers by releasing the strain and absorbing the generated threading dislocations in the thin Si overlay of the SOI substrate. Meanwhile, selective area growth has also been applied on the stripe-patterned and dot-patterned Si(1 1 1) by MOCVD [310,311] and MBE [312]. It opens new chances for growing low-defect GaN heterostructures. Other growth methods such as ultrahigh vacuum chemical vapor deposition [313], radio-frequency-magnetron sputtering [314], and ion-beam assisted evaporation [315] have also been investigated.

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5. GaAs The material property advantages of GaAs as a substrate for GaN epitaxy include its isoelectronic structure (i.e. both GaAs and GaN are III–V compounds), the shared element (Ga), the potential to convert the surface of GaAs to GaN, and parallel cleavage planes between the film and substrate. The technological advantages include a well-established process technology, several readily available substrate orientations of both polar and non-polar varieties, and low resistance ohmic contacts. There are several disadvantages to GaAs as a substrate for GaN epitaxy, including a large lattice constant and thermal expansion coefficient mismatch, a poor thermal conductivity, and perhaps most problematic, low thermal stability. The principle areas of research on GaAs as a substrate for GaN epitaxy include: obtaining pure (wurtzite-free) zincblende GaN on GaAs(0 0 1); growing thick wurtzite GaN films on GaAs(1 1 1) substrates; developing growth processes which maintain smooth and abrupt GaN/GaAs interfaces and avoid GaAs decomposition; and minimizing the dislocation density and improving the crystal quality in the GaN epitaxial films. Maintaining a smooth surface during all stages of GaN film growth is particularly important for all of GaN on GaAs, as it is a major cause of crystal defects and polytype inclusions. GaAs(0 0 1) is one of the few semiconductor substrates on which metastable zincblende GaN epitaxial films readily form, and many researchers have investigated the best ways of avoiding any inclusion of the wurtzite polytype into these films. GaAs is much more readily wet etched than sapphire, making GaN films easier to separate from GaAs than sapphire. GaAs also has a lower yield strength than sapphire so cracking and plastic deformation are more likely to occur in the substrate rather than the GaN film. Potentially, GaAs(1 1 1) substrates are better templates for creating freestanding thick GaN films for subsequent epitaxy and device fabrication, with the ultimate goal of eliminating the problems associated with heteroepitaxy. Since the decomposition rate of GaAs in NH3 or an ultrahigh vacuum rapidly increases at temperatures above 700 8C, this imposes limits on the epitaxial growth temperature of GaN, and hence, its maximum growth rate. Even a small amount of GaAs decomposition is detrimental to zincblende GaN epitaxy, as surface roughening or faceting enhances the onset of wurtzite growth. Because MBE is capable of depositing epitaxial GaN films at a lower temperature, it has been more commonly employed than MOCVD or HVPE when GaAs is the substrate. There has been some success in circumventing the maximum allowed temperature, once the GaAs substrate is completely encapsulated with GaN deposited at a LT, thereby making MOCVD and HVPE more viable. 5.1. Structure and properties of GaAs GaAs has the same structure as zincblende GaN (see Fig. 2). Table 13 lists the physical, chemical, thermal, mechanical and optical properties of GaAs of importance to the GaN epitaxy. GaAs is less stable than SiC or sapphire. Above 800 8C its decomposition rate to liquid gallium and arsenic vapor is considerable. Fig. 25 shows the thermodynamic temperature–pressure relation for GaAs. The upper curves are the vapor pressures of the pure components (i.e. PGa over pure gallium and PAs2, PAs4 over pure arsenic) and the lower curves are the vapor pressures of the components over GaAs. In particular, the activation energy of arsenic dimers desorption is between 2.6 and 3.9 eV [324,325]. It is common practice in MBE to anneal the GaAs substrate at 600 8C in the presence of an As flux to desorb its native oxide and prevent the preferential desorption of As from the surface. The kinetics of GaAs decomposition has been modeled by Karpov and Maiorov [325,326], and its relevance to GaN growth has been considered [327].

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97

Table 13 Properties of GaAs at room temperature Properties

Value

Reference

Lattice constant (nm) Density (g/cm3) Melting point (8C) Heat capacity (J/g K) Thermal conductivity (W/cm K) Thermal diffusivity (cm2/s) Thermal expansion(linear)  10
6 K
1 Percentage change in lattice constant: Da/a0 (300–1200 K) Bulk modulus (GPa) Young’s modulus (GPa) Poisson’s ratio Refractive index Static dielectric constant Electrical resistivity (undoped) (O cm) Energy band gap (eV) Electron mobility (cm2/V s) Hole mobility (cm2/V s)

a ¼ 0.56536 5.32 1240 0.327 0.45 0.26 6.03 0.5876 75.0 85.5 0.31 3.66 13.1 1.0  104 1.423–1.430 8500 400

[316] [317] [318] [319] [320] [320] [321] [19] [322] [323] [323] [323] [323] [323] [323] [323] [323]

5.2. Bulk crystal growth of GaAs Currently there are two prominent GaAs crystal growth techniques: liquid encapsulated Czochralski (LEC) and vertical gradient freeze (VGF). Intense efforts have been undertaken

Fig. 25. T–P phase diagram of GaAs (after [328]).

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worldwide to overcome problems such as crystal defects, impurities and micro-inhomogeneities by improving these techniques [329–331]. Today, GaAs wafers are available in many diameters (up to 150 mm), several orientations, doping types, and impurity concentrations. Both silicon and tellurium are used as n-type dopants in GaAs substrates, which are available with electron carrier concentrations 1016 to 1018 cm
3. Zinc is the standard p-type dopant and substrates with hole carrier concentrations of 1018 to 1019 cm
3 are available. Commercial LEC GaAs typically has an etch pit density (EPD) <104 cm
2, and electron mobility >4000 cm2/V s. VGF grown GaAs substrates have a lower defect density, with EPD almost an order of magnitude less. Both (0 0 1) and (1 1 1) types of GaAs with different vicinal degrees are available commercially. 5.3. GaN epitaxy on GaAs(0 0 1) substrates In principle, zincblende structure GaN possesses superior electronic properties for device applications, such as a higher mobility, isotropic properties due to its cubic symmetry and high optical gain. In practice, such advantages have not been realized due to the difficulty in producing low defect content, polytype pure material. To produce zincblende GaN requires substrates with four-fold symmetry, that is (0 0 1) substrates. Zincblende GaN has been prepared on several substrates including Si, 3C-SiC, GaP, and MgO, but the most commonly employed substrate for this is GaAs(0 0 1). Since zincblende GaN was first grown on GaAs substrates in 1992 [332], many additional studies have been done. The results of studies through 1996 were summarized by Okumura et al. [243]. A summary of the most recent works is given in Table 14. In certain cases both zincblende and wurtzite structures of GaN films may coexist because of the small difference between their energies of formation [344,345]; therefore, the control of purity of zincblende phase in the GaN film is a major concern and critical problem for the GaN epitaxy on GaAs(0 0 1) substrate. 5.3.1. Surface preparation The direct nucleation of GaN on as-received GaAs is difficult due to the shorter bond length in GaN compared to GaAs, and the presence of a native oxide. A nitridation step alone has proved insufficient to improve GaN nucleation as it produces a rough surface; and subsequently deposited GaN epitaxial films contain DPBs and polytype inclusions. The quality of the nitridated GaAs can be further compromised if residual oxygen is present on the surface during nitridation. In practice, a GaAs prelayer is usually necessary to improve GaN wetting and provide an atomically smooth substrate surface before nitridation [346–348]. This also reduced the effects of residual oxygen on the surface which affects nitridation [349]. The growth temperature for a GaAs prelayer is usually about 600 8C and the thickness is typically 200 nm, but up to 1 mm [350] has been used. Sometimes, several minutes of annealing under As4 flux is also utilized after the initial growth of GaAs layer to improve the surface smoothness further [347]. The presence of As ambient pressure seems to protect the surface from degradation at high substrate temperature and to enhance subsequent zincblende phase GaN growth [351]. The best GaN films are produced when a thin GaAs layer is first deposited and then nitridated before GaN nucleation. This helps to maintain an atomically smooth surface, which is critical for minimizing defects. Trampert et al. [346] and Brandt et al. [347] found that GaN nucleates in the metastable zincblende phase on atomically smooth 100 nm thick GaAs(0 0 1) prelayer by plasma-assisted MBE, with the same orientation as the substrate. An atomically abrupt GaN/GaAs interface was produced, although there is a high density of dislocations generated in the GaN layer at the interface, identified as stacking faults and microtwins propagated along the {1 1 1} planes, as seen in Fig. 26.

