JOURNAL OF APPLIED PHYSICS 107, 043527 共2010兲
High-temperature molecular beam epitaxial growth of AlGaN/GaN on GaN
templates with reduced interface impurity levels
G. Koblmüller,1,3,a兲 R. M. Chu,2 A. Raman,2 U. K. Mishra,2 and J. S. Speck1
1
Department of Materials, University of California, Santa Barbara, California 93106, USA
Department of Electrical and Computer Engineering, University of California, Santa Barbara, California
93106, USA
3
Walter Schottky Institut, Technische Universität München, D-85748 Garching, Germany
2
共Received 27 October 2009; accepted 7 December 2009; published online 25 February 2010兲
We present combined in situ thermal cleaning and intentional doping strategies near the substrate
regrowth interface to produce high-quality AlGaN/GaN high electron mobility transistors on
semi-insulating 共0001兲 GaN templates with low interfacial impurity concentrations and low buffer
leakage. By exposing the GaN templates to an optimized thermal dissociation step in the
plasma-assisted molecular beam epitaxy environment, oxygen, carbon, and, to lesser extent, Si
impurities were effectively removed from the regrowth interface under preservation of good
interface quality. Residual Si was further compensated by C-doped GaN via CBr4 to yield highly
resistive GaN buffer layers. Improved N-rich growth conditions at high growth temperatures were
then utilized for subsequent growth of the AlGaN/GaN device structure, yielding smooth surface
morphologies and low residual oxygen concentration with large insensitivity to the 共Al+ Ga兲N flux
ratio. Room temperature electron mobilities of the two-dimensional electron gas at the AlGaN/GaN
interface exceeded ⬎1750 cm2 / V s and the dc drain current reached ⬃1.1 A / mm at a +1 V bias,
demonstrating the effectiveness of the applied methods. © 2010 American Institute of Physics.
关doi:10.1063/1.3285309兴
I. INTRODUCTION
AlGaN/GaN high electron mobility transistors 共HEMTs兲
have recently gained large attention for high-power and highfrequency electronic devices due to their significant merits
associated with the large breakdown field, the high saturation
velocity, and the high sheet charge density of the
polarization-induced two-dimensional electron gas 共2DEG兲.
To fabricate nitride-based HEMT materials, molecular beam
epitaxy 共MBE兲 has become an established growth technique
for state-of-the-art devices1–4 as it offers many benefits, such
as sharp interface and doping control, low background impurity incorporation, and powerful in situ diagnostics for precise growth control.
Despite these promises, direct growth by conventional
plasma-assisted 共PA兲-MBE, typically performed under
metal-rich conditions at fairly low temperatures, was often
hampered by the huge requirements to provide low threading
dislocation densities 共TDDs兲 in the GaN buffer for low device leakage along with smooth, homogeneous AlGaN/GaN
heterointerfaces. Only through complex multistep buffer layers for effective threading dislocation 共TD兲 reduction during
heteroepitaxial nucleation on foreign substrates 关i.e.,
sapphire5 and SiC 共Ref. 6兲兴 and optimized growth conditions
around the narrow boundary for Ga droplet formation were
successes achieved with atomically smooth metal and surface pit-free heterointerfaces.7,8
Most recently, growth in nitrogen-rich environment by
PAMBE 共Ref. 9兲 and ammonia-based MBE 共Refs. 10 and 11兲
delivered significant advances. In particular, these were
a兲
Electronic mail: [email protected].
0021-8979/2010/107共4兲/043527/9/$30.00
wider growth windows and potentially higher growth temperatures
共T ⬎ 760 ° C兲
with
suppressed
thermal
decomposition9–13 and lower impurity incorporation,14 resulting in TDDs reduced to ⬍5 ⫻ 109 cm−2,15,16 the highest reported room temperature 共RT兲 bulk electron mobilities 共␮
⬎ 1150 cm2 / V s兲 9 in MBE-grown GaN and exceptional
AlGaN/GaN HEMT power performance on SiC and Si.11,17
In addition, under N-rich growth conditions reduced leakage
currents were reported due to the lack of excess Ga decoration along TDs.18
Ultimately, to increase the intrinsic properties of the
2DEG in AlGaN/GaN structures and to reduce leakage paths,
the number of charged TDs propagating through the device
needs to be further minimized via higher-quality substrates.
Both the use of semi-insulating 共SI兲 or freestanding GaN
templates grown by metal organic vapor phase epitaxy or
hydride vapor phase epitaxy on sapphire resulted recently in
RT 2DEG mobilities ⬎1900 cm2 / V s and record ␮
⬃ 167 000 cm2 / V s at 0.3 K in MBE-regrown AlGaN/GaN
heterostructures due to the overall lower TDDs 共⬃107 – 8
⫻ 108 cm−2兲.19–21
Although with further development of SI and selfsupporting low-TDD GaN templates, parasitic conduction at
the homoepitaxial regrowth interface 共RI兲 presents an independently serious problem. This arises from the typical interface contamination 共mainly O, Si, and C impurities兲 adsorbed during air exposure in standard substrate loading
procedures or at the beginning of regrowth, requiring routes
to eliminate the n-type conductive channel close to the MBEGaN-template interface. Wet chemical etching in HF, HCl,
KOH, and sulfuric acids, together with thermal desorption
via in situ vacuum annealing at elevated temperatures,
107, 043527-1
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043527-2
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Koblmüller et al.
