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Structural and morphological evolution of GaN grown by metalorganic chemical vapor deposition on SiC substrates using an AlN initial layer
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Journal of Crystal Growth 273 (2004) 38–47 www.elsevier.com/locate/jcrysgro
Structural and morphological evolution of GaN grown by metalorganic chemical vapor deposition on SiC substrates using an AlN initial layer B. Morana,, F. Wua,b, A.E. Romanova,1, U.K. Mishrac, S.P. Denbaarsa,b,c, J.S. Specka,b a Materials Department, University of California, Santa Barbara, CA 93106-5050, USA JST/ERATO, UCSB Group, University of California, Santa Barbara, CA 93106-5050, USA c Electrical and Computer Engineering Department, University of California, Santa Barbara, CA 93106-5050, USA b
Received 5 February 2004; accepted 16 August 2004 Communicated by C.R. Abernathy Available online 7 October 2004
Abstract The morphological and structural evolution is presented for GaN grown by metalorganic chemical vapor deposition on 25 nm thick or 150 nm thick AlN initial layers on (0 0 0 1) 4 H and 6 H SiC substrates. The 25 nm thick AlN on SiC was a rough, partially coalesced film, whereas the 150 nm thick AlN on SiC was smooth and was characterized by a step-terrace structure. Both the 25 and 150 nm thick AlN layers on SiC were nearly free of elastic strains. For both AlN initial layers, the GaN films grew by a coarse islanding mechanism. Measurement of the GaN (0 0 0 2) interplanar spacing shows that these islands were largely strain relaxed throughout the growth process. Plan view transmission electron microscopy (TEM) showed a well developed misfit dislocation network at the GaN island/AlN interface. Cross sectional TEM revealed that these islands are free of threading dislocations (TDs) prior to coalescence. A simple, lowangle grain boundary model based on island misorientation at coalescence provides reasonable agreement with TD densities measured using plan view TEM. An electron mobility value as high as 818 cm2/V s has been measured at room temperature confirming a high degree of material quality. r 2004 Elsevier B.V. All rights reserved. PACS: 61.72.y Keywords: A1. Line defects; A3. Metalorganic chemical vapor deposition; B1. Nitrides
Corresponding author. Tel.: +1-805-8938869; fax: +1-805-
893-5263. E-mail address: [email protected] (B. Moran). 1 Permanent address: Ioffe Physico-Technical Institute, St. Petersburg, Russia.
1. Introduction The lack of a native substrate has focused research effort in the wide band gap nitrides into
0022-0248/$ - see front matter r 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.jcrysgro.2004.08.012
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the hetereoepitaxial growth on a variety of different substrates. The in-plane lattice mismatch of 16% for GaN and sapphire and 3.4% for GaN and SiC leads to the formation of a high density of threading dislocations (TDs) in the epitaxial GaN layers (108–1011 cm2), the presence of which have been shown to be detrimental to the performance of both optoelectronic and electronic devices [1,2]. High quality, c-plane epitaxial GaN thin films were produced on sapphire substrates using MOCVD by employing a two-step growth technique [3,4]. This method produced a Volmer–Weber growth mode, in which TDs form along low-angle grain boundaries of coalescing nuclei [5,6]. In recent years, GaN-based optoelectronic devices grown in this manner have become commercially available from many manufacturers. High power, high-frequency GaN-based electronic devices are a comparatively undeveloped field, though significant development activities have resulted in promising reports of high-performance GaN-based transistors [7,8]. Due to self-heating effects, GaN-based transistors require significant thermal dissipation; it has been demonstrated that SiC (thermal conductivity, k3.4 W/cm K at 300 K) [9] is clearly a more suitable substrate than sapphire (k0.42 W/cm K at 300 K) [10] for these applications [11]. There are reports of the growth of GaN on SiC without the use of any intermediate buffer layer. Sakaski and Matsuoka successfully deposited GaN directly on SiC by MOCVD [12], followed by a few groups in the 1990s [13,14]. However, as demonstrated by Lahreche et al. [15], the growth of GaN directly on SiC under typical growth conditions results in a rough, islanded GaN layer after long growth time due to poor surface wetting. The use of an AlN initial layer promotes surface wetting and reduces the lattice mismatch (2.4%) between AlN and GaN. Reports of high-performance AlGaN/GaN HEMT devices grown by MOCVD on SiC employ AlN initial layers [16–18]. For instance, Shealy et al. [18] describe an AlN initial layer as an important step for realizing highperformance HEMTs. An early report by Weeks et al. [19] suggested that the growth temperature of AlN in the range of 500–1050 1C resulted in a polycrystalline AlN layer leading to poor quality
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in subsequently deposited GaN; a high-temperature (HT) initial layer of AlN grown at temperatures greater than 1100 1C was used to produce monocrystalline AlN and therefore resulted in higher quality GaN films. Davis et al. [20] provided evidence that GaN deposition on AlN results predominantly in the formation of GaN islands which eventually coalesce and form a continuous film. Previous reports have identified a large density (1011 cm2) of dislocations in the GaN/AlN interfacial region, primarily composed of half loops (Burgers vector b=7[0 0 0 1]) and both mixed character (b ¼ 13h1 1 2 3i) and pure edge-type dislocations (b ¼ 13h1 1 2 0i). The TD density tended to decrease to 109 cm2 at the film surface [21,22] through dislocation reaction processes after 1 mm of GaN growth. By using room temperature electron mobility as a metric, GaN of very high quality has been produced using a HT AlN initial layer on SiC as room temperature electron mobility values as high as 900 cm2/V s have been reported [23]. Despite the body of work on GaN on SiC growth, details of the GaN growth mode and extended defect generation remain relatively undeveloped. The work described in this paper uses growth interruption experiments for two representative III-nitride nucleation conditions on SiC in an MOCVD growth environment. Characterization of these films at various stages of evolution allows for an identification of the growth mode, provides insight into the relaxation of lattice mismatch, and presents evidence for the mechanism of TD generation.