Method

Orientation

N source

Buffer layer or nitridation

Growth temperature (8C)

Growth V/III

Main phase and thickness (nm) or rate

Characterization

Reference

MBE

(0 0 1)

RF N2 300 W

620

1.0

b, 50–100

TEM

[333]

MBE MBE MBE

(0 0 1) (0 0 1) (0 0 1)

RF N2 200 W RF N2 RF N2

580 720 740

Varying 1.0 0.5–1

b, 0.1 ml/s b, 1600 b, N/A

HRSEM CL: 3.26, 3.17 and 3.08 eV Raman, peaks at 555 and 741 cm
1

[334] [335] [336]

MBE

(0 0 1)

RF N2

700

1.0

b, 400

(0 0 1)

RF N2 500 W

o/2y (0 0 2) Brag peak FWHM 11 and 35 arcmin o (0 0 2) FWHM 10 arcmin

[337]

MBE MOCVD

(0 0 1)

NH3

Eight monolayer GaN BL under As4 N/A N/A 10–20 monolayer of GaAs BL at 600 8C, then 10–20 monolayer of GaN nucleation at V/III ¼ 4, 600 8C 600 8C GaAs BL under As flux 200 nm GaAs BL at 600 8C, then nitridation 10–20 nm 550 8C GaN BL

MOCVD MOCVD

(0 0 1) (0 0 1)

NH3 DMHy

MOCVD

(0 0 1)

DMHy

a

3 min 580 8C nitridation 20 nm GaN BL at 600 8C under V/III ratio of 100 100 nm 700 8C GaAs BL, 20 nm 575 8C GAN BL at V/III ¼ 100

b, 500, 150 nm/h

700 750–790

400–1200

b, 800

550 600–820

50

b, 1 mm/h b, 0.5 mm/h

900

25

b, 2 mm

PL 387 nm, (0 0 2) o scan FWHM 0.25–0.78 TEM, TED XRD PL peak at 3.26 eV with FWHM 13 meV at 5.3 K

[338,339] [340] [341] [342] [343]

MBE, molecular beam epitaxy; MOCVD, metalorganic chemical vapor deposition; RF, radial frequency plasma; DMHy, dimethylhydrazine; BL, buffer layer; HRSEM, high resolution scanning electron microscopy; RHEED, reflection high energy electron diffraction; FWHM, full width half maximum; PL, photoluminescence; TEM, transmission electron microscope; TED, transmission electron diffraction; a, wurtzite phase; b, zincblende phase.

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Table 14 Most recent works of GaN expitaxy on GaAs since 1997a

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Fig. 26. (a) SAD pattern and (b) cross-sectional HRTEM image of the interface region of the GaN/GaAs(0 0 1) heterostructure taken along the h1 1 0i direction. The arrows in (b) depict an atomically flat part of the interface containing regularly arranged edge-type misfit dislocations. (c) Magnified HRTEM micrograph showing the presence of dislocations (arrows) (after [346]).

These results indicate that first depositing a GaAs prelayer help the nucleation of GaN on a GaAs substrate and may result in zincblende growth. MBE systems are generally well equipped to deposit both GaAs and GaN, hence it has been natural for researchers to consider the benefits of first depositing a thin GaAs film on the GaAs substrate before GaN growth. In contrast, most CVD reactors are not equipped to do an analogous experiment with GaN on GaAs substrates. One concern would be the residual doping effect of Si on the GaAs film. 5.3.2. GaAs nitridation As with other substrates, nitridation has been applied to GaAs to improve GaN nucleation and to mitigate the lattice constant mismatch. The nitridation of the GaAs substrate results in the formation of a thin nitride surface layer, which provides a template for the epitaxial growth and improves the quality of the GaN layer. This nitride layer also helps to suppress GaAs decomposition. Successful nitridation requires the formation of an atomically smooth layer. This is challenging, as the surface tends to form facets during nitridation. Generally, GaAs substrates are nitridated either by exposure to a nitrogen plasma (MBE) or annealing in ammonia (MOCVD and MBE). Fig. 27 shows a schematic of the mechanistic steps involved in GaAs nitridation. Using the Auger electron spectroscopy to investigate the processes during the initial stage of nitridation of the GaAs(0 0 1) surface by RF generated atomic nitrogen, Aksenov et al. [352] found that the nitridation process was greatly dependent on the temperature. At a LT (<200 8C), the nitridation was hindered by kinetic limitation, while at a HT (>500 8C) simultaneous etching of the surface occurred along with the nitridation process. Complete nitridation produces a pure GaN layer [353], but can lead to a highly facetted interface between the GaN layer and GaAs substrate. The root-mean-square roughness of the surface after nitridation can be controlled on the order of 0.25 nm [354,355]. GaAs(0 0 1) nitridation does not occur uniformly over the surface, but proceeds fastest along {1 1 1} facets into the underlying GaAs layer and leads to faceted GaN/GaAs interface [356,357]. Two methods have been employed to improve the GaN quality at the interface: one is to use the highest growth rate possible to quickly bury the interface; the other is to maintain an As flux during the growth of the first few monolayers of the GaN, which helps to prevent the deterioration of the GaAs surface [333].

L. Liu, J.H. Edgar / Materials Science and Engineering R 37 (2002) 61–127

Fig. 27. Schematic diagram showing the model for nitridation of GaAs(0 0 1) (after [353]).