yielded only limited reduction of either O or C
impurities.22–25 Instead, annealing studies in NH3 ambient at
700– 800 ° C resulted sometimes in sufficiently clean and
stoichiometric 共0001兲 GaN surfaces, although not performed
in a closed MBE growth environment.25 Most of these in situ
annealing methods were additionally constricted by the very
low thermal dissociation temperature of the 共0001兲 GaN surface, i.e., ⬃750 ° C in vacuum.12,13,24
More effective compensation of the interfacial impurities
yielding highly insulating GaN buffer layers in the regrown
MBE-AlGaN/GaN structures was demonstrated by the use of
different acceptors, such as berrylium19 and carbon,26,27 or
through graded iron concentration profiles in the underlying
SI GaN templates.17 Nevertheless, compensation-doped GaN
buffer layers need to be well balanced in thickness and doping concentration for sufficient isolation and to minimize dispersion effects in the 2DEG at the AlGaN/GaN heterointerface.
In the present study, we merged many of the critical
aspects mentioned above to fabricate high-quality AlGaN/
GaN heterostructures on low-dislocation-density 共0001兲 GaN
templates with improved RI properties and low buffer leakage. Essentially, we developed a sophisticated thermal dissociation procedure prior to GaN buffer nucleation, eliminating
both oxygen and carbon impurities at the RI directly in the
MBE growth environment and without the use of wet chemical etchants. The common propensity for residual Si remaining at the RI was solved by additional compensation doping
via highly insulating C-doped GaN buffer layers. Growth of
the AlGaN/GaN heterostructures on top demonstrated further
the significant merits of the N-rich growth conditions employed in the thermal decomposition regime, with RT 2DEG
mobilities ⬎1750 cm2 / V s and growth surfaces free from
detrimental excess Ga.
II. EXPERIMENTS
As substrates we used semi-insulating 共Fe-doped兲
⬃3.5-␮m-thick 共0001兲 GaN templates grown on c-plane sapphire by MOCVD 共Lumilog兲. The TDD and resistance were
specified as ⬃5 ⫻ 108 cm−2 and ⬃10 M⍀ at RT, respectively. For sample preparation, the 2 in. GaN substrates’
backsides were metalized with ⬃500-nm-thick Ti to enable
radiative heating, then cleaved into 0.5–2 in. pieces, and subsequently degreased in acetone, methanol, and isopropanol.
For reference, two samples were additionally etched in buffered hydrofluoric and hydrochloric acids 共HF and HCl兲 to
investigate the effect of potentially reduced substrate surface
contamination. After blow drying under N2 ambient, the
samples were immediately loaded into the loadlock of the
MBE chamber and were outgassed for 1 h at 400 ° C prior to
growth.
All further experiments were carried out in situ in a
Gen-II MBE system equipped with standard effusion cells
for Ga and Al and a Vecco Unibulb radio frequency plasma
source for supplying active nitrogen. The N2 gas was specified with 7N purity 共99.99999%兲 and cleaned via two inert
gas purifiers at the gas flow inlet prior to entering the plasma
source. The substrate temperature was measured in situ via
FIG. 1. 共Color online兲 共a兲 Schematic of typical MBE–grown AlGaN/GaN
heterostructure on MOCVD-GaN template 共as referred to also in Figs. 7 and
8兲, utilizing C-doped GaN buffer layer, N-rich growth conditions in the GaN
thermal decomposition regime, and 共b兲 an optimized thermal cleaning step
at the RI prior to layer growth. This step consists of alternating cycles of 1
min 共1⬘兲 GaN growth and 2 min 共2⬘兲 thermal decomposition under vacuum
at T = 780 ° C. Labels “g.r.” and “d.r.” refer to the growth and decomposition
rates, respectively. The net amount of evaporated GaN during one such
thermal cleaning cycle corresponds to 2.5 nm.
an optical pyrometer. For intentional carbon doping, we used
a commercially available carbon tetrabromide 共CBr4兲 sublimation system, where the variable foreline pressure, regulated through automated valve control, was employed to control the CBr4 flow. Except for the CBr4 flow 共in foreline
pressure units of mTorr兲, we gave all other molecular fluxes
in 共0001兲 GaN growth rate units 共nm/min兲, as calibrated
from thickness measurements of Ga- and N-limited GaN
films at temperatures of negligible thermal decomposition.28
For conversion, we note that in wurztite GaN c / 2
= 0.259 nm or 1.14⫻ 1015 GaN/ cm2 areal density refer to 1
ML. Moreover, two specific N fluxes were used, i.e.,
⬃3.5 nm/ min 关250 W/0.3 SCCM 共SCCM denotes standard
cubic centimeter per minute兲兴 during the initial thermal dissociation procedure and subsequent C-doped GaN buffer
growth, and ⬃5 nm/ min 共300 W/0.4 SCCM兲 for the undoped N-rich/high-T grown AlGaN/GaN heterostructure on
top. Since the Ga/N flux ratio for the AlGaN/GaN structure
was at ⬃0.75– 0.9,29 the growth rate during the entire structure was nearly identical. A schematic of the typical sample
structure and the thermal dissociation step are illustrated in
Fig. 1.