2. Experimental procedure Epitaxial AlN and GaN films were grown by MOCVD at 1050 1C in a Thomas Swan closecoupled showerhead reactor on 4 H and 6 H SiC substrates. The substrates were baked in H2 for 10 min at growth temperature prior to film deposition. AlN was grown at 76 Torr, using trimethylaluminum (TMAl) and NH3 at a growth rate of 0.75 A˚/s. GaN was grown at 760 Torr using trimethylgallium (TMGa) and NH3 at an
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approximate growth rate of 3.5 A˚/s. H2 was used as a carrier gas in all cases. Where applicable, silicon doping of GaN was carried out using disilane (Si2H6) as a precursor. AlN layers of 25 and 150 nm in thickness were first deposited on the SiC substrate prior to GaN growth to generate two comparative sample sets. These thicknesses were chosen as they generated two disparate GaN island nucleation scenarios. GaN films grown subsequently on these layers will be known respectively as ‘Material 25’ and ‘Material 150’. Their morphological and structural evolution was observed by performing growth interruption experiments similar to the work of Fini et al. [6]. Growth times of 50, 200, and 600 s were selected to observe the initial stages of GaN growth for both Material 25 and 150 prior to film coalescence. Cross sectional and plan view specimens for transmission electron microscopy (TEM) were prepared by wedge polishing and argon ion milling. Scattering contrast TEM images were recorded with either a JEOL 2000 FX or an FEI Tecnai G2 Sphera; high-resolution TEM (HRTEM) imaging was performed on a JEOL 2010HR. The operating voltage for all TEM imaging was 200 kV. Scanning electron microscopy (SEM) was performed using an FEI Sirion operating at 2 kV. Atomic force microscopy (AFM) was conducted using a Digital Instruments D-3000 SPM in tapping mode. X-ray diffraction scans were recorded using a Philips X-Pert MRD in triple-axis configuration. The (0 0 0 l) interplanar spacing for each evolution sample was determined by measuring the relative separation of the GaN (0 0 0 2) peak from the SiC (0004) or (0006) peak (4H or 6H SiC substrates, respectively) in an o–2y scan. Rocking curves (o-scans) of the GaN ð2 0 2 1Þ peak were performed in a skew geometry (realized by tilting the sample by 751) to determine the twist mosaic in the samples. An X-ray reciprocal space map was recorded of Material 150 near the GaN ð1 0 1 5Þ reflection, with the ensuing room temperature a lattice parameter of AlN being estimated using the ð1 0 1 10Þ reflection of 4H SiC as an angular reference. Electron mobility was measured using indium metal contacts and standard van der Pauw
geometry over 5 mm 5 mm.
a
sample
dimension
of
3. Results Particular consideration was given to the resultant AlN surface morphology prior to GaN deposition. The 25 nm thick AlN layer, shown in Fig. 1a, was a partially coalesced, islanded film with an average grain size of 30 nm. Though the tops of the AlN appeared to be flat, the overall RMS roughness of the layer was 1.6 nm. The AFM image in Fig. 1b shows that 150 nm thick AlN film was fully coalesced and had a surface consisting of atomic steps and terraces with an RMS roughness of 0.3 nm. The TD density in the 150 nm thick AlN, as measured by plan view TEM, was 4 1010 cm2. Details of the morphological evolution, strain relaxation, and microstructure of these AlN films on SiC will be reported elsewhere [24]. One micrometer thick GaN films were grown under identical conditions on both 25 or 150 nm AlN thick initial layers. AFM images shown in Fig. 1c and d of the GaN layers (Materials 25 and 150, respectively) demonstrate that both GaN films had a smooth surface morphology with easily identified atomic steps and terraces. A measured step termination density for Materials 25 and 150 was 4 108 and 1 108 cm2, respectively. The total TD density in Material 25 and 150, quantified using plan view TEM, was 1 109 and 2 108 cm2, respectively. Crosssectional TEM images of these layers with imaging conditions of g ¼ 0002 and g ¼ 1 1 2 0 (not shown) were compared to determine the character of observed TDs. Though an exhaustive statistical analysis was not performed, examination of these images indicated that a majority of the TDs were of pure edge-type (b ¼ 13h1 1 2 0i) and the balance had mixed character (b ¼ 13h1 1 2 3i). o rocking curves determined by X-ray diffraction for the onaxis (0 0 0 2) peak had full width half maximum (FWHM) values of 179 and 153 arcsec and off-axis ð2 0 2 1Þ peak widths of 260 and 218 arcsec for Material 25 and 150, respectively.
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Fig. 1. AFM height images (2 2 mm) of (a) 25 nm thick and (b) 150 nm thick AlN and 5 5 mm AFM height images of (c) 1 mm thick Material 25 and (d) 1 mm thick Material 150.
AFM micrographs of the growth interruption series of both Material 25 and 150 are shown in Fig. 2 to provide morphological and geometric detail of the island nucleation and growth process. Each island had primarily a flat (0 0 0 1) surface, which was sustained through film coalescence. SEM images (not shown) showed that the islands were truncated hexagonal pyramids with f1 0 1 1g inclined facets. Table 1 provides the measured GaN island aspect ratio and density. The density is reported for 50 s of growth in for both Materials 25 and 150 as an approximation of the GaN island density because there was no morphological evidence for coalescence in either material at this early stage of growth. Significant coalescence was first observed in Material 25 after 200 s of growth, and the film was nearly continuous after 600 s of growth. Comparatively, clustering was first widely observed at 600 s in Material 150, with complete film coalescence occurring beyond 1200 s of
growth. Analysis several of AFM scans showed that the average island diameter at coalescence was 0.9 mm for Material 25 and 2.5 mm for Material 150. Diffraction contrast TEM studies were performed on representative cross-section samples to reveal the microstructure evolution. An isolated island of Material 25 after 50 s of GaN growth, shown in Fig. 3, was found to be free of TDs. The GaN does show significant strain contrast in both the g ¼ 0 0 0 2 and g ¼ 1 1 2 0 image—this contrast is attributed to both the strain fields of misfit dislocations (MDs) at the GaN/AlN interface and the rough nature of this interface. The contrast from the SiC at the interface appears to correspond to the island structure of the AlN. Fig. 4a shows two coalesced GaN islands of Material 150. The g ¼ 0 0 0 2 image has little contrast—consistent with a smooth GaN/AlN interface. The g ¼ 1 1 2 0 image shows the high density of pure edge
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Fig. 2. AFM images (15 15 mm) of GaN growth interruption study. Material 25: (a) and (b) are height images, (c) is a tip amplitude image. Material 150: (a) is a height image, (e) and (f) are tip amplitude images.