Conflicting conclusions have been made as to the effect of GaAs nitridation on the resulting phase of deposited GaN. Some studies claim that a nitrogen-deficient condition leads to the formation of metastable zincblende GaN [340,348,358,359]. Kikuchi et al. [348] carried out an extensive investigation on the GaAs substrate nitridation effect by molecular beam epitaxy using RFradial nitrogen source. In their work, the GaAs(1 1 1)A (Ga-polarity) substrate resulted in wurtzite GaN independent of substrate nitridation. Zincblende GaN was produced on the GaAs(0 0 1) substrates without nitridation and wurtzite GaN grown on sufficiently nitridated substrates. Their result is in contrast to others [360,361] which they attribute to the different nitrogen source in their work. Sato [359] used the plasma-assisted nitrogen radicals to improve the growth rate for MOCVD of GaN on a GaAs substrate, finding that the zincblende phase is grown under Ga-rich conditions, and wurtzite phase grown under is a N-rich condition. Other papers report contrary results to the above cases, instead assuming that the active nitrogen atoms impinging on the GaAs surface diffuse into the substrate and replace the arsenic atoms occupying the group V lattice of zincblende structure, forming a zincblende GaN template [337,356]. This exchange process is energetically favorable if we consider the bond strength only. Two displaced arsenic atoms will recombine to form a dimer which desorbs from the growth surface. Under this assumption, sufficient nitridation is necessary to provide a zincblende GaN template. Clearly the effect of nitridation on the phase of GaN is a complicated one. It may depend on the source of nitrogen, radials or ions, nitrogen sufficiency (flux, V/III ratio), and nitridation temperature. Li et al. [337] proposed a simple kinetic growth model for MBE growth: that the existence of excess Ga atoms at the surface, which is formed due to the desorption of arsenic at HT, results in wurtzite GaN nucleation. This may explain why supplying an excess As or N at the initial growth and higher growth temperature are important for obtaining purely zincblende GaN. But excessive nitridation at HT will lead to wurtzite growth due to the increasingly rough morphology as a result of the surface decomposition of the GaAs [354]. So enough nitridation but still keeping a

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smooth surface is critical for providing a zincblende growth template. Due to the different nitrogen source and nitridation procedure, different phase selection after nitridation may occur. In addition, an arsenic flux during nitridation also plays an important role in the phase selection. The existence of arsenic was beneficial to the formation of zincblende GaN. Jung and Kumagai [350] found that a high background pressure of As suppresses nitridation and results in a dominant zincblende GaN phase. The effect of arsenic seems to support the assumption of deficient nitridation bringing about zincblende GaN. Actually, it may suppress the formation of gallium at the GaAs surface, hence, bringing about smooth surface and zincblende GaN nucleation. 5.3.3. Buffer layer The deposition of a buffer layer of GaN at LT following the initial nitridation generally improves the GaN quality. AlN is generally extremely difficult to deposit as pure zincblende material, so it is not used as a buffer layer in zincblende GaN growth on GaAs substrates. Other buffer layer such as AlAs has lead to single-phase hexagonal GaN growth [362,363]. Usually the initial deposition of GaN is conducted at relatively LTs (550–650 8C) compared to that for epitaxial film growth. This film can protect the GaAs surface at relatively high growth temperatures (>700 8C) in addition to enhancing GaN nucleation and initiating growth [340,364]. High resolution TEM and electron diffraction measurements show that a GaN buffer layer by HVPE at 550 8C is polycrystalline as deposited, but becomes a single crystal upon thermal annealing at 850 8C for 10 min, as shown in Fig. 28 [365]. The thickness of this initial layer and growth temperature have a strong effect on the crystalline quality of the zincblende GaN quality in MBE or MOCVD growth [337,366]. Usually only few GaN monolayers are used as buffer layer. In the MBE growth by Li et al. [337] (using a radio frequency N2 plasma discharge source), a two-monolayer thick GaN interfacial layer deposited at 600 8C resulted in better quality films without any wurtzite phase for the subsequent growth at 700 8C. Thin initial layers were superior to thick layers because the quicker temperature increase for the thinner buffer layer will enhance the migration of excess gallium and result in zincblende nucleation/growth. LT GaN buffer layers are also utilized in MOCVD growth of GaN on GaAs substrate. In Sato’s plasma assisted MOCVD growth [359], a very thin initial GaN layer grown at 550 8C for 2 min was used to deposit cubic GaN under Ga-rich condition at a temperature of 600–700 8C. In Wu et al.’s work [367,368], a 20 nm GaN buffer layer was deposited at 575 8C while the growth temperature was 900 8C for the MOCVD. The buffer layer was found essential to achieve good quality film,

Fig. 28. (a) High-resolution TEM image and (b) electron diffraction pattern taken near the interface between GaN buffer layer and GaAs substrate (after [365]).

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which has FWHM 70 meV for PL excitonic transition at 3.274 eV (300 K). Xu et al. [369] deposited LT GaN buffer layers under different conditions and found that the surface roughness decreased as the growth temperature increased, while the interstice of islands in buffers increased. Higher temperatures and longer deposition time for GaN buffer layer in MOCVD caused more stacking faults line and introduced more hexagonal phase inclusions [366]. The buffer layer deposition conditions (both temperature and time) changed the hexagonal phase content and determined the polytype purity and quality in zincblende GaN. The GaN growth temperature can control the phase formation of the epitaxial layer [370,371]. Presumably HT (>850 8C) prefers wurtzite growth. So both the deposition temperature and thickness of the buffer layer are critical to obtain a pure zincblende GaN film. Exposure of the layer surface to As flux during the growth of the first few GaN monolayers was also beneficial for the growth. It produced a remarkably flat GaN/GaAs interface [333] and pure zincblende GaN on the GaAs(0 0 1) substrate [338,372]. Li and coworkers [337,339] obtained high quality zincblende GaN with FWHM of rocking curve as only 10 arcmin by deposition an optimized GaN initial layer (two-monolayer-thick) under As pressure in MBE. 5.3.4. Substrate misorientation Only a few studies have been published on the effect of substrate misorientation on zincblende GaN on GaAs(0 0 1) [348]. Nakayama et al. [373] studied the substrate misorientation dependence of the wurtzite phase inclusion in GaN MOCVD. The sample grown on the tilted GaAs(0 0 1) surface toward ½1 1 0 showed an enhanced generation of the wurtzite domain on the ð1 1 1Þ face, whereas the wurtzite domain on the (1 1 1) face was suppressed. The sample grown on the tilted surface toward [1 1 0] showed the generation of the wurtzite domains on both the (1 1 1) and (1 1 1) faces with an equal magnitude. They proposed that the generation of the wurtzite domains may be suppressed when the exposure of the (1 1 1) or (1 1 1) faces by the thermal damage of the substrate surface is reduced. With a LT GaN buffer layer on 28 miscut GaAs(0 0 1), Yang et al. [374] obtained a GaN epitaxy film that was primarily the zincblende phase, deposited by HVPE method with low quality (FWHM 4.88). But, this 28 miscut GaAs(0 0 1) surface produced both wurtzite and zincblende GaN thin film, as shown in Fig. 29 [375], depending on the growth conditions of MBE. The zincblende GaN growth on patterned GaAs(0 0 1) substrates has also been investigated [376]. Lateral epitaxial overgrowth may reduce the dislocation density greatly, as has been achieved on sapphire and 6H-SiC. The incorporation of wurtzite GaN was reduced by using ½0 1 1-oriented stripes to suppress the formation of (1 1 1)B facets.

Fig. 29. High resolution TEM images of (a) wurtzite layer with nitridation, (b) zincblende GaN layer without nitridation, grown on GaAs(0 0 1) misorientated by 28 toward [0 0 1] (after [375]).