Reflection high energy electron diffraction 共RHEED兲
was employed to monitor the surface properties throughout
the thermal dissociation procedure at the RI and the film
growth. The effects of these individual growth steps on the
concentration of interfacial impurities 共O, C, and Si兲 and
intentional carbon dopants were further analyzed by secondary ion mass spectroscopy 共SIMS兲. The surface morphologies of the thermally treated GaN templates and the final
AlGaN/GaN heterostructure were characterized by atomic
force microscopy 共AFM兲. The quality of the AlGaN/GaN
heterointerfaces and the Al content in the AlGaN top layer
was further analyzed using high-resolution x-ray diffraction
共HRXRD兲. Finally, test structures were fabricated using contact photolithography, including van der Pauw 共vdP兲 patterns
for measurements of the 2DEG mobilities and test patterns
for the evaluation of isolation buffer leakage currents and
maximum drain currents.
III. RESULTS
A. Thermal cleaning of GaN template
The in situ thermal dissociation procedure applied to the
共0001兲 GaN template required a careful counterbalance be-
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043527-3
Koblmüller et al.
J. Appl. Phys. 107, 043527 共2010兲
FIG. 3. 3 ⫻ 3 ␮m2 atomic force micrographs of MOCVD-GaN templates,
which received 共a兲 no thermal cleaning 共i.e., untreated sample兲, 共b兲 a 10cycle thermal cleaning step, and 共c兲 a 20-cycle thermal cleaning step. The z
scale for all images is 5 nm. Note that the surface pits visible after the
thermal cleaning procedure performed at 780 ° C reveal the underlying
TDDs; however, no significant degradation in surface roughness with increasing number of cleaning cycles was observed.
FIG. 2. 共Color online兲 SIMS profiles of AlGaN/GaN layers on MOCVDGaN template showing 共a兲 high levels of oxygen 共O兲, carbon 共C兲 and silicon
共Si兲 at the RI for standard regrowth conditions after in situ annealing of the
GaN template below 750 ° C; 共b兲 reduced levels of 关O兴 and 关C兴 to below the
detection limit for regrowth subsequent to the proposed thermal cleaning
step consisting of ten cycles at 780 ° C.
tween 共i兲 effective thermal desorption of surface impurities
and 共ii兲 preservation of smooth surface morphology. In general, we found that keeping the GaN template in vacuum at
substrate temperatures below the onset of thermal dissociation, i.e., ⬃750 ° C,12,13 for duration of several minutes to up
to 1 h, resulted in no discernible reduction of the O-, C-, and
Si-related surface impurities. This is evident in the SIMS
profiles of Fig. 2共a兲, showing increased levels of O 共5.5
C 共5.0⫻ 1019 cm−3兲,
and
Si 共2.0
⫻ 1019 cm−3兲,
19
−3
⫻ 10 cm 兲 at the RI between the GaN template and the
top MBE-GaN layer grown at 700 ° C. These large O, C, and
Si concentrations remained also unaffected by etching the
GaN template either in buffered HF or in HCl acids prior to
loading into the MBE chamber.
On the other hand, elevated substrate temperatures of
T ⬎ 760 ° C yielded a characteristic layer-by-layer thermal
dissociation behavior in vacuum evidenced by RHEED intensity oscillations recorded along the 关112̄0兴 azimuth with
periods corresponding to the evaporation of individual MLs
共not shown here兲. According to our previous investigation,
the layer-by-layer thermal dissociation rate obeyed an exponential behavior with substrate temperature.30 However, the
layer-by-layer-like evaporation ceased readily with the appearance of slightly facetted surfaces after the dissociation of
several MLs, in particular for T ⬎ 800 ° C. The critical time
for the onset of facet formation—as observed by RHEED—
was approximately 1.5 min at 800 ° C and 2.5 min at
780 ° C, equivalent to the evaporation of ⬃15– 20 ML of
GaN based on our in situ measurements.30
Utilizing such single-step thermal dissociation procedure
below the limits of facet formation did not reproducibly desorb the surface impurities to below the detection limit of
SIMS. To access lower impurity concentration at the GaN
template surface, we developed an optimized thermal cleaning procedure that consisted of alternating vacuum thermal
dissociation and GaN regrowth cycles at constant 780 ° C.
As depicted in Fig. 1共b兲, after raising the substrate temperature to 780 ° C, the cycles were initiated with 1 min of GaN
growth under slightly Ga-rich conditions, followed by 2 min
of thermal dissociation in vacuum, with closed Ga and N
shutters. This sequence was repeated several times, where
the shorter GaN regrowth step essentially helped to preserve
the surface from increased facet formation. The net growth
rate of the GaN regrowth step was 1.5 nm/min 共i.e., 3.5
nm/min supplied N flux minus 2 nm/min thermal dissociation rate at 780 ° C兲, while the thermal dissociation step delivered a constant GaN desorption rate of 2 nm/min. Thus,
one cycle resulted in the net GaN evaporation of 2.5 nm, i.e.,
⬃10 ML.