Table 1 Measured average values for GaN island aspect ratio and density Growth time (s)
Material 25 Aspect ratio (width/height)
50 200 600
16.3 9.8 Coalesced film
TDs (b ¼ 13h1 1 2 0i) in the AlN and also shows a segment of a single TD—we believe that this TD formed was a result of island–island coalescence. In the case of Material 25 in a late stage of coalescence (600 s, Fig. 4b), TDs of both pure edge and mixed character were unambiguously observed. Plan view TEM micrographs of 1 mm thick fully coalesced films are shown in Fig. 5. The TD density of Material 25 and 150 was 1 109 and 2 108 cm2.
Material 150 Density (#/cm2) 7
8.1 10 3.6 107 Coalesced film
Aspect ratio
Density
6.1 6.6 8.5
8.4 106 8 106 8 106
An X-ray reciprocal space map was recorded near the GaN ð1 0 1 5Þ reflection of a 1 mm thick Material 150 layer and is shown in Fig. 6. The estimated a lattice constant calculated from this data of 3.1143 A˚ for the 150 nm thick AlN was within 0.01% of its relaxed bulk value [19]. Given the similar a-axis thermal expansion coefficient of AlN and SiC (4.2 106 cm2) [21], it is assumed that the 150 nm thick AlN was nearly fully relaxed at growth temperature.
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Fig. 3. Bright field cross-section TEM image of single Material 25 GaN island (after 50 s of growth) in g ¼ 0002 (top) and g ¼ 1 1 2 0 (bottom) two beam imaging conditions.
Fig. 4. Bright field cross-section TEM images recorded with either g ¼ 0002 or g ¼ 1 1 2 0 imaging conditions of (a) Material 150 after 200 s of growth and (b) Material 25 after 600 s of growth. The arrow in the GaN island in (a) indicates a TD segment.
o–2y X-ray diffraction scans were recorded to determine the GaN (0 0 0 2) absolute peak separation from the relevant SiC substrate peak ((0 0 0 4) for 4H SiC and (0 0 0 6) for 6H SiC). Based upon this data, the (000 l) interplanar spacing and thus
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the measured c lattice constant can be determined for the GaN films at the various stages of film evolution. The value for the a lattice constant and thus the in-plane strain in each sample was determined assuming purely biaxial stress using elastic constants as calculated by Kim et al. [25] and bulk unstrained lattice constant ( values published in the literature (a ¼ 3:1881 A, ( c ¼ 5:1844 A—these lattice constants were used due to the undoped and low background impurity character of the GaN films produced in this study) [26]. It is important to understand the strain in the GaN film at the growth temperature (Tg). Without the availability of in situ monitoring, the strain in the GaN film was estimated by using the in-plane strain at room temperature and then accounting for the thermal strain developed during postgrowth cool down by the differences in the temperature-dependent coefficients of thermal expansion (CTE), Da, between GaN and SiC [27,28]. Fig. 7 shows the in-plane strain data for Materials 25 and 150 at RT and extrapolated to Tg. There are several points to be made regarding the data in this graph. Both of the thick films (i.e., long growth times) of Material 25 and 150 were under tensile strain at RT; after extrapolation to Tg, these films were nearly strain-free. Regarding the Tg data only, there was also no particular point at which coalescence events appeared to markedly change the strain state of the films; rather, a more gradual trend of relaxation of compressive stress was observed as the growth time increased. Finally, the strain in the GaN extrapolated to Tg even at a short growth time of 50 s was low, with values of 2 103 (0.2%) and 8 104 (0.08%) for Materials 25 and 150, respectively. These low residual strain values are indicative of a high degree of relaxation of the 2.4% compressive lattice mismatch strain between GaN and AlN. TEM studies of both Materials 25 and 150 show MDs at the GaN/AlN interface with a periodicity consistent with largely relaxed GaN throughout the growth. Fig. 8 shows a plan view weak beam image of an uncoalesced Material 150 film. The 150 nm thick AlN layer has a TD density of 4 1010 cm2, as determined from regions without GaN coverage. A MD network was clearly
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Fig. 5. Plan view, bright field g ¼ 1 1 2 0 TEM images of (a) 1 mm thick Material 25 (b) 1 mm thick Material 150.
Fig. 7. Measured room temperature in-plane strain and the inplane strain extrapolated to the growth temperature for each sample in the growth evolution study.
Fig. 6. Reciprocal space map of Material 150 (GaN thickness is 1 mm). The map was taken near the asymmetric (1 0 1 5) GaN reflection in coplanar geometry with shallow incidence and rocking curve configuration. ‘‘r.l.u’’ refers to dimensionless reciprocal lattice units (lq 10000).
seen at the GaN/AlN interface. From detailed analysis of the images, the average MD line direction was determined to be h1 1 2 0i: The average MD spacing in the h1 1 2 0i directions was 19 nm for the image shown in Fig. 8. As we will discuss below, the 2.4% mismatch between GaN and AlN can be fully accommodated by a triangular MD array with h1 1 2 0i line direction and Burgers vector 13h1 1 2 0i with MD spacing of 17.3 nm. Thus, the MD spacing of 19 nm is
consistent with 90% relaxation in the GaN islands. Based on our high-resolution imaging (not shown) and diffraction contrast studies, we tentatively conclude that the MD dislocations have Burgers vector 13h1 1 2 0i (further details on the nature of the dislocations at the GaN/AlN interface will be presented elsewhere [24]). MDs were also observed by inclined view weak beam images on cross-section samples at the GaN/AlN interface both for coalesced films and isolated islands. A silicon-doped GaN layer grown on top of Material 25 and 150 template layers provided an appreciable number of carriers to examine the transport properties of these films. Bulk GaN
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Fig. 9. The dependence of electron mobility on measured carrier concentration for several samples of Material 25 and 150.