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5.4. GaAs(1 1 1) substrates for wurtzite phase growth and templates of free-standing epitaxial films Generally, GaAs(1 1 1) substrates are used for wurtzite GaN growth. Both GaAs(1 1 1)A and GaAs(1 1 1)Bð1 1 1Þ can result in single phase wurtzite GaN in MBE growth [372,377,378]. Foxon’s group found that (1 1 1)B substrates produced better zincblende GaN (FWHM of XRD o scan 720 s) than (1 1 1)A [372,377,378] in MBE. For the growth on (1 1 1)A oriented substrates, a thin amorphous layer at the substrate surface resulted polarity inversion and higher defect densities. If GaAs decomposition at temperature 850 8C can be avoided, GaN epitaxy by MOCVD or HVPE is attractive as very high growth rates are possible. GaAs(1 1 1) have been used for growing free-standing wurtzite GaN [370–382]. It was demonstrated that GaAs substrates can withstand GaN growth at 1000 8C by introducing an intermediate GaN layer grown at about 850 8C [381,382]. Although the buffer layer of GaN was rough, it became smooth with GaN growth at 1000 8C, indicating this a promising method for the preparation of freestanding GaN substrates. Kumagai et al. [383] found that numerous pinholes formed in the GaN film when grown on a GaAs(1 1 1)B surface and arsenic desorption occurred through them from the GaAs substrate in HVPE. In contrast, a GaN layer with a mirror-like surface was grown on (1 1 1)A substrate. Hong et al. [384] also reported that a better interface characteristics for GaN grown on GaAs(1 1 1)A than on GaAs(1 1 1)B, where zincblende GaN film was deposited by MOCVD.

6. LiGaO2 Lithium gallate (LiGaO2) is the most closely lattice matched substrate currently being considered for GaN heteroepitaxy, with an average lattice constant mismatch of only 0.9% in the basal plane. Consequently, it is expected that GaN epitaxial layers can be grown on LiGaO2 without using buffer layers. There is evidence that the quality of very thin layers of GaN grown on LiGaO2 is better than films of similar thickness on sapphire or 6H-SiC substrates [385]. LiGaO2 is a polar crystal which may facilitate control of the GaN film polarity. LiGaO2 is easily etched, a beneficial property for transferring the GaN film to another substrate for better heat dissipation, combining with integrated circuits, or for producing thick free-standing GaN epitaxial films for subsequent homoepitaxy. It is relatively inexpensive in comparison to SiC. The main disadvantages to LiGaO2 are its low thermal stability under MOCVD growth conditions, low thermal conductivity, high thermal expansion coefficients, and that it is an electrical insulator (i.e. conductive LiGaO2 substrates are not possible). However, if a thick layer (300 mm) of GaN is grown by HVPE and subsequently the LiGaO2 is removed by etching, then the LiGaO2 properties are irrelevant for future device applications. 6.1. Structure and properties of LiGaO2 LiGaO2 (lithium gallium oxide, LGO) has the space group of Pna21 (no. 33) and is classified as orthorhombic, but the crystal structure of LiGaO2 is very similar to the III-nitride wurtzite structure, as shown in Fig. 30. LiGaO2 is the I–III–V analog of ZnO, with one half of the Zn replaced by Li, and the other half replaced by Ga. Unlike ZnO, LiGaO2 melts under atmospheric pressure, and hence, large single crystals can be pulled from the melt using the Czochralski method. The metal atoms are ordered, alternating in the [0 1 0] direction between Li and Ga layers. The crystal structure deviates slightly from the hexagonal symmetry because of the need of accommodating two different metallic atoms of Ga and Li. However, the main virtue of LiGaO2 is that it has a very small lattice

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Fig. 30. The crystal structures of LiGaO2 and GaN (after [387]).

constant mismatch with GaN (averaging 0.9% at room temperature: 1.9% in the a-direction and
0.19% in the b-direction). The atomic relationship between GaN and LiGaO2 in the ideal case is shown in Fig. 31 [386]. One advantage of LiGaO2 over sapphire is that it has polarity along the c-axis or [0 0 1] direction. Surfaces with either anion (oxygen) or cation (gallium and lithium) termination, respectively, are found on the (0 0 1) substrate. The oxygen terminated surface (A-face) can be etched easily while the metal terminated surface (B-face) is difficult to etch in an aqueous solution of nitride acid (H2O:HNO3 ¼ 1:1) [388,389]. It was speculated that the surface oxygen atoms have a dangling bond with two electrons while surface metal atoms have no dangling bond [389]. Kropewnicki et al. [390] developed a basic etchant with different pH values and the maximum etch rate of 0.25 mm/min to the cation face was obtained at pH 9.2 and temperature 50 8C. In contrast to an acid etch, a basic solution etches the cation side more rapidly than the anion side. An entire LGO substrate was chemically removed in <5 min in a base solution without etching gallium nitride at all [385]. This easy etching character of LiGaO2 is essential to make it an attractive material for compliant substrate technology.

Fig. 31. Example of the exact fit of GaN atoms over the LiGaO2 lattice if there were no distortion (after [386]).

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Table 15 The average linear thermal expansion coefficient of LiGaO2 Reference

Temperature range (K)

a[1 0 0]  10
6 K

a[0 1 0]  10
6 K

a[0 0 1]  10
6 K

Nanamatsu et al. [392] Neumann et al. [393] Ishii et al. [394] Rawn and Chaudhuri [395]

293–473 293–473 293–1073 293–473

6 12.6 1.7 7.0

9 15.7 11.0 15.6

7 7.5 4.0 10.7

a[1 2 0]  10
6 K

3.8

In general, the properties of LGO have not been well quantified. The modulus of LGO is about 150 GPa and hardness about 10.0 GPa as measured by nanoindentation [391]. There is still some discrepancy about the thermal expansion coefficient of LiGaO2. Since LiGaO2 is not symmetric along the [1 0 0] and [0 1 0] directions, the coefficients along these two directions are different. The thermal expansion coefficients as measured by different authors are listed in Table 15. The strong anisotropy may lead to both compressive and tensile strain in the GaN epilayer on LiGaO2 since GaN has the thermal expansion coefficient of 5:59  10
6 K
1. Such anisotropic strain has not been confirmed experimentally. A significant disadvantage of using LiGaO2 directly as a substrate is that the thermal conductivity of LiGaO2 is presumed to be small (values have not been measured), which may also limit its device application. 6.2. Bulk crystal growth of LiGaO2 LiGaO2 melts congruently at 1585 8C, therefore, the requirements for its crystal growth are relatively modest—no HT or high pressure apparatus are necessary. Accordingly, good quality wafers are available in relatively large diameters at a price comparable to sapphire. Crystals up to 50 mm in diameter, and 200 mm long have been produced at pull rates of 2–5 mm/h by the Czochralski method from the mixture of Li2CO3 and Ga2O3 [386]. Crystals have been grown in [1 0 0], [0 1 0], and [0 0 1] orientations. A large boule of LiGaO2 is shown in Fig. 32 [396]. In early studies, bulk LiGaO2 crystals occasionally had polar-twin defects, arising from polarity inversion along the c-axis [397]. In recent years, the size and density of these domains have been greatly reduced [398].

Fig. 32. A large single crystal boule of LiGaO2 pulled from the melt at Crystal Photonics Inc. (after [396]).