The SIMS profile in Fig. 2共b兲 demonstrates the effective
reduction of both oxygen and carbon impurities at the GaN
template interface by using ten cycles of the proposed procedure 共i.e., 25 nm net amount of evaporated GaN兲. Reproducible reduction by three orders of magnitude of both O and
C levels to the detection limit of SIMS 共关C , O兴 ⬃ mid
− 1016 cm−3兲 was achieved for cycle periods of 5, 10, and 20,
corresponding to ⬃12.5– 50 nm thickness etched thermally
from the GaN template surface. Note that the Si impurity
level at the interface was barely affected by this procedure.
This was not surprising due to the extremely low vapor pressure of Si at these temperatures, and strategies to overcome
Si interface contamination will be discussed below.
The surface morphology during the repeated regrowth/
thermal dissociation cycles showed no significant roughening or faceting effects. This was confirmed by the continuously streaky, i.e., two-dimensional 共2D兲, RHEED patterns
and ex situ by AFM scans of the treated GaN template surfaces after immediate cool-down 共Fig. 3兲. Surprisingly, a
comparison with the untreated GaN template surface evidenced no increase in surface root-mean-square 共rms兲 roughness even for increased cycle periods, with typical rms val-
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043527-4
Koblmüller et al.
ues of less than 0.5 nm over a 3 ⫻ 3 ␮m2 area. However, the
step-flow growth features present on the untreated GaN template surface were diminished, and morphologies with 2D
islandlike growth features formed instead. These are clearly
obvious in Fig. 3共c兲, showing a large density of 1 ML 共i.e.,
⬃0.3 nm兲 high islands of arbitrary shape, which is consistent with the preferential 2D layer-by-layer growth and dissociation mode observed under high-T MBE growth.9,12,29,30
Furthermore, the treated GaN template surfaces exhibited several randomly distributed shallow surface pits, with
depths of ⬃1 – 2 nm and densities of 4 ⫻ 108 cm−2 关Fig.
3共b兲兴 and 6 ⫻ 108 cm−2 关Fig. 3共c兲兴. These surface pit densities are identical to the total TDD of the GaN template, confirming the common association of shallow surface pits in
MBE-grown GaN with the surface termination of TDs.28,31
We thus proved that the thermal cleaning procedure preferentially etches along these surface defects, unveiling the underlying TD network, similar to wet chemical or photoelectrochemical techniques.32 Given the similar surface pit
widths and depths in Figs. 3共b兲 and 3共c兲, no correlation between the total amount of GaN evaporated and surface pitsize effects could be identified. This provides, moreover, an
effective thermal etching route of 共0001兲 GaN layers without
significant changes in surface morphology and applicable
even toward larger etch depths.
B. Compensation doping of residual Si
Although the dominant oxygen and carbon impurities
were eliminated from the RI by this procedure, residual Si
with concentrations of 1018 – 2 ⫻ 1019 cm−3 were still
present. These 关Si兴 are expected to yield yet increased buffer
leakage,21,32 as shown further below, limiting the isolation of
the 2DEG at the AlGaN/GaN interface from the parasitic
capacitances at the MBE-GaN/GaN-template RI. To achieve
superior electrical isolation for better rf performance of the
AlGaN/GaN devices, we investigated the use of C-doped
GaN 共GaN:C兲 buffer layers, which were also recently employed for AlGaN/GaN structures on sapphire,33 SiC,26,34
and Si 共111兲 共Ref. 27兲 substrates. As illustrated in Fig. 1共a兲,
the doped GaN:C layer was 400 nm thick and directly grown
on the GaN template after the completion of the thermal
dissociation cycles. Variable CBr4 foreline pressures were
used to identify the carbon incorporation efficiency for the
GaN buffer layers. Also, the growth conditions for the
C-doped GaN layers were varied between the two sophisticated GaN growth regimes, i.e., Ga-rich, low-T 共Ga/ N = 1.7,
T = 680 ° C兲 and N-rich, high-T 共Ga/ N = 0.85, T = 790 ° C兲.
As indicated earlier, the GaN:C buffer layer growth was followed then by undoped ⬃30-nm-thick AlGaN/ 1-␮m-thick
GaN device layer grown under constant N-rich, high-T conditions 共Ga/ N = 0.85, T = 790 ° C兲.