electron mobility values for several 1.5–2.0 mm thick samples are plotted v s. carrier concentration in Fig. 9. The maximum measured values were 804 and 818 cm2/V s for carrier concentrations of 5 1016 cm3 for Materials 25 and 150, respectively. In both Material 25 and 150, we observed this expected fall-off in mobility with an increase in Si doping. Several groups have shown that scattering due to charged dislocation lines does not limit the transverse room temperature electron mobility for GaN with TD densities in the range 108–109 cm2 and free electron concentrations in the range of 1016–1017 cm3 [29–31]. Thus, it is not surprising that Materials 25 and 150 have a similar room temperature electron mobility despite having nearly an order of magnitude different TDD.
system. It is clear from the AFM data that the GaN on AlN grew via a coarse islanding mechanism. The X-ray diffraction data indicates that the AlN layers were largely relaxed on the SiC. In turn, the GaN was nearly fully relaxed throughout its growth on AlN. Additionally, the low strain when extrapolated to Tg in each of these films provided evidence to conclude a Volmer–Weber growth mode. The presence of MD arrays at the GaN/AlN interface for uncoalesced islands clearly demonstrates that the islands were largely relaxed. The lattice mismatch between GaN and AlN can be accommodated by a triangular array of MDs. If the array line direction is h1 1 2 0i and the MD Burgers vectors are 13h1 1 2 0i; then the Burgers vector will make an angle of 301 with respect to the MD line direction (e.g, b ¼ 13h2 1 1 0i and |b|=a, where a is the lattice constant, and line direction h1 2 1 0i), the edge component of the MD parallel pffiffi to the film/substrate interface is bedge;== ¼ 23a: For a triangular MD grid, the misfit accommodation e is given by [32,33]
4. Discussion
¼
The data presented above provides details of the mechanism for growth of this specific materials
where l is the MD spacing in the grid. For fully relaxed GaN on AlN, ¼ 2:4% and Eq. (1) gives a
Fig. 8. Plan view g ¼ 1 1 2 0 weak-beam image of Material 150 with 200 s of GaN deposition (the nominal GaN thickness was 70 nm). The AlN shows a TD density of 4 1010 cm2. The GaN island showed a clear MD mesh with average line direction h1 1 2 0i and MD spacing of 19 nm.
3 bedge;== ; 2 l
(1)
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MD spacing of 17.3 nm. The spacing of 19 nm, determined from Fig. 8, shows that the 200 s GaN island, had relaxed to at least 90%. The degree of relaxation calculated using data from X-ray diffraction and extrapolated to growth temperature for Material 150 was between496% (50 s of growth) and 99% (coalesced film). An analysis for Material 25 follows in a similar manner; the relaxation data measured by X-ray diffraction predicts491% (50 s of growth) and 99% (coalesced film) relaxation. As is identified earlier by TEM, TDs were observed only in GaN films that have undergone island coalescence; as was determined in the GaN/ sapphire system [5], this asserts that TD generation is primarily a consequence of island coalescence. For GaN growth on the two different AlN layers, the initial GaN island density varies by an order of magnitude. As these islands continued to grow laterally, it follows that the density of uncoalesced islands largely determine their average diameter at coalescence, Dg. Additionally, a previous report [34] has demonstrated that the FWHM of inclined peaks such as the ð2 0 2 1Þ provide a reasonable approximation for y, the degree of twist misorientation in the GaN film. The pure edge TD density may be estimated from the average island size at coalescence, Dg, and the average twist misorientation, y. Rotational misorientation, with respect to the surface normal, of two merging GaN islands is accommodated by the generation of a low-angle grain boundary that consists of pure TDs with line directions parallel to the surface normal. The average dislocation spacing L is simply given as b L¼ ; y
(2)
where b is the magnitude of the Burgers vector of a pure edge TD in GaN and is equal to the a-axis lattice constant. The TD density is simply given as TD y ¼ : area bDg
(3)
The experimentally determined Dg values were 0.9 and 2.5 mm and y values were 260 and 218 arcsec (estimated from the ð2 0 2 1Þ rocking curve width) for Materials 25 and 150, respec-
tively. Using Eq. (3), an estimated TD density of 4.3 108 cm2 was calculated for Material 25, and a density of 1.4 108 cm2 was calculated for Material 150. Both of these values predict a dislocation density approximately one-half of the value reported above measured using plan view TEM. Step termination on a crystal surface must be associated with a screw-component TD. In the case of GaN films, which typically have few pure screw component (b=7[0001]) TDs, the areal density of surface step terminations is an excellent measure of the mixed-character (b ¼ 13h1 1 2 3i TD density [35]. Subtracting the density of mixedcharacter TDs from the overall measured TD density yields a density for pure-edge type TD’s of 6 108 and 1 108 cm2 for Material 25 and 150, respectively—these values are in close agreement with those determined simply from a lowangle grain boundary model. In our earlier work on GaN growth on sapphire, we observed that isolated islands often had a single mixed character TD. Isolated TDs (i.e., those not associated with low-angle interfaces) will not strongly contribute to the broadening of an X-ray rocking curve, rather, isolated dislocations give rise to static diffuse scattering away from the Bragg peak. Thus, we speculate that a substantial fraction of the mixed character TDs form earlier in the growth process, possibly within isolated islands.