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6.3. GaN epitaxy on LiGaO2 substrates The orientation relationship between GaN and LiGaO2 has been studied by several groups, yielding [1 1 2 0] GaN || [0 1 0] LiGaO2 and [0 0 0 1]GaN || [0 0 1]LiGaO2 [388,399–401]. Despite the small lattice constant mismatch between GaN and LiGaO2, there was still a significant density of threading dislocations and stacking faults at the interface [400]. A threading dislocation density of about 6  108 cm
2 was shown by TEM for films grown by MBE, although the GaN/LiGaO2 interface was highly disordered [398,402]. However, this value was still much lower at the interface than that of GaN grown on other substrates. The close lattice matching between GaN and LiGaO2 supports the direct epitaxial growth of GaN films; no LT nucleation layer is required. Generally, the FWHMs of GaN films on LiGaO2 are much lower than that deposited directly on other substrates. For a film thickness of 0.4 mm deposited by MBE, the FWHM value of (0 0 0 2) X-ray peak of GaN thin film is 260 arcsec on B-domain LiGaO2 and 1840 arcsec directly on sapphire [387]. For a 0.28 mm GaN film by MBE, the FWHM value of (0 0 0 4) X-ray peak is 145 arcsec, 0.5 mm film 103 arcsec, and 1.0 mm film 80 arcsec [385,403,404]. Doolittle et al. [403] also reported desorption of Li and Ga atoms degraded the surface crystal quality for temperatures >840 8C under MBE conditions. Because of the small atom size of lithium, it is a fast diffuser in the solid-state, and lithium from the substrate is typically incorporated into the GaN film. In addition, diffusion and reaction between the LiGaO2 and GaN produces a transitional interface layer such as Li5GaO2 that is amorphous or nanocrystalline [398,405]. The results of GaN epitaxy by MOCVD on LiGaO2 have been of mixed success. GaN epitaxy by MOCVD is two-dimensional and smooth on LiGaO2, while it is discontinuous, forming islands and columnar on sapphire by single-step method [406]. Thus, the film quality of thin (<0.5 mm) GaN layers is much better than films of similar thickness grown directly on sapphire. However, MOCVD grown GaN films deposited at 1000 8C pealed off as soon as they were in contact with the water vapor in ambient atmosphere [406]. Moreover, hydrogen attacks LiGaO2, thus, nitrogen must be used as a carrier gas in MOCVD growth of GaN in order to achieve high structural high quality GaN on LiGaO2 [407]. With the controlled low growth temperature (850 8C) and extremely low input partial pressures of hydrogen, Duan et al. [399] successfully deposited GaN on the both domains of LiGaO2 by MOCVD without the problem of peeling off. In addition, the problem of diffusion of lithium into the GaN film can be prevented by substrate nitridation before MOCVD growth, which also leads to the reconstruction of the substrate surface and to the formation of a thin layer with the same orientation as the substrate [408]. In the past, it has been difficult to obtain LiGaO2 crystals that are of single polarity (cation or oxygen terminated only). Gallium nitride epitaxy is only possible on purely cation terminated crystals, as film adhesion on oxygen or mixed terminated crystal is very poor. Doolittl and Brown [388] concluded that GaN films grow epitaxially only on the metal face (B-face) by MBE at 700 8C; the films did not adhere to the oxygen face (A-face). Doolittle et al. [403] also found that GaN films peeled or cracked shortly after MBE growth without proper identification of the cation growth face. Although recently some techniques, such as ECR–MBE [409] or special care with LT and low hydrogen partial pressure in MOCVD [399], have been applied to grow GaN on both domains of substrates, recent improvements in LiGaO2 crystal growth have reduced the problems associated with antiphase domains. Due to the extremely high etching selectivity for GaN/LiGaO2, GaN on LGO system forms the desirable template for a viable thin film/compliant GaN substrate [385,402,410]. The compliant substrate provides a new way for homoepitaxial growth of GaN, although the large thin GaN film still tended to crack due to the strain [402]. In addition, Kryliouk and coworkers [408,411] grown single GaN substrate by HVPE on LiGaO2 without cracks and residual strain. The film thickness was about

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Fig. 33. An SEM image of the rear (first grown) surface of a free-standing GaN layer which had been grown on LiGaO2 by HVPE, with subsequent removal of the LiGaO2 by wet chemical etching (after [412]).

100–300 mm, and the average FWHM value of the (0 0 0 4) X-ray peak ranged from 100 to 300 arcsec. An example of the rear surface of a thick GaN film after removal from the LiGaO2 substrate by wet chemical etching is shown in Fig. 33 [412]. The surface proved to be mirror smooth. Therefore, LiGaO2 shows great promise for allowing the preparation of free-standing GaN wafers for homoepitaxy.

7. AlN Aluminum nitride (AlN) is an ideal material for III-nitride substrate applications. AlN and GaN have the same structure (isomorphic), and their lattice constants and thermal expansion coefficients are very similar, thus, abrupt compositional junctions that are relatively free of strain can be produced. An attractive property of AlN is that it is capable of supporting epitaxy on non-c-plane substrates. It can be wet etched in KOH, and is also electrically insulating—ideal for microwave devices. In addition, the isomorphic character of AlN substrates for GaN epitaxy is beneficial to eliminating SMB, which is inevitable due to the steps on the non-isomorphic substrate even with a similar lattice constant such as SiC. Compared to GaN bulk crystals, AlN is even better for its high thermal conductivity (almost double that of GaN), closer lattice match for high A1-content AlGaN epitaxy (essential for short-wavelength emitters and detectors), and better performance for certain devices such as radio frequency transistors due to its insulating nature. To date, bulk AlN crystals are not available commercially, but intense research to improve its crystal growth suggests that AlN substrates will be available in a few years. 7.1. Structure and properties of AlN AlN normally has the wurtzite structure, although epitaxial layers of zincblende structure AlN have been made [47,413]. Wurtzite AlN has the space group of P63mc (no. 186) as same as wurtzite GaN shown in Fig. 1. The (0 0 0 1) surfaces of AlN are polar, which has an important effect on its etching, bulk crystal growth and GaN epitaxy.

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Table 16 Properties of wurtzite AlN at room temperature Properties

Value

Reference

Lattice constant (bulk crystal, nm) Density (g/cm3) Melting point Heat capacity (J/mol K) Thermal conductivity (bulk crystal) (W/cm K) Thermal expansion coefficient (linear)  10
6 K
1 Percentage change in lattice constants (300 –1400 K) Bulk modulus (GPa) Young’s modulus (GPa) Refractive index Dielectric constant Band gap (bulk crystal, eV)

a ¼ 0.31106, c ¼ 0.49795 3.28 2800 8C at 100 atm 29.83 2.85 a ¼ 2.9, c ¼ 3.4 Da/a0: 0.6415, Dc/c0: 0.5349 201 344.83 2.15  0.5 eð0Þ : 8:5  0:2; eð1Þ : 4:68, 4.84 6.2

[416] [417] [418] [419] [420] [419] [19] [421] [422] [423] [423] [424]

Electrical resistivity (O cm) Undoped n-, p-type

107 to 1013 400, 103 to 105

[425]