The SIMS data in Fig. 4 show representative C-doping
profiles in relation to the interfacial 关Si兴 for these two buffer
layer growth regimes. Most importantly, carbon incorporation was at least two orders of magnitude higher under
Ga-rich/low-T conditions compared to N-rich/high-T conditions, although the effective buffer layer growth rates were
similar. Also, under the N-rich/high-T conditions, the carbon
J. Appl. Phys. 107, 043527 共2010兲
FIG. 4. 共Color online兲 SIMS profiles of AlGaN/GaN layers on MOCVDGaN template highlighting the C concentration relative to 关Si兴 in C-doped
GaN buffer layers grown under 共a兲 Ga-rich/low-T conditions 共Ga/ N = 1.7,
T = 680 ° C兲 with CBr4 flow of 25 mTorr and 共b兲 N-rich/high-T conditions
共Ga/ N = 0.85, T = 790 ° C兲 with CBr4 flow of 50 mTorr. 共c兲 Comparison of
the C incorporation efficiencies for these two different growth regimes as a
function of CBr4 flow.
concentration remained largely constant over a large range of
CBr4 foreline pressures. We therefore assume that carbon
incorporation is much more sensitive to growth temperature
than the Ga/N flux ratio, as noted also in a recent study based
on lower GaN growth temperatures.35
The buffer leakage was investigated by measurements of
the drain-source I-V curves from isolation patterned AlGaN/
GaN HEMT structures with the 2DEG etched away at the
gate region. Details on our standard AlGaN/GaN HEMT device dimensions and processing steps are reported
elsewhere.4,26 Figure 5 shows the drain-source I-V curves
from isolation patterned samples grown 共a兲 without the thermally etched RI and without the C-doped GaN buffer layer
共standard sample兲, 共b兲 with the RI thermally etched by 25 nm
共i.e., ten thermal dissociation cycles兲, but without the
FIG. 5. Drain-source buffer leakage currents from isolation patterned
samples dependent on different MBE-GaN/MOCVD-GaN-template conditions 共i.e., with and without thermally cleaned RI and with C-doped GaN
buffer layer—the C doping level in this particular sample was ⬃2
⫻ 1019 cm−3 as measured by SIMS兲. The dashed line at 20 V bias is a guide
to the eye for comparison between samples.
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043527-5
Koblmüller et al.
C-doped GaN buffer layer, and 共c兲 with both the thermally
etched RI 共again 25 nm兲 and the C-doped GaN buffer layer.
The carbon concentration 关C兴 in this particular GaN:C buffer
layer was ⬃2 ⫻ 1019 cm−3, as measured by SIMS, although
buffer layers investigated with slightly less or higher 关C兴
gave similar results.
We observed a substantial decrease in buffer leakage
current over several orders of magnitude among these three
different samples. While the standard sample showed relatively high buffer leakage currents of ⬃100 mA/ mm at 20
V, these leakage currents were sequentially reduced to
⬃1 mA/ mm at 20 V for the sample with its RI thermally
etched and even more notably to ⬃1 ␮A / mm at 20 V when
a C-doped GaN buffer layer was additionally introduced. In
comparison, standard samples with additional treatment of
buffered HF or HCl prior to growth 共not shown here兲 gave
similarly high buffer leakage currents as in the case of the
untreated samples. On the other hand, increasing the number
of thermal dissociation cycles by a factor of 2, i.e., increasing
the etching depth at the GaN template to ⬃50 nm, could not
yield further improvement in buffer leakage currents. This
suggests that the residual 关Si兴 at the RI is the main cause of
buffer leakage in samples without C-doped GaN buffer layers. Despite the obvious merits of the thermal etching procedure, it is important to note that sufficiently low buffer leakage currents for device operation 共Ⰶ1 mA兲 require
additional compensation of the residual Si by using C-doped
GaN buffer layers.
J. Appl. Phys. 107, 043527 共2010兲
C. High-temperature AlGaN/GaN growth on GaN
template
FIG. 6. 共Color online兲 10⫻ 10 ␮m2 共1 ⫻ 1 ␮m2 inset兲 atomic force micrographs of AlGaN layers grown on N-rich/high-T GaN buffer layer/
MOCVD-GaN templates at constant T = 780 ° C and an 共Al+ Ga兲 / N flux
ratio of 共a兲 1.6 共Ga-rich兲 and 共b兲 0.75 共N-rich兲. The thickness, Al content,
and surface rms roughness of these AlGaN layers are 共a兲 28 nm, x共Al兲
= 0.39, rms= 0.7 nm; 共b兲 31 nm, x共Al兲 = 0.34, rms= 0.6 nm, respectively.
Note that the surface morphology is largely independent of the 共Al
+ Ga兲 / N flux ratio. 共c兲 SIMS plot showing O and Al profiles across a sequence of ⬃80 nm AlGaN/250 nm GaN layers grown under different
growth conditions, initiated at the GaN template with Ga-rich/low-T conditions 共Ga/ N = 1.8, T = 700 ° C兲 and varied consecutively to Ga-rich/high-T
共Ga/ N = 1.8, T = 780 ° C兲, N-rich/high-T 共Ga/ N = 0.7, T = 780 ° C兲, and
N-rich/low-T 共Ga/ N = 0.7, T = 700 ° C兲 conditions 共constant Al flux of 1
nm/min兲.
Despite the recently highlighted N-rich/high-T growth
regime with optimized surface morphologies and electron
mobilities for homoepitaxial 共0001兲 GaN layers on GaN
templates,9 the effect of these growth conditions on the
growth of AlGaN layers has remained unexplored. Representative AFM images in Fig. 6 illustrate the influence of different 共Al+ Ga兲 / N flux ratios 关in the range of 0.75⬍ 共Al
+ Ga兲 / N ⬍ 1.6兴 on the surface morphology of AlGaN/GaN
layers grown at a constant growth temperature T = 780 ° C.