5. Summary We explored GaN growth on a representative rough AlN initial layer (Material 25) and a smooth AlN initial layer (Material 150). X-ray analysis showed that the AlN layers were relaxed. Despite the significant difference in the morphology of the AlN initial layer, the initial GaN on the AlN grew as three-dimensional, relaxed islands in both cases (misfit dislocations were observed at the GaN/AlN interface). TDs were observed to form at the coalescence boundaries between adjacent GaN islands—these TDs accommodated misorientation between islands. A simple model based on grain misorientation has closely predicted the density of edge TDs. We conclude both AlN on SiC and
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GaN on relaxed AlN grew by a Volmer–Weber growth mode. [12] [13]
Acknowledgements [14]
The authors wish to thank P. Waltereit, P. Fini, and T. Katona for useful discussions and technical assistance. This work was supported in part by the Office of Naval Research under the CANE MURI program (H. Dietrich, program manager) and made use of the MRL Central Facilities supported by the National Science Foundation under Award No. DMR00-80034. The work of F. Wu was supported by the JST/ERATO program at UCSB. References
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Journal of Crystal Growth 273 (2004) 38–47 www.elsevier.com/locate/jcrysgro
Structural and morphological evolution of GaN grown by metalorganic chemical vapor deposition on SiC substrates using an AlN initial layer B. Morana,, F. Wua,b, A.E. Romanova,1, U.K. Mishrac, S.P. Denbaarsa,b,c, J.S. Specka,b a Materials Department, University of California, Santa Barbara, CA 93106-5050, USA JST/ERATO, UCSB Group, University of California, Santa Barbara, CA 93106-5050, USA c Electrical and Computer Engineering Department, University of California, Santa Barbara, CA 93106-5050, USA b
Received 5 February 2004; accepted 16 August 2004 Communicated by C.R. Abernathy Available online 7 October 2004
Abstract The morphological and structural evolution is presented for GaN grown by metalorganic chemical vapor deposition on 25 nm thick or 150 nm thick AlN initial layers on (0 0 0 1) 4 H and 6 H SiC substrates. The 25 nm thick AlN on SiC was a rough, partially coalesced film, whereas the 150 nm thick AlN on SiC was smooth and was characterized by a step-terrace structure. Both the 25 and 150 nm thick AlN layers on SiC were nearly free of elastic strains. For both AlN initial layers, the GaN films grew by a coarse islanding mechanism. Measurement of the GaN (0 0 0 2) interplanar spacing shows that these islands were largely strain relaxed throughout the growth process. Plan view transmission electron microscopy (TEM) showed a well developed misfit dislocation network at the GaN island/AlN interface. Cross sectional TEM revealed that these islands are free of threading dislocations (TDs) prior to coalescence. A simple, lowangle grain boundary model based on island misorientation at coalescence provides reasonable agreement with TD densities measured using plan view TEM. An electron mobility value as high as 818 cm2/V s has been measured at room temperature confirming a high degree of material quality. r 2004 Elsevier B.V. All rights reserved. PACS: 61.72.y Keywords: A1. Line defects; A3. Metalorganic chemical vapor deposition; B1. Nitrides
Corresponding author. Tel.: +1-805-8938869; fax: +1-805-
893-5263. E-mail address: [email protected] (B. Moran). 1 Permanent address: Ioffe Physico-Technical Institute, St. Petersburg, Russia.
1. Introduction The lack of a native substrate has focused research effort in the wide band gap nitrides into
0022-0248/$ - see front matter r 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.jcrysgro.2004.08.012
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the hetereoepitaxial growth on a variety of different substrates. The in-plane lattice mismatch of 16% for GaN and sapphire and 3.4% for GaN and SiC leads to the formation of a high density of threading dislocations (TDs) in the epitaxial GaN layers (108–1011 cm2), the presence of which have been shown to be detrimental to the performance of both optoelectronic and electronic devices [1,2]. High quality, c-plane epitaxial GaN thin films were produced on sapphire substrates using MOCVD by employing a two-step growth technique [3,4]. This method produced a Volmer–Weber growth mode, in which TDs form along low-angle grain boundaries of coalescing nuclei [5,6]. In recent years, GaN-based optoelectronic devices grown in this manner have become commercially available from many manufacturers. High power, high-frequency GaN-based electronic devices are a comparatively undeveloped field, though significant development activities have resulted in promising reports of high-performance GaN-based transistors [7,8]. Due to self-heating effects, GaN-based transistors require significant thermal dissipation; it has been demonstrated that SiC (thermal conductivity, k3.4 W/cm K at 300 K) [9] is clearly a more suitable substrate than sapphire (k0.42 W/cm K at 300 K) [10] for these applications [11]. There are reports of the growth of GaN on SiC without the use of any intermediate buffer layer. Sakaski and Matsuoka successfully deposited GaN directly on SiC by MOCVD [12], followed by a few groups in the 1990s [13,14]. However, as demonstrated by Lahreche et al. [15], the growth of GaN directly on SiC under typical growth conditions results in a rough, islanded GaN layer after long growth time due to poor surface wetting. The use of an AlN initial layer promotes surface wetting and reduces the lattice mismatch (2.4%) between AlN and GaN. Reports of high-performance AlGaN/GaN HEMT devices grown by MOCVD on SiC employ AlN initial layers [16–18]. For instance, Shealy et al. [18] describe an AlN initial layer as an important step for realizing highperformance HEMTs. An early report by Weeks et al. [19] suggested that the growth temperature of AlN in the range of 500–1050 1C resulted in a polycrystalline AlN layer leading to poor quality
39
in subsequently deposited GaN; a high-temperature (HT) initial layer of AlN grown at temperatures greater than 1100 1C was used to produce monocrystalline AlN and therefore resulted in higher quality GaN films. Davis et al. [20] provided evidence that GaN deposition on AlN results predominantly in the formation of GaN islands which eventually coalesce and form a continuous film. Previous reports have identified a large density (1011 cm2) of dislocations in the GaN/AlN interfacial region, primarily composed of half loops (Burgers vector b=7[0 0 0 1]) and both mixed character (b ¼ 13h1 1 2 3i) and pure edge-type dislocations (b ¼ 13h1 1 2 0i). The TD density tended to decrease to 109 cm2 at the film surface [21,22] through dislocation reaction processes after 1 mm of GaN growth. By using room temperature electron mobility as a metric, GaN of very high quality has been produced using a HT AlN initial layer on SiC as room temperature electron mobility values as high as 900 cm2/V s have been reported [23]. Despite the body of work on GaN on SiC growth, details of the GaN growth mode and extended defect generation remain relatively undeveloped. The work described in this paper uses growth interruption experiments for two representative III-nitride nucleation conditions on SiC in an MOCVD growth environment. Characterization of these films at various stages of evolution allows for an identification of the growth mode, provides insight into the relaxation of lattice mismatch, and presents evidence for the mechanism of TD generation.