AlN combines properties such as high thermal conductivity, low thermal expansion coefficient, high electrical resistivity, good dielectric properties, and excellent oxidation resistance. It is inherently stable under GaN epitaxial growth conditions, both by MOCVD and MBE. The properties of aluminum nitride depend critically on material purity. Because of the strong affinity of aluminum for oxygen and the very negative Gibbs free energy of formation of Al2O3 (
1:58  106 J/mol), oxygen is a very common contaminant in AlN. Oxygen reduces its thermal conductivity and causes point defects and stacking faults. Slack reported the lowest oxygen concentration in AlN as 4:3  1019 cm
3 [414]. The maximum thermal conductivity has been theoretically estimated as 3.20 W/cm K. Lower thermal conductivity values are attributed to the presence of oxygen impurities and secondary phases at the grain boundaries [415]. Table 16 summarizes physical, chemical, thermal, mechanical and optical properties of wurtzite AlN. 7.2. Bulk crystal growth of AlN Bulk AlN crystals have been prepared by sublimation, vaporization, and solution routes [414]. Of these, sublimation has been the most successful method for producing bulk AlN crystals [414,424,426–428]. Fig. 34 shows temperature–pressure diagram for aluminum and nitrogen system according to the thermodynamic data [429]. The nitrogen vapor pressure over AlN is six-orders of magnitude lower than that over GaN, thus, it is possible to grow bulk AlN crystals at atmosphere or subatmosphere pressure. Melt growth of AlN is not practical due to its extremely high melting points, estimated to exceed 2800 8C. Nevertheless, due to the high reactivity of aluminum gas at high sublimation temperatures (about 2000 8C) of aluminum nitride, there are still some problems that need to be solved such as a high purity AlN source and an inert crucible for AlN growth [414]. The largest AlN crystal of 470 mm3 reported by far were achieved by Slack and McNelly [424,430]. The source material was prepared by the direct reaction of aluminum and nitrogen at 1850 8C and the crystal grew at the rate of 0.3 mm/h with free nucleation in the sharp tip of the sealed tungsten crucible at the temperature of 2250 8C in a nitrogen atmosphere. Tungsten was most durable of several materials tested including graphite, SiC, Mo, and Ta [414]. Tanaka et al. [431] prepared AlN crystals freely nucleated on the graphite crucible wall and investigate their axis

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Fig. 34. Temperature–pressure diagram for aluminum and nitrogen system.

orientations. Three crystal shapes were obtained: needles with hexagonal cross-sections; flat platelets; and needles with rectangular cross-sections. The FWHM of (0 0 0 2) rocking curve for the needle-shaped crystal with a rectangular cross-section was as low as 12 arcsec. These crystal shapes were dependent on the growth temperature, and crystal sizes were limited to less than 10 mm3. Currently, the group of Schowalter and coworkers [432–435] continues the most successful AlN growth of Slack and McNelly from the mid 1970s [424,430] and large (15 mm diameter) boules have been prepared with FWHM 100 arcsec for (1 1 2 2) rocking curve and density of dislocations <104 cm
2 [436]. Chemical mechanical polishing method has been applied on those AlN wafers and surface roughness of 1.41.6 nm was achieved. 6H-SiC has been utilized as the seed crystal of AlN sublimation due to the easy nucleation and orientation control. Balkas et al. [437] grew seeded single crystalline platelets of AlN <1 mm thick on a 10 mm  10 mm 6H-SiC in a resistively heated graphite furnace at the nitrogen pressure of 500 Torr. At high growth temperature between 2150 and 2250 8C, the growth rates was as high as 0.5 mm/h for a source to seed separation distance of 1–5 mm. But individual hexagonal AlN crystals ( 2 mm  2 mm) were observed on the 6H-SiC seed, and the incorporation of Si, C and O was confirmed due to the utilization of a silicon carbide coated graphite crucible. In other groups, Sarney et al. [438] also grew AlN on 6H-SiC in the temperature range between 2150 and 2200 8C, with nitrogen pressures of 400–410 Torr and a separation distance 4 mm. The growth rate was varied between 10 and 50 mm/h. But the same problems as those seen by Balkas et al. were also observed, i.e. the cracks and high screw dislocations due to stress from the thermal expansion coefficients mismatch between AlN and SiC. Karpov and coworkers [439,440] grew AlN layers 60–80 mm thick of 10 mm  10 mm size in a sublimation–sandwich system with SiC as the seed under both atmospheric pressure and in vacuum (10
4 Torr). The growth temperature was maintained at 1700– 2300 8C and the separation distance was about 1–5 mm. The maximal growth rate 90 mm/hr was achieved. But the authors did not present any characterizations of these crystals. Edgar and coworkers carried out the sublimation growth in a semi-sealed tungsten crucible at a temperature

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around 1800 8C and pressure 400 Torr, and the transportation in the crucible, nucleation behavior and AlN buffer layer effects have been studied in detail [441–445]. They identified three major problems with seeding AlN on SiC faces: (1) the thermal expansion coefficient of AlN exceeds that of SiC, while its critical shear stress is lower. Thus, the AlN films are highly stressed and frequently cracked; (2) SiC is unstable and decomposes at the HT necessary for AlN crystal growth. Furthermore, Si and C from the SiC seed can incorporate into the growing AlN film; (3) AlN sublimed onto SiC does not completely wet the surface, but proceeds by three-dimensional nucleation. This creates subgrain boundaries in the resulting AlN film. J. Nause of CERMET Inc. tried to grow AlN directly from a melt [446]. This high pressure, RF, melt growth process will solve the problems of both container reaction and impurity inclusions. The process will reach the temperature of 2600 8C and pressure of 10–100 atm. The AlN crystals of cmsized areas have been prepared with high quality (FWHM 27 arcsec). Schlesser and Sitar [447] recently reported using vaporization method [448] to grow AlN bulk crystals. Pure aluminum and nitrogen were used as the source materials. These are inherently higher purity sources than commercial AlN, which typically contains on the order of 1.0 wt.% oxygen. Higher growth rates are also possible due to the higher vapor pressure of aluminum over pure aluminum compared to AlN. At the temperature of 2100 8C, AlN c-plates up to 50 mm2 with large growth rates of 5 mm/h in the c-plane and 0.2 mm/h along the c-axis were obtained for 2 h’ run. Raman spectroscopy indicated crystal quality increased drastically with increasing temperature over the temperature range studied, 1800–2100 8C. The crystal morphology changed from predominately needles to platelets as the temperature was increased. Some new techniques for AlN growth, such as HVPE [449] and metal nitridation in supercritical ammonia methods [450,451] have also been explored recently. For HVPE growth, the crystal quality is worse than the one by the sublimation method. The latter method still produced too small crystals (the largest crystals obtained were only 3 mm thick and 25 mm long in [450]) and need further improvement. It is worth pointing out that there is a competition for the first commercialization of AlN bulk crystals and GaN bulk crystals. The success of bulk SiC by sublimation stimulated the growth of AlN by sublimation in 1990s, and suggested its growth is more feasible since GaN bulk crystals can not be grown by sublimation due to the extremely high dissociation pressure of GaN. There is still of course a great deal of interest in GaN bulk growth. GaN bulk crystal growth seemed likely than AlN growth because GaN can be grown in melt. Another direction of GaN bulk growth by HVPE is also beneficial from the extensive experience of GaN epitaxy. More attention was focused on GaN growth since GaN has a larger commercial market. 7.3. GaN epitaxy on AlN substrates The only research for GaN epitaxy on an AlN substrate was carried out by the group of Schowalter and coworkers [435–443]. With the grown single crystals of AlN up to 15 mm in diameter, the substrates were cut from these single crystals boules with both the (1 1 2 0) (A-face) and the (0 0 0 1) (C-face) orientation. High quality AlN and AlxGa1
xN epitaxial layers were deposited on these substrates by MOCVD. Ion channeling along the [1 1 2 0] axis gave a channeling minimum yield of 1.5% for an AlN layer and 2.2% for an Al0.5Ga0.5N layer indicating an extremely high quality epitaxial layer grown on the A-face AlN substrate. AFM characterization showed that the AlN epilayer has increased surface roughness compared with the substrate and the Al0.5Ga0.5N epilayer surface was faceted with individual facet heights reaching almost half the film thickness (1 mm film thickness). AFM images of the AlN substrate, AlN epilayer, and Al0.5Ga0.5N

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Fig. 35. The AFM images of: (a) the AlN substrate after chemical mechanical polishing; (b) after growth of a 0.7 mm-thick AlN epilayer; (c) after growth of a 1 mm-thick Al0.5Ga0.5N epilayer. Scan area, 1 mm2; height scales, 1.5, 10, and 400 nm, respectively (after [434]).

epilayer are shown in Fig. 35 [434]. Double crystal X-ray diffraction with FWHM of 17 arcmin for (1 1 2 0) rocking curve, and preliminary electrical results as resistivity of 20 O cm and a mobility of 20 cm2/V s reported for Si-doped, 1 mm-thick Al0.5Ga0.5N epilayer, demonstrated the potential of AlN substrates.