Note that the varying 共Al+ Ga兲 / N flux ratios cover growth
conditions from N-rich to Ga-rich growth via adjusting both
Al and Ga fluxes such that the resulting ⬃30-nm-thick AlGaN layers had comparable Al content in the range of 0.34
⬍ x共Al兲 ⬍ 0.39. Here, the underlying GaN layer was grown
under an identical Ga/N flux ratio of 0.75, such that the
nucleation of the consecutive AlGaN layers was not obscured by the underlying GaN morphology.
Surprisingly, the AFM images show no discernible differences in morphology among the individual AlGaN
samples, yielding atomically flat surfaces with rms roughnesses typically between 0.6 and 0.8 nm 共10⫻ 10 ␮m2 scan
area兲. Also, all AlGaN surfaces replicate the underlying
TDDs, indicated by the shallow surface pit densities of ⬃5
⫻ 108 cm−2. This wide growth window for atomically
smooth AlGaN layers is in stark contrast to the observed
sensitivity of AlGaN surface morphology and surface defect
structure to the 共Al+ Ga兲 / N flux ratio at lower growth
temperatures.36,37
Moreover, we investigated the incorporation of unintentional oxygen into both GaN and AlGaN layers grown at
high temperatures under both N-rich and Ga-rich conditions.
The SIMS data in Fig. 6共c兲 show the oxygen and Al profiles
共as guide to the eye兲 across an AlGaN/GaN sample structure
共see inset兲 grown under the two distinct high-T 共N-rich versus Ga-rich兲 growth conditions, relative to low-T growth
conditions. Initiating the growth within this structure by
Ga-rich/low-T conditions and changing these sequentially to
Ga-rich/high-T, N-rich/high-T, and N-rich/low-T conditions,
care has been taken to maintain a smooth growth front up to
the typically kinetically roughened N-rich/low-T grown
layer. The layer thicknesses for each individual growth condition sequence were 250 nm for GaN and 80 nm for AlGaN,
while the Al content of the AlGaN layer was adjusted to
x共Al兲 ⬃ 0.2 to prevent large tensile strain formation and
cracking of the AlGaN layers.
A comparison among these AlGaN/GaN layer sequences
indicates that the oxygen levels are very similar among the
individual AlGaN and GaN layers grown under
Ga-rich/low-T, Ga-rich/high-T, and N-rich/high-T condi-
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043527-6
Koblmüller et al.
FIG. 7. 共Color online兲 ␻-2␪ high-resolution XRD scan across the symmetric
共0002兲 reflection for an AlGaN/GaN HEMT structure on the MOCVD-GaNtemplate showing complete overlap of experimental 共dotted兲 and simulated
共line兲 data. Best fits of AlGaN thickness and Al content yielded 24 nm and
x共Al兲 = 0.38. The insets in the upper left corner show representative AFM
images with z scale of 5 nm and surface rms roughness of less than 1 nm.
tions, though oxygen incorporation was slightly higher in
AlGaN than in GaN. The N-rich/low-T growth conditions
yielded significantly higher oxygen levels 共about one order
of magnitude higher兲, which is in agreement with previous
reports38 and will be discussed later.
Due to the independence of the AlGaN surface morphology with the 共Al+ Ga兲 / N ratio at high growth temperatures,
we fabricated a simple AlGaN/GaN-on-GaN-template
HEMT structure using an 共Al+ Ga兲 / N flux ratio of 0.9,
which yielded an AlGaN thickness of 24 nm and an Al content of 0.38 at the given conditions. Both AlGaN thickness
and Al content were calculated based on simulations of the
peak separation between GaN and AlGaN peaks in the ␻-2␪
HRXRD scans under the assumption of a coherently strained
AlGaN layer 共Fig. 7兲. Indicative of the high-quality AlGaN/
GaN heterostructure, we observed thickness fringes up to at
least second order in the HRXRD data. As expected, this was
also proven by the atomically smooth surface morphology
with rms roughness of less than 1 nm 共see inset: AFM image兲. RT Hall measurements on several vdP patterns fabricated on this sample yielded an average 2DEG Hall mobility
of 1760 cm2 / V s at a sheet density of 1.2⫻ 1013 cm−2 and a
sheet resistance of 295 ⍀ / 䊐. We attribute the high mobility
observed in this sample to a combination of decently low
TDD 共mid-108 cm−2兲, good isolation, and low dispersion effects of the 2DEG via reduced impurities at the RI 共using the
thermal dissociation step兲 and the 400-nm-thick C-doped
GaN buffer layer, as well as the improved N-rich/high-T
growth conditions. The high-quality transistor structure was
further corroborated by the dc drain current characteristics
shown in Fig. 8. The saturation drain current was measured
as high as 1.1 A/mm for VGS = +1 V, although self-heating
effects might be responsible for the slight decrease in drain
current at higher voltages. We note that these data compare
favorably with the best devices grown on both GaN
templates,19 freestanding GaN 共Ref. 21兲 and SiC
substrates.4,8,11
IV. DISCUSSION
This work demonstrates that despite the rapid development of ever lower TDDs of self-supporting GaN templates,
J. Appl. Phys. 107, 043527 共2010兲
FIG. 8. 共Color online兲 DC current-voltage curves of the representative
AlGaN/GaN HEMT on the GaN template. The gate voltage was varied
between +1 and ⫺6 V in steps of ⫺1 V.