2. Experimental procedure Epitaxial AlN and GaN films were grown by MOCVD at 1050 1C in a Thomas Swan closecoupled showerhead reactor on 4 H and 6 H SiC substrates. The substrates were baked in H2 for 10 min at growth temperature prior to film deposition. AlN was grown at 76 Torr, using trimethylaluminum (TMAl) and NH3 at a growth rate of 0.75 A˚/s. GaN was grown at 760 Torr using trimethylgallium (TMGa) and NH3 at an
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approximate growth rate of 3.5 A˚/s. H2 was used as a carrier gas in all cases. Where applicable, silicon doping of GaN was carried out using disilane (Si2H6) as a precursor. AlN layers of 25 and 150 nm in thickness were first deposited on the SiC substrate prior to GaN growth to generate two comparative sample sets. These thicknesses were chosen as they generated two disparate GaN island nucleation scenarios. GaN films grown subsequently on these layers will be known respectively as ‘Material 25’ and ‘Material 150’. Their morphological and structural evolution was observed by performing growth interruption experiments similar to the work of Fini et al. [6]. Growth times of 50, 200, and 600 s were selected to observe the initial stages of GaN growth for both Material 25 and 150 prior to film coalescence. Cross sectional and plan view specimens for transmission electron microscopy (TEM) were prepared by wedge polishing and argon ion milling. Scattering contrast TEM images were recorded with either a JEOL 2000 FX or an FEI Tecnai G2 Sphera; high-resolution TEM (HRTEM) imaging was performed on a JEOL 2010HR. The operating voltage for all TEM imaging was 200 kV. Scanning electron microscopy (SEM) was performed using an FEI Sirion operating at 2 kV. Atomic force microscopy (AFM) was conducted using a Digital Instruments D-3000 SPM in tapping mode. X-ray diffraction scans were recorded using a Philips X-Pert MRD in triple-axis configuration. The (0 0 0 l) interplanar spacing for each evolution sample was determined by measuring the relative separation of the GaN (0 0 0 2) peak from the SiC (0004) or (0006) peak (4H or 6H SiC substrates, respectively) in an o–2y scan. Rocking curves (o-scans) of the GaN ð2 0 2 1Þ peak were performed in a skew geometry (realized by tilting the sample by 751) to determine the twist mosaic in the samples. An X-ray reciprocal space map was recorded of Material 150 near the GaN ð1 0 1 5Þ reflection, with the ensuing room temperature a lattice parameter of AlN being estimated using the ð1 0 1 10Þ reflection of 4H SiC as an angular reference. Electron mobility was measured using indium metal contacts and standard van der Pauw
geometry over 5 mm 5 mm.
a
sample
dimension
of
3. Results Particular consideration was given to the resultant AlN surface morphology prior to GaN deposition. The 25 nm thick AlN layer, shown in Fig. 1a, was a partially coalesced, islanded film with an average grain size of 30 nm. Though the tops of the AlN appeared to be flat, the overall RMS roughness of the layer was 1.6 nm. The AFM image in Fig. 1b shows that 150 nm thick AlN film was fully coalesced and had a surface consisting of atomic steps and terraces with an RMS roughness of 0.3 nm. The TD density in the 150 nm thick AlN, as measured by plan view TEM, was 4 1010 cm2. Details of the morphological evolution, strain relaxation, and microstructure of these AlN films on SiC will be reported elsewhere [24]. One micrometer thick GaN films were grown under identical conditions on both 25 or 150 nm AlN thick initial layers. AFM images shown in Fig. 1c and d of the GaN layers (Materials 25 and 150, respectively) demonstrate that both GaN films had a smooth surface morphology with easily identified atomic steps and terraces. A measured step termination density for Materials 25 and 150 was 4 108 and 1 108 cm2, respectively. The total TD density in Material 25 and 150, quantified using plan view TEM, was 1 109 and 2 108 cm2, respectively. Crosssectional TEM images of these layers with imaging conditions of g ¼ 0002 and g ¼ 1 1 2 0 (not shown) were compared to determine the character of observed TDs. Though an exhaustive statistical analysis was not performed, examination of these images indicated that a majority of the TDs were of pure edge-type (b ¼ 13h1 1 2 0i) and the balance had mixed character (b ¼ 13h1 1 2 3i). o rocking curves determined by X-ray diffraction for the onaxis (0 0 0 2) peak had full width half maximum (FWHM) values of 179 and 153 arcsec and off-axis ð2 0 2 1Þ peak widths of 260 and 218 arcsec for Material 25 and 150, respectively.
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Fig. 1. AFM height images (2 2 mm) of (a) 25 nm thick and (b) 150 nm thick AlN and 5 5 mm AFM height images of (c) 1 mm thick Material 25 and (d) 1 mm thick Material 150.
AFM micrographs of the growth interruption series of both Material 25 and 150 are shown in Fig. 2 to provide morphological and geometric detail of the island nucleation and growth process. Each island had primarily a flat (0 0 0 1) surface, which was sustained through film coalescence. SEM images (not shown) showed that the islands were truncated hexagonal pyramids with f1 0 1 1g inclined facets. Table 1 provides the measured GaN island aspect ratio and density. The density is reported for 50 s of growth in for both Materials 25 and 150 as an approximation of the GaN island density because there was no morphological evidence for coalescence in either material at this early stage of growth. Significant coalescence was first observed in Material 25 after 200 s of growth, and the film was nearly continuous after 600 s of growth. Comparatively, clustering was first widely observed at 600 s in Material 150, with complete film coalescence occurring beyond 1200 s of
growth. Analysis several of AFM scans showed that the average island diameter at coalescence was 0.9 mm for Material 25 and 2.5 mm for Material 150. Diffraction contrast TEM studies were performed on representative cross-section samples to reveal the microstructure evolution. An isolated island of Material 25 after 50 s of GaN growth, shown in Fig. 3, was found to be free of TDs. The GaN does show significant strain contrast in both the g ¼ 0 0 0 2 and g ¼ 1 1 2 0 image—this contrast is attributed to both the strain fields of misfit dislocations (MDs) at the GaN/AlN interface and the rough nature of this interface. The contrast from the SiC at the interface appears to correspond to the island structure of the AlN. Fig. 4a shows two coalesced GaN islands of Material 150. The g ¼ 0 0 0 2 image has little contrast—consistent with a smooth GaN/AlN interface. The g ¼ 1 1 2 0 image shows the high density of pure edge
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Fig. 2. AFM images (15 15 mm) of GaN growth interruption study. Material 25: (a) and (b) are height images, (c) is a tip amplitude image. Material 150: (a) is a height image, (e) and (f) are tip amplitude images.
