Nucleation layer evolution in metal-organic chemical vapor deposition
grown GaN
X. H. Wu, D. Kapolnek, E. J. Tarsa, and B. Heying
Materials Department, College of Engineering, University of California, Santa Barbara, California 93106
S. Keller, B. P. Keller, and U. K. Mishra
Electrical and Computer Engineering Department, College of Engineering, University of California,
Santa Barbara, California 93106
S. P. DenBaars
Materials Department and Electrical and Computer Engineering Department, College of Engineering,
University of California, Santa Barbara, California 93106
J. S. Specka)
Materials Department, College of Engineering, University of California, Santa Barbara, California 93106
Received 25 September 1995; accepted for publication 8 January 1996
The structure and morphology of lowgrowth temperature GaNnucleation layers have been studied
using atomic force microscopy AFM , reflection high energy electron diffraction RHEED , and
transmission electron microscopy TEM . The nucleation layers were grown at 600°C by
atmospheric pressure metalorganic chemical vapor deposition MOCVD on c-plane sapphire. The
layers consist of predominantly cubic GaN c-GaN with a high density of stacking faults and twins
parallel to the film/substrate interface. The average grain size increases with increasing layer
thickness and during the transition from low temperature 600°C to the high temperatures
1080°C necessary for the growth of device quality GaN. Upon heating to 1080°Cthe nucleation
layer partiallyconverts tohexagonal GaN h-GaN while retaininga highstackingfault density. The
mixed cubic-hexagonal character of the nucleation layer region is sustained after subsequent
high-temperature GaNgrowth. © 1996 American Institute of Physics. S0003-6951 96 01710-8
The high interfacial energy associated with GaN thin creased prior to bulk filmgrowth.7 9 However, detailed char-
films on sapphire substrates leads to three-dimensional island acterization of the nucleation layer microstructure has yet to
be reported. Since the nucleation layer serves as the crystal-
growth of the GaN. As a result, a two-step metalorganic
lographic and morphological template for subsequent GaN
chemical vapor deposition MOCVD growth process has
deposition, the structure of this layer must be understood for
been developed to promote two-dimensional GaNgrowth on
the ultimate optimization of GaN device structures on sap-
sapphire.1,2 This process involves the growth of GaNor AlN
phire. In this letter, we present the characterization of the
nucleation layers at low temperatures, followed by the
morphology and crystallinity of as-grown GaN nucleation
growth of the device quality GaN filmat high temperature.
layers and the microstructural evolution of these layers as
High nucleation densities are more readily achieved at low
they are heated to typical GaNgrowth temperatures.
temperatures, where the supersaturation of the growth spe-
The GaNnucleation layers were grown on c-plane sap-
cies above the substrate is high and adatoms have low sur-
phire using a horizontal flowatmospheric pressure metalor-
face mobilities. In contrast, stabilization of the hexagonal
ganic chemical vapor deposition MOCVD reactor Thomas
phase of GaN h-GaN is achieved at high temperatures
Swan Co., Ltd. . The substrates were first cleaned with sol-
where surface mobilities are sufficient to facilitate step-flow
vents and subjected to an in situ pretreatment in flowing H2
growth.3
at 1050°C. The nucleation layers were grown at 600°Cus-
The 14%lattice mismatch between h-GaN a 3.180
ing trimethylgallium TMGa and ammonia (NH3). Some
Å, c 5.166 Å and basal plane sapphire -Al2O3 , a
nucleation layers were quenched to roomtemperature imme-
4.755 Å, a/ 3 2.748 Å, c 12.991 Å leads to unavoid-
diately following growth, while others were heated to
ably high misfit dislocation densities. As a consequence, high
1080°C and held for several seconds prior to quenching.
threading dislocations densities ( 1010 cm 2) are often ob-
This was done to investigate the evolution of the nucleation
served throughout GaN films grown on sapphire.4 6 Re-
layer during different stages of the standard GaNfilmgrowth
cently, we have been able to successfully grow films with
process. The nucleation layers were characterized using
threading dislocation densities of 7 108 cm 2 by opti-
atomic force microscopy AFM in tapping mode, reflection
mizing the nucleation layer growth conditions.3 Previous
high energy electron diffraction RHEED , and cross-
studies have demonstrated that nucleation layer growth con-
sectional transmission electron microscopy TEM . RHEED
ditions strongly affect the electrical and optical properties of
studies were performed using 7 keVelectrons in a separate
GaNdevice layers.2,7,8 In general, the structure of the nucle-
molecular beamepitaxy chamber.
ation layer is believed to evolve as the temperature is in-
The initial stages of GaN nucleation on sapphire were
first examined using AFM. Figure 1 shows AFM micro-
a
Electronic mail: [e-mail] graphs of as-grown nucleation layers of 2 nm Fig. 1 a and
Appl. Phys. Lett. 68 (10), 4 March 1996Ź 0003-6951/96/68(10)/1371/3/$10.00Ź ©1996American Institute of Physics 1371
Copyright ©2001. All Rights Reserved.