parasitic conduction at the RI plays a still dominant role for
the intrinsic properties 共current leakage paths, 2DEG Hall
mobilities, and dc characteristics兲 of direct-regrown AlGaN/
GaN HEMT structures. Since both origin and magnitude of
residual free carriers at the near-RI region in these structures
seem to affect the buffer leakage characteristics drastically
共see Fig. 5兲, the aim and design of experiments were therefore threefold: 共i兲 to clarify the possible origin and magnitude of residual free carriers, 共ii兲 to reduce the buffer leakage
via thermal cleaning of the GaN template and via carbon
doping of the GaN buffer, and 共iii兲 to investigate the effect of
high-temperature growth conditions on the incorporation of
unintentional impurities 共specifically oxygen兲.
Although a number of groups have investigated the first
two points,19–27 there is yet no standard route of preparing
the GaN template and eliminating the parasitic conductive
channel at the RI. In the simple thermal 共i.e., dry兲 cleaning
procedure of the GaN template introduced in this study, we
avoided any wet chemical and/or UV/ O3 treatments, as commonly utilized for GaN surface cleaning prior to Ohmic and
Schottky contact deposition.23,39 This was motivated mainly
by the limited reproducibility of these treatments, which
were found to be heavily dependent on the sequence and
solution ratios of the applied chemicals 共mainly HCl-, HF-,
H2SO4-, and H3PO4-based solutions兲 and their frequent
traces of additional elements 共such as Cl and F兲 on the
surfaces.22,24,25,40 Furthermore, UV/ O3 treatments were reported to increase the O concentration on the surface despite
the effectiveness in removing C.24,25,40 Nevertheless, common to most of these studies were the findings that thermal
desorption by high-temperature 共⬎700 ° C兲 annealing in
vacuum, NH3, N2, or N2 / H2 plasmas after wet chemical
and/or UV/ O3 treatments, provided a more effective 共although not complete兲 removal of residual contaminants,
while no comments on the effects on surface morphology
and electrical properties of devices fabricated were made.
In our thermal cleaning method, large requirements were
posed on balancing the effective desorption of surface impurities during the evaporation of GaN and preserving atomically smooth surface morphology. This becomes very impor-
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043527-7
Koblmüller et al.
tant as long-period high-temperature thermal cleaning 共at
temperatures beyond thermal dissociation兲 is known to introduce large damage to the surface, usually making the GaN
material electrically unsuitable.41 We noticed that the critical
time period for thermal dissociation in vacuum before deterioration of the GaN surface 共i.e., ⬃2.5 min at
780 ° C—equivalent to ⬃15– 20 ML of evaporated GaN兲
was insufficient to thermally desorb all residual O and C to
below the detection limit of SIMS. In contrast, the implementation of a short Ga-rich regrowth step or multiple alternating regrowth/dissociation cycles at 780 ° C reduced both
O and C more effectively. In part, this can be explained by
the ready reaction of surface Ga with residual O, forming
highly volatile Ga2O3.38 We note that applying pure Gaassisted thermal cleaning without active nitrogen 共i.e., “Ga
polishing”42兲 might also be sufficient for the desorption of
residual O from the GaN template at 780 ° C; however, this
simultaneously leads to rough surface morphologies due to
the increased thermal dissociation rates under large excess
Ga.30 It is further important to note that the thermal cleaning
procedure was optimized for a fixed N flux of 5 nm/min.
With increases in effective N flux, we expect decreases in the
thermal dissociation rates such that the thermal cleaning step
could shift to much higher temperatures and lead possibly to
more enhanced thermal desorption of both O and C.