Table 1 Measured average values for GaN island aspect ratio and density Growth time (s)
Material 25 Aspect ratio (width/height)
50 200 600
16.3 9.8 Coalesced film
TDs (b ¼ 13h1 1 2 0i) in the AlN and also shows a segment of a single TD—we believe that this TD formed was a result of island–island coalescence. In the case of Material 25 in a late stage of coalescence (600 s, Fig. 4b), TDs of both pure edge and mixed character were unambiguously observed. Plan view TEM micrographs of 1 mm thick fully coalesced films are shown in Fig. 5. The TD density of Material 25 and 150 was 1 109 and 2 108 cm2.
Material 150 Density (#/cm2) 7
8.1 10 3.6 107 Coalesced film
Aspect ratio
Density
6.1 6.6 8.5
8.4 106 8 106 8 106
An X-ray reciprocal space map was recorded near the GaN ð1 0 1 5Þ reflection of a 1 mm thick Material 150 layer and is shown in Fig. 6. The estimated a lattice constant calculated from this data of 3.1143 A˚ for the 150 nm thick AlN was within 0.01% of its relaxed bulk value [19]. Given the similar a-axis thermal expansion coefficient of AlN and SiC (4.2 106 cm2) [21], it is assumed that the 150 nm thick AlN was nearly fully relaxed at growth temperature.
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Fig. 3. Bright field cross-section TEM image of single Material 25 GaN island (after 50 s of growth) in g ¼ 0002 (top) and g ¼ 1 1 2 0 (bottom) two beam imaging conditions.
Fig. 4. Bright field cross-section TEM images recorded with either g ¼ 0002 or g ¼ 1 1 2 0 imaging conditions of (a) Material 150 after 200 s of growth and (b) Material 25 after 600 s of growth. The arrow in the GaN island in (a) indicates a TD segment.
o–2y X-ray diffraction scans were recorded to determine the GaN (0 0 0 2) absolute peak separation from the relevant SiC substrate peak ((0 0 0 4) for 4H SiC and (0 0 0 6) for 6H SiC). Based upon this data, the (000 l) interplanar spacing and thus
43
the measured c lattice constant can be determined for the GaN films at the various stages of film evolution. The value for the a lattice constant and thus the in-plane strain in each sample was determined assuming purely biaxial stress using elastic constants as calculated by Kim et al. [25] and bulk unstrained lattice constant ( values published in the literature (a ¼ 3:1881 A, ( c ¼ 5:1844 A—these lattice constants were used due to the undoped and low background impurity character of the GaN films produced in this study) [26]. It is important to understand the strain in the GaN film at the growth temperature (Tg). Without the availability of in situ monitoring, the strain in the GaN film was estimated by using the in-plane strain at room temperature and then accounting for the thermal strain developed during postgrowth cool down by the differences in the temperature-dependent coefficients of thermal expansion (CTE), Da, between GaN and SiC [27,28]. Fig. 7 shows the in-plane strain data for Materials 25 and 150 at RT and extrapolated to Tg. There are several points to be made regarding the data in this graph. Both of the thick films (i.e., long growth times) of Material 25 and 150 were under tensile strain at RT; after extrapolation to Tg, these films were nearly strain-free. Regarding the Tg data only, there was also no particular point at which coalescence events appeared to markedly change the strain state of the films; rather, a more gradual trend of relaxation of compressive stress was observed as the growth time increased. Finally, the strain in the GaN extrapolated to Tg even at a short growth time of 50 s was low, with values of 2 103 (0.2%) and 8 104 (0.08%) for Materials 25 and 150, respectively. These low residual strain values are indicative of a high degree of relaxation of the 2.4% compressive lattice mismatch strain between GaN and AlN. TEM studies of both Materials 25 and 150 show MDs at the GaN/AlN interface with a periodicity consistent with largely relaxed GaN throughout the growth. Fig. 8 shows a plan view weak beam image of an uncoalesced Material 150 film. The 150 nm thick AlN layer has a TD density of 4 1010 cm2, as determined from regions without GaN coverage. A MD network was clearly
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Fig. 5. Plan view, bright field g ¼ 1 1 2 0 TEM images of (a) 1 mm thick Material 25 (b) 1 mm thick Material 150.
Fig. 7. Measured room temperature in-plane strain and the inplane strain extrapolated to the growth temperature for each sample in the growth evolution study.
Fig. 6. Reciprocal space map of Material 150 (GaN thickness is 1 mm). The map was taken near the asymmetric (1 0 1 5) GaN reflection in coplanar geometry with shallow incidence and rocking curve configuration. ‘‘r.l.u’’ refers to dimensionless reciprocal lattice units (lq 10000).
seen at the GaN/AlN interface. From detailed analysis of the images, the average MD line direction was determined to be h1 1 2 0i: The average MD spacing in the h1 1 2 0i directions was 19 nm for the image shown in Fig. 8. As we will discuss below, the 2.4% mismatch between GaN and AlN can be fully accommodated by a triangular MD array with h1 1 2 0i line direction and Burgers vector 13h1 1 2 0i with MD spacing of 17.3 nm. Thus, the MD spacing of 19 nm is
consistent with 90% relaxation in the GaN islands. Based on our high-resolution imaging (not shown) and diffraction contrast studies, we tentatively conclude that the MD dislocations have Burgers vector 13h1 1 2 0i (further details on the nature of the dislocations at the GaN/AlN interface will be presented elsewhere [24]). MDs were also observed by inclined view weak beam images on cross-section samples at the GaN/AlN interface both for coalesced films and isolated islands. A silicon-doped GaN layer grown on top of Material 25 and 150 template layers provided an appreciable number of carriers to examine the transport properties of these films. Bulk GaN
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Fig. 9. The dependence of electron mobility on measured carrier concentration for several samples of Material 25 and 150.