FIG. 1. Tapping mode AFMimages of GaNnucleation layers: a as-grown
layer with nominal 2 nmthickness, b as-grown layer with nominal 20 nm
thickness, c nominal 20 nmthick layer after heating to 1080°C.
20 nm Fig. 1 b nominal thickness. The average grain size
increases from25 to 33 nmand the rms roughness increases
from 3.11 to 6.18 nm as the nominal thickness increased
from2 to 20 nm. The morphology of the films after heating
to 1080°C Fig. 1 c will be discussed below.
The as-grown 20 nmthick nucleation layer surfaces pro-
duced RHEEDpatterns corresponding primarily to 111 ori-
ented cubic GaN c-GaN , with 180° rotational twinning
Å»
about the surface normal, as evident in the 110 azimuth
c
RHEED pattern in Fig. 2 a calculated diffraction patterns
with one c-GaN variant are also shown in Fig. 2 . The
Å» Å» c
RHEEDpattern was spotty in the 21 1 azimuth the sub-
script c refers to cubic indexing and the subscript h refers to
hexagonal , as shown in Fig. 2 b , thus indicating that the
surface is three dimensional. In contrast, the RHEEDpattern
FIG. 3. Cross-section TEMdiffraction and imaging on as-grown nucleation
Å» layers a , b , and c and on nucleation layers after high-temperature
observed along the 110 azimuth was somewhat streaky,
c
exposure d , e , and f . The images and diffraction patterns were re-
suggesting a smooth surface. However, cross-section TEM
Å»
corded either near or down a 110 zone axis. a Selected area diffraction
c
Å»
diffraction showed that the streaking observed along 110 c
pattern corresponding to predominantly c-GaN. b Bright field image. A
high density faulting density is clearly observed in the GaNgrains. Faceting
is not associated with surface morphology but rather with a
is also clearly observed. c High resolution image recorded down a
high density of stacking faults and twins in the close-packed
Å»
110 zone axis showing stacking disorder in the grains. d Selected area
c
111 c-GaN planes parallel to the film/substrate interface.
diffraction pattern. The diffraction pattern corresponds to mixed c-GaN/h-
Å»
GaN. e Bright field image. Ahigh faulting density is clearly observed in
Chevroning was evident in 110 RHEEDpatterns but not
c
the GaNgrains. The grains appear more rounded than the as-grown nucle-
Å» Å» c
in 21 1 RHEED patterns. Chevroning is associated with
ation layer. f High resolution image showing stacking disorder in the
grains. The outer shell of the grains is predominantly h-GaN.
facet formation and demonstrates that the facet planes lie in a
Å»
110 zone.
c
Cross-sectionTEManalysis confirmed the highly faulted
but predominantly cubic character of the 20 nmthick nucle-
ation layers and further showed the highly faceted grain mor-
phology Fig. 3 . The selected area diffraction pattern shown
in Fig. 3 b corresponds to a mixture of c-GaN and h-GaN
with the orientation relationship 111 001 ,Ź and
c h
Å»
110 100 . The regions of pure c-GaNor h-GaNwere
c h
composed of only a few close-packed layers, as shown in
both the bright field image of Fig. 3 a and the high resolu-
tion image in Fig. 3 c . The qualitative difference in intensi-
ties between the unique cubic and hexagonal reflections in
Å»
the 110 zone axis diffraction patterns indicates that this
c
layer is predominantly c-GaN. This may also be confirmed
FIG. 2. RHEEDpatterns fromas-grown nominally 20 nmthick nucleation
Å»
layers experimental patterns top, schematic patterns bottom . a 110 by examination of the relative fraction of c-GaNin the high
c
100 azimuth zone axis . In this orientation, stacking disorder between
h
resolution image Fig. 3 c . The cubic and hexagonal re-
Å» Å» c h
cubic and hexagonal regions leads to streaking. b 21 1 210 azi-
gions can be easily distinguished in cross-section high reso-
muth zone axis . In this orientation, cubic and hexagonal stacking se-
Å»
lution images recorded down a 110 zone axis since the
quences cannot be distinguished and thus no stacking disorder streaks are c
observed. cubic regions show 111 cross fringes that are inclined
c
1372Ź Appl. Phys. Lett., Vol. 68, No. 10, 4 March 1996ŹWu et al.
Copyright ©2001. All Rights Reserved.