8. GaN substrates Gallium nitride itself is the best choice as a substrate for GaN epitaxy and device fabrication, as it eliminates all problems associated with heteroepitaxy. Homoepitaxy offers better control compared to heteroepitaxy over the crystal polarity, dopant concentration, and strained layer thickness. Gallium nitride homoepitaxy of smooth films in a two-dimensional growth mode does not require nitridation or buffer layers, as is required for sapphire and silicon carbide. Although there are several techniques under development for producing bulk GaN crystals including growth by vapor phase transport [452,453], growth from supercritical fluids [454], and growth from sodium fluxes [455], only high pressure growth from solutions and hydride vapor phase epitaxy have produced large area crystals, which will be discussed in the later sections. 8.1. HNPS GaN substrates The most extensive studies of GaN homoepitaxy have employed bulk crystals produced at the High Pressure Research Center in Warsaw, Poland. Grzegory and coworkers produce bulk crystals from gallium melts saturated with 1 at.% nitrogen at temperatures up to 1700 8C and nitrogen pressures of 20,000 atm, named the high nitrogen pressure solution (HNPS) growth technique [456–458]. The properties of these GaN crystals are summarized in Table 17. The structural properties of the bulk crystals are excellent with very narrow X-ray rocking curve widths for all crystals planes, low dislocation densities, and negligible strain. The optical and electrical properties of the bulk crystals are poor compared to heteroepitaxial GaN layers on sapphire and other substrates, a consequence of high residual impurities and point defect concentrations (believed to be primarily Ga vacancies). Both conductive and semi-insulating bulk GaN crystals have been used as substrates for GaN homoepitaxy. Unintentionally, doped GaN produced by the HNPS method typically has a high electron

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113

Table 17 Properties of bulk GaN crystal produced by HNPS [456,459,460] Property of bulk GaN crystals produced by HNPS

Undoped GaN

Mg doped GaN

(0 0 2) XRD rocking curve width (arcsec) Resistivity (O cm) Carrier concentration (cm
3) Mobility (cm2/V s) Oxygen concentration (cm
3) Ga vacancy concentration (cm
3) Dislocation density (cm
2)

30–40 10
3 to 10
2 3  1019 to 6  1019 30–90 1019 to 1020 2  1017 10 to 105

20–30 104 to 106 1019 to 1020 6  1016

concentrations, greater than 1019 cm
3, from high oxygen concentrations in the crystals. The high electron concentration causes a slight dilation in the c-lattice constant in comparison to low electron concentration material. Consequently, homoepitaxial GaN layers with low electron carrier concentrations are slightly strained when deposited on such high electron concentration substrates [461]. The donor activity of the oxygen can be compensated by adding magnesium, making the bulk crystals semi-insulating [459]. Adding magnesium also reduces the gallium vacancy concentration (VGa) and the crystal lattice constants [460]. The XRD FWHM from all crystal planes becomes narrower indicating that magnesium doped bulk GaN crystal have a better structural quality than undoped bulk crystals [456]. The Ga and N polar faces of bulk GaN crystals have distinctive chemical properties which impact substrate preparation, and the structural, electrical, and optical properties of the homoepitaxial films [462]. For as-grown crystals, the Ga-face is atomically smooth and chemically inert. The N-face is relatively rough, with pyramids several tens of nanometers in height. For crystals produced by the HNPS method, the defect density is much higher near the N-face than the Ga-face [462]. N-polarity (0 0 0 1) surfaces can be made atomically flat by chemo-mechanical polishing, but gallium polarity surfaces are chemically inert, and can only be mechanically polished. Consequently, removing all remnant polishing damage from Ga-face crystals is more difficult. In contrast to the inferior electrical and optical properties of the substrates, for homoepitaxial GaN layers on these substrates these properties are the best reported for any GaN including crystal quality, photoluminescence, and electrical properties [463–465]. The photoluminescence peaks from the homoepitaxial layers can be extremely narrow, less than 0.1 meV, indicating very high quality material [464]. The result from epitaxial films grown by MOCVD and MBE are summarized in Table 18. Because AlxGa1
xN and InxGa1
xN epitaxial layers grown on GaN substrates can be defect-free, thicker layers can be produced without stress-relaxation, than on similar layers on sapphire or silicon carbide [467,468]. Table 18 Summary of photoluminescence results for homoepitaxial layers on HNPS substratesa PL peak identity

LT peak position (eV)

Peak FWHM (meV)

Polarity

Growth technique

Reference

DBE ABE DBE DBE DBE DBE

3.4709 3.4655 3.4714 3.4714 3.4712 3.4720

0.114 0.095 0.35 0.9

Ga

MOCVD

Ga N Ga

MOCVD

[464] [464] [465] [465] [466] [466]

a

0.3

DBE, exciton bound to a neutral donor; ABE, exciton bound to a neutral acceptor.

MBE

114

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The impurity concentration and the composition of epitaxial alloy films are dependent on the substrate polarity. Prystawko et al. [469] demonstrated that donor impurities, such as silicon and oxygen, more readily incorporate to higher concentrations in films deposited on N-polarity crystals than on Ga-polarity crystals. On the other hand, acceptors are incorporated at higher concentrations on Ga-polarity crystals than on N-polarity crystals in the same growth run. Similarly, the composition of AlGaN and InGaN films is affected by which polarity the film is deposited on. Prystawko et al. [468] observed that the Al content was higher for AlGaN deposited on Ga-polarity (0 0 0 1) substrates compared to N-polarity substrates, while the In content was higher on the N-polarity substrates for films deposited under the same process conditions by MOCVD. Misorientation of the substrate affects the growth mode and morphology of the homoepitaxial films. On crystals exactly oriented to the (0 0 0 1) plane, deposited films were smooth on Ga-polarity crystals, but were rough with a high density of hexagonal hillocks on N-polarity crystals [470]. Zauner et al. [470] attributed the hillocks to inversion domains of Ga-polarity which grow at a faster rate than the surrounding N-polarity material. The inversion domains nucleate at thin platelets containing oxygen, possibly produced during chemical–mechanical polishing. The hexagonal hillock density was reduced by a factor of 100 by using N-polarity GaN crystals with intentionally misoriented by 48 off the (0 0 0 1) plane. The direction of the misorientation was important—as crystals tilted toward the [1 1 2 0] direction had narrower photoluminescence peaks than crystals tilted toward the [1 0 1 0] direction. To date, only a few devices have been fabricated on HNPS GaN. Pelzmann et al. [471] made homoepitaxial blue LEDs which were twice as bright as heteroepitaxial devices. Mayer et al. [472] demonstrated UV LEDs emitting at 371 nm with a line width of 8 nm. Khan et al. [473] found HFET on HNPS GaN substrates had slightly better linearity than similar devices on SiC, but was more power-limited by self-heating effects, due the lower thermal conductivity of GaN compared to SiC. Although the HNPS growth technique produces excellent bulk GaN crystals, it is a difficult process to scale up to produce large crystals at high throughput rates due to the high pressure required and the low maximum growth rates (0.1 mm/h or less). Current plans are to scale this process to produce 25 mm diameter crystals within a year. 8.2. Free-standing GaN epitaxial layers produced by HVPE Thick free-standing epitaxial GaN films deposited by HVPE and separated from the substrate is another approach for producing substrates for GaN homoepitaxy [474–476]. The equipment required for HVPE is relatively simple, and the growth rates are comparably high, up to 100 mm/h, making this method economically viable. Much larger area GaN crystals have been produced by HVPE, up to 50 mm in diameter and 300 mm thick, compared to HNPS GaN crystals [476]. A variety of substrates have been employed as templates for supporting the GaN film including 6H-SiC [477], sapphire [474–476], LGO [410,411], NDGaO3 [478], spinel (MgAl2O4) [44], GaAs [479], and Si [480]. The substrate is subsequently removed and the free-standing GaN film is polished and etched to prepare it for epitaxy. Once removed from the template substrate, both the Ga- and N-polarity surfaces are accessible. Such free-standing epitaxial films have the same problems always characteristic of heteroepitaxy as listed in Table 19, but to a lesser degree because the film is so thick: the dislocation densities and optical properties improve with film thickness [481]. The structural quality of the HVPE substrates is not as good as HNPS GaN crystals, as indicated by wider XRD rocking curve FWHM and higher dislocation densities, but the electrical and optical properties are better, as the residual background