electron mobility values for several 1.5–2.0 mm thick samples are plotted v s. carrier concentration in Fig. 9. The maximum measured values were 804 and 818 cm2/V s for carrier concentrations of 5 1016 cm3 for Materials 25 and 150, respectively. In both Material 25 and 150, we observed this expected fall-off in mobility with an increase in Si doping. Several groups have shown that scattering due to charged dislocation lines does not limit the transverse room temperature electron mobility for GaN with TD densities in the range 108–109 cm2 and free electron concentrations in the range of 1016–1017 cm3 [29–31]. Thus, it is not surprising that Materials 25 and 150 have a similar room temperature electron mobility despite having nearly an order of magnitude different TDD.
system. It is clear from the AFM data that the GaN on AlN grew via a coarse islanding mechanism. The X-ray diffraction data indicates that the AlN layers were largely relaxed on the SiC. In turn, the GaN was nearly fully relaxed throughout its growth on AlN. Additionally, the low strain when extrapolated to Tg in each of these films provided evidence to conclude a Volmer–Weber growth mode. The presence of MD arrays at the GaN/AlN interface for uncoalesced islands clearly demonstrates that the islands were largely relaxed. The lattice mismatch between GaN and AlN can be accommodated by a triangular array of MDs. If the array line direction is h1 1 2 0i and the MD Burgers vectors are 13h1 1 2 0i; then the Burgers vector will make an angle of 301 with respect to the MD line direction (e.g, b ¼ 13h2 1 1 0i and |b|=a, where a is the lattice constant, and line direction h1 2 1 0i), the edge component of the MD parallel pffiffi to the film/substrate interface is bedge;== ¼ 23a: For a triangular MD grid, the misfit accommodation e is given by [32,33]
4. Discussion
¼
The data presented above provides details of the mechanism for growth of this specific materials
where l is the MD spacing in the grid. For fully relaxed GaN on AlN, ¼ 2:4% and Eq. (1) gives a
Fig. 8. Plan view g ¼ 1 1 2 0 weak-beam image of Material 150 with 200 s of GaN deposition (the nominal GaN thickness was 70 nm). The AlN shows a TD density of 4 1010 cm2. The GaN island showed a clear MD mesh with average line direction h1 1 2 0i and MD spacing of 19 nm.
3 bedge;== ; 2 l
(1)
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MD spacing of 17.3 nm. The spacing of 19 nm, determined from Fig. 8, shows that the 200 s GaN island, had relaxed to at least 90%. The degree of relaxation calculated using data from X-ray diffraction and extrapolated to growth temperature for Material 150 was between496% (50 s of growth) and 99% (coalesced film). An analysis for Material 25 follows in a similar manner; the relaxation data measured by X-ray diffraction predicts491% (50 s of growth) and 99% (coalesced film) relaxation. As is identified earlier by TEM, TDs were observed only in GaN films that have undergone island coalescence; as was determined in the GaN/ sapphire system [5], this asserts that TD generation is primarily a consequence of island coalescence. For GaN growth on the two different AlN layers, the initial GaN island density varies by an order of magnitude. As these islands continued to grow laterally, it follows that the density of uncoalesced islands largely determine their average diameter at coalescence, Dg. Additionally, a previous report [34] has demonstrated that the FWHM of inclined peaks such as the ð2 0 2 1Þ provide a reasonable approximation for y, the degree of twist misorientation in the GaN film. The pure edge TD density may be estimated from the average island size at coalescence, Dg, and the average twist misorientation, y. Rotational misorientation, with respect to the surface normal, of two merging GaN islands is accommodated by the generation of a low-angle grain boundary that consists of pure TDs with line directions parallel to the surface normal. The average dislocation spacing L is simply given as b L¼ ; y
(2)
where b is the magnitude of the Burgers vector of a pure edge TD in GaN and is equal to the a-axis lattice constant. The TD density is simply given as TD y ¼ : area bDg
(3)
The experimentally determined Dg values were 0.9 and 2.5 mm and y values were 260 and 218 arcsec (estimated from the ð2 0 2 1Þ rocking curve width) for Materials 25 and 150, respec-
tively. Using Eq. (3), an estimated TD density of 4.3 108 cm2 was calculated for Material 25, and a density of 1.4 108 cm2 was calculated for Material 150. Both of these values predict a dislocation density approximately one-half of the value reported above measured using plan view TEM. Step termination on a crystal surface must be associated with a screw-component TD. In the case of GaN films, which typically have few pure screw component (b=7[0001]) TDs, the areal density of surface step terminations is an excellent measure of the mixed-character (b ¼ 13h1 1 2 3i TD density [35]. Subtracting the density of mixedcharacter TDs from the overall measured TD density yields a density for pure-edge type TD’s of 6 108 and 1 108 cm2 for Material 25 and 150, respectively—these values are in close agreement with those determined simply from a lowangle grain boundary model. In our earlier work on GaN growth on sapphire, we observed that isolated islands often had a single mixed character TD. Isolated TDs (i.e., those not associated with low-angle interfaces) will not strongly contribute to the broadening of an X-ray rocking curve, rather, isolated dislocations give rise to static diffuse scattering away from the Bragg peak. Thus, we speculate that a substantial fraction of the mixed character TDs form earlier in the growth process, possibly within isolated islands.
5. Summary We explored GaN growth on a representative rough AlN initial layer (Material 25) and a smooth AlN initial layer (Material 150). X-ray analysis showed that the AlN layers were relaxed. Despite the significant difference in the morphology of the AlN initial layer, the initial GaN on the AlN grew as three-dimensional, relaxed islands in both cases (misfit dislocations were observed at the GaN/AlN interface). TDs were observed to form at the coalescence boundaries between adjacent GaN islands—these TDs accommodated misorientation between islands. A simple model based on grain misorientation has closely predicted the density of edge TDs. We conclude both AlN on SiC and
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GaN on relaxed AlN grew by a Volmer–Weber growth mode. [12] [13]
Acknowledgements [14]
The authors wish to thank P. Waltereit, P. Fini, and T. Katona for useful discussions and technical assistance. This work was supported in part by the Office of Naval Research under the CANE MURI program (H. Dietrich, program manager) and made use of the MRL Central Facilities supported by the National Science Foundation under Award No. DMR00-80034. The work of F. Wu was supported by the JST/ERATO program at UCSB. References
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