70.5° to the film/substrate interface, whereas the hexagonal grains are composed of nearly equal proportions of c-GaN
regions showvertical (010)h fringes. and h-GaN. Both scattering contrast and high resolution im-
Although the primary orientation between c-GaN and ages Fig. 3 f , confirmthat the grains have lost their dis-
Å»
tinct facets and now appear somewhat rounded. However,
h-GaN is 111 001 , and 110 100 , the cubic
c h c h
the mosaic in the grains determined in plan-viewTEMmain-
phase also displays 180° rotational twinning about 111 c
Ż10 c h
tains a range of 1° to 3°. The conversion fromc-GaN to
i.e., 1 100 ). Since, the reciprocal lattice points for
Å» Å» c h
h-GaNis predominant near the free surfaces of the grains as
c-GaNand h-GaNare coincident for a 21 1 210 zone
shown in Fig. 3 f . Finally, the inclined twins in the cubic
axis, as shown in the schematic diffraction pattern in Fig.
regions were not observed after the high-temperature step.
2 b , it is not possible to distinguish between cubic or hex-
These results demonstrate that the GaNnucleation layer
agonal stacking or resolve stacking disorder in this orienta-
Å» Å» c is fully crystalline and epitaxial upon growth. Polytypismin
tion. The spotty RHEED patterns recorded along 21 1
GaN thin films has been previously reported. However, the
210 therefore, correctly convey the rough nature of the
h
predominantly cubic nature of the as-grown nucleation layer
GaN surface. In contrast, the reciprocal lattice points for
Å» has not been reported. While GaN growth at high tempera-
c-GaN and h-GaN do not coincide in the 110 100
c h
tures results in a homogeneous h-GaNfilm, these results sug-
zone axis, as shown in Fig. 2 a . Correspondingly, the
gest that the low-temperature epitaxy is more complicated.
Å»
110 100 zone axis is sensitive to stacking sequence
c h
Currently, we can only speculate on the physical basis for the
and stacking disorder. The streaks observed in RHEEDpat-
formationof c-GaNat lowtemperatures. It wouldappear that
Å»
terns 110 100 therefore originate fromstacking dis-
c h
the formation of the cubic phase may either be associated
order and not surface morphology.
with the relatively high pressures used for these growths or a
The morphology of the grains seen in cross-sectionTEM
reduction in free-surface energy of the initial GaNislands.
is consistent with a 111 contact plane at the film/substrate
The nucleation layer substantially coarsens during the
interface and a well-defined facet angle of 55° with respect
temperature increase to 1080°C, even though the duration of
to the substrate. This angle is very close to the angle between
the heating cycle was only about 2 min. The (001)h surfaces
(111)c and (100)c planes. However, the grains are not proper
become stabilized under these conditions, resulting in the
single crystals because of their mixed cubic-hexagonal na-
development of flat surfaces in the grains. However, the con-
ture and high density of twins. Thus, a proper crystallo-
version to h-GaNis not accompanied by an appreciable de-
graphic assignment to the inclined facets is neither possible
crease in stacking fault density or in mosaic in the grains.
nor appropriate. The faceted morphology of the as-grown
Despite this, subsequent high-temperature GaN growth on
nucleation layer was also observed in plan-viewTEM. In this
these nucleation layers leads to specular films that show a
case, both triangular and hexagonal islands were seen in pro-
well-defined step-terrace structure. The typical threading dis-
jection. The island edges were all oriented along one of three
2
location density for a 1 m thick film is 7 108 cm .3
110 direction separated by 120° . For the triangular is-
c
Ongoing studies are directed at understanding the high-
lands, two orientations were observed corresponding to the
temperature growth evolution fromrough three-dimensional
180° twin orientations of c-GaN. The as-grown nucleation
surfaces to smooth two-dimensional surfaces and the mecha-
layers displayed a 1° to 3° in-plane mosaic spread as mea-
nisms of threading dislocation reduction.
sured fromselected area electron diffraction patterns.
The authors gratefully acknowledge the support of the
In addition to planar faulting parallel to the film/
Army Research Office through a contract supervised by Dr.
substrate interface, cubic twins were also observed on in-
John Zavada. Partial funding for this research was provided
Å»
clined (111)c planes, as shown by high resolution TEM Fig.
by the NSF Materials Research Laboratories DMR 91-
Å»
3 c . This result is not surprising as (111)c and (111)c are
23048 .
both proper close-packed planes in the cubic phase.
The morphological evolution of a 20 nmthick nucleation
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6
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Appl. Phys. Lett., Vol. 68, No. 10, 4 March 1996ŹWu et al. 1373
Copyright ©2001. All Rights Reserved.