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115

Table 19 Summary of the best reported properties for HVPE GaN substrates Properties of free-standing GaN crystals produced by HVPE

Value

Reference

(0 0 2) XRD rocking curve width (arcsec) Maximum resistivity with Zn doping (O cm) Minimum carrier concentration (cm
3) Maximum room temperature mobility (cm2/V s) Dislocation density (cm
2) LT PL peak position, Ga-face (DBE) (eV) LT PL peak FWHM, Ga-face (DBE) (meV) RMS surface roughness, Ga-face (nm) RMS surface roughness, N-face (nm)

69 1  1012 8  1015 903 5  105 3.4710 1.1 0.4 1

[482] [483] [482] [484] [482] [482] [482] [485] [482]

carrier concentrations can be much lower (<1016 cm
3) and the electron mobility higher [482]. The low residual impurity concentrations makes control of the electrical properties by the addition of impurities more reliable. Resistivities can vary widely from semi-insulating for compensated GaN produced by Zn doping [483], to n-type conductive substrates produced by doping with Si. However, the best structural quality material is nominally undoped; doping tends to degrade the crystal’s quality. Several challenges remain to be solved with HVPE GaN substrates. Substrates are typically bowed, a consequence of the thermal expansion mismatch between the GaN film and the substrate. Large-scale surface morphologies may be present on the as-grown GaN surface. Consequently, substrates must be polished to planarize the substrate and to produce a flat smooth surface. As with HNPS GaN, polishing processes for both surfaces of both polarities have only begun to be optimized. Homoepitaxial films deposited on such substrates by MOCVD or MBE can have better properties than the substrate itself. MOCVD GaN deposited on a free-standing HVPE GaN substrate by Miskys et al. [486] had much narrower LT photoluminescence line widths than similar layers deposited directly on sapphire substrates. Later results by this same group demonstrated improved properties for the MOCVD homoepitaxial layer, and are summarized in Table 20 [485]. Combining the low dislocation density of a thick HVPE GaN layer (still on the sapphire substrate in this case) with MBE’s ability to produce abrupt compositional junctions, Manfra et al. [487] produced very high density and high mobility two dimensional electron gases with AlGaN/GaN structures. Table 20 Properties of an MOCVD homoepitaxial GaN films on a free-standing HVPE GaN substrate [485] Properties of homoepitaxial GaN films on HVPE GaN

Value

Layer thickness (mm) Polarity ( 0 0 2) XRD FWHM 2y
o (arcsec) Dislocation density (cm
2) Carrier concentration (cm
3) Room temperature electron mobility (cm2/V s) Dislocation density (cm
2) LT photoluminescence peak position (DBE) (eV) LT photoluminescence (DBE) FWHM (meV) RMS surface roughness (nm)

2 Ga 17 2  107 5  1016 310 2.0  107 3.474 0.55 0.2

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9. Summary and perspective The list of materials which have been successfully employed as substrates for GaN epitaxy continues to grow. For most of these materials, it is questionable whether the potential advantages justify the enormous effort which will be required to develop a viable process to produce good quality films. The best processes for GaN epitaxy on sapphire slowly evolved over many years as there were no competitive substrates. In contrast, any new substrate material must compete with the well-established technologies of sapphire and silicon carbide. The new substrate materials will have to demonstrate clear advantages or unique capabilities over sapphire and silicon carbide. The well-established process technology for GaN epitaxy on sapphire and silicon carbide ensures that these will remain the most common substrates for the foreseeable future. Sapphire will probably be the substrate of choice for low cost and low power devices, while silicon carbide will be preferred for high power devices requiring good heat dissipation. Incremental improvements in crystal quality of gallium nitride film quality on both substrates will continue to be made as epitaxial growth processes are refined. Several new processes utilizing well-developed substrates, which would have merit for GaNbased devices, are currently under intense study to address the general problems of heteroepitaxy. Both compliant [488,489] and porous substrates [490–494] can greatly reduce or entirely eliminate stress in epitaxial films, thereby reducing its dislocation density. A compliant substrate consist of a thick mechanical host substrate (MHS), a layer which is viscous at the epitaxial growth temperature such as SiO2 or a low melting temperature metal, and a thin single crystal layer providing a template for epitaxial growth. Stress is confined to the template layer, allowing the epitaxial layer to grow defect-free. A porous substrate contains a high density of nanometer size holes in the top surface of the substrate. A continuous epitaxial film grows from solid regions of the substrate covering the pores. Stress is reduced in the epitaxial film due to the discontinuous nature of the substrate. Lateral epitaxial overgrowth and its many variations [79,495] will continue to be advanced and simplified, improving the quality of GaN films on inexpensive, large area substrates, and reducing the process costs. For all techniques, it remains to be seen whether the considerable process complexity necessary to prepare the substrate is justified by improvements in material quality or reductions in the cost of devices. In the long-term, GaN substrates will probably be developed as there are very definite incentives for homoepitaxy. Gallium nitride would meet most substrate requirements for zero or low thermal expansion and lattice constant mismatch. It could be made either electrically conductive or semiinsulating in a variety of crystal orientations and polarities. Currently, there are several groups seeking to commercialize bulk GaN crystals by a variety of growth techniques including high pressure solution, sublimation, and the ammonothermal methods. The information available on these commercial ventures (to data, generally very little) has been reviewed in the section on bulk GaN growth. Already some bulk GaN crystals and free-standing epitaxial HVPE GaN crystals are commercially available, and impressive results for homoepitaxy have been demonstrated. Initially, the cost of GaN substrates will be quite high, but costs will drop as simple crystal growth technique and the associated process technology are developed. Should bulk GaN wafers become available, future reviews of substrates for GaN epitaxy may be restricted to only GaN itself.

Acknowledgements Support for this work from the Office of Naval Research and the National Science Foundation is greatly appreciated.

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