Obviously, in Fig. 2, the thermal dissociation cycles also
exposed surface pits associated with underlying TDs. Since
both the surface pit density matched the total TDD of the
underlying GaN template and no discernible differences in
surface pit sizes were observed for various periods of high-T
dissociation/regrowth cycles, no relation of preferential thermal dissociation on the type of dislocation could be elucidated. This is in contrast to the typically observed pit-size
differences and their association with individual types of
TDs in MBE-grown GaN layers at low growth
temperature.31 It is important to note that nearly constant
surface pit sizes were also observed in thick GaN epilayers
grown by MBE either under N-rich9 or Ga-rich conditions at
high growth temperatures of ⬃780 ° C.43 We therefore attribute the appearance of these surface pits to specific surface
kinetics under these high-T conditions and assume that their
sizes might be predefined at the onset of high-T GaN nucleation or during the initial thermal dissociation step. Nevertheless, we expect that both the duration of the thermal dissociation step and the Ga/N flux ratio during the intermediate
regrowth steps are important parameters governing both
widening/facetting and filling mechanisms of surface pits,
respectively. In the extreme case of prolonged thermal dissociation at high enough temperatures, thermal dissociation
along the sidewalls of these surface pits was found to dominate over the 2D layer-by-layer-like dissociation from the
planar 共0001兲 GaN surface, resulting in highly facetted GaN
surfaces.30 During the short intermediate regrowth cycles,
however, the Ga-rich conditions seemed to promote a high
enough growth rate across these pits preventing rapid surface
pitting. The equilibrium between the two counteracting
mechanisms could be related to the same Ga-adlayer medi-
J. Appl. Phys. 107, 043527 共2010兲
ated surface diffusion enhancement, as recently found for the
small pit sizes of 共0001兲 GaN surface morphologies grown
by PAMBE under Ga-rich conditions.28,31,44
With residual Si remaining largely unaffected by the
thermal cleaning procedure 关Fig. 2共b兲兴, the buffer leakage
was not reduced to the ⬃␮A / mm levels ultimately required
for device operation 共Fig. 5兲. As expected, the negligible
thermal desorption of Si can be related to its very low vapor
pressure in the investigated temperature region. This could
be overcome only by significantly larger temperatures
共⬎1100 ° C兲 in combination with pre-oxidized Si 共i.e., SiOx兲
via UV/ O3 or wet chemical etching in HF-based solutions.
We have attempted both these methods prior to thermal
cleaning; however, the removal of residual Si was unsuccessful due to the limited substrate temperature heating
共⬃900 ° C兲 of our MBE system.
To compensate residual Si at the RI effectively, carbon
was introduced as an intentional dopant to realize semiinsulating GaN buffer layers.26,27,33,34 The marked difference
in carbon incorporation and compensation effect within
GaN:C buffer layers grown under either Ga-rich/low-T conditions or N-rich/high-T conditions can be explained twofold. First, the incorporation efficiency of carbon via the
CBr4 source was recently found to be strongly temperature
dependent and much less dependent on the Ga/N flux ratio.35
This is particularly corroborated by the independence of the
generally low C incorporation at variable CBr4 fluences under high-temperature growth conditions, as found in Fig.
4共c兲. Second, the low compensation effect of carbon under
N-rich/high-T GaN growth conditions is not expected to
arise from changes in the amphoteric behavior of carbon in
GaN.45,46 As GaN layers grown under N-rich/high-T conditions show similar unintentional n-type conductivity as those
grown under standard Ga-rich/low-T conditions,9 the formation of substitutional CN as a shallow acceptor should hold
true independent of the given growth conditions.
Another possible effect for low compensation of residual
impurities and higher buffer leakage could be related to
codoping of unintentional oxygen under N-rich conditions,
as evidenced by Green et al.35 However, the SIMS profiles of
N-rich layers grown under the much higher temperatures in
this study did not indicate any measurable increases in oxygen concentration within the GaN:C buffer layers. It is worth
noting that the SIMS profiles in Figs. 4共a兲 and 4共b兲 show
both a distinct low intensity Si peak 共关Si兴 ⬃ 共2 – 3兲
⫻ 1017 cm−3兲 at the beginning and end of the GaN:C buffer
layer, respectively. In the case of the Ga-rich/low-T GaN:C
buffer layer growth, this can be attributed to residual Si being carried by the Ga-rich GaN:C growth surface, leaving a
spike where growth was interrupted and conditions changed
to N-rich/high-T conditions. Thus, residual Si can either diffuse from the RI and be carried on the metal-rich growth
surface or be buried by the N-rich growth surface, such that
the Si from the RI becomes redistributed between the RI
itself and the GaN:C buffer layer. Very similar observations
were found for progressive and inhibited Si diffusion at Alrich versus N-rich AlN surfaces.26
Although the N-rich/high-T growth conditions have now
been widely investigated for GaN epilayers,9,12,29,30 their ef-
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043527-8
J. Appl. Phys. 107, 043527 共2010兲
Koblmüller et al.
fect on the growth of AlGaN layers as presented in this study
requires further discussion. The apparent independence of
the AlGaN surface morphology on the 共Al+ Ga兲 / N flux ratio
at the high growth temperatures of ⬃780– 790 ° C 关at least
for AlGaN layers with Al content of less than 40% and
0.75⬍ 共Al+ Ga兲 / N ⬍ 1.6兴 suspects at first glance that the
growth kinetics mimic those observed for high-T GaN
growth by PAMBE.9,43 The common surface characteristics
given in this wide growth window are consistent with typical
2D layer-by-layer growth features and the lack of stepped
terraces and hexagonal spiral hillocks 共as under low growth
temperatures兲. It has been argued that these 2D surface morphologies achieved under high growth temperature were
mainly due to thermally activated surface diffusion, rather
than adlayer mediated diffusion9 since no surfactant Ga adlayer could be confirmed from in situ RHEED and desorption mass spectrometry analysis.12 Thus, we suggest that for
low-to-intermediate Al-content AlGaN layers, the surface
growth kinetics are dominated by the growth temperature,
while the effects of varying Al flux are negligible, provided
that Al adatoms are fairly mobile at these high growth
temperatures.47
The similarities in high-T growth kinetics between GaN
and AlGaN layers are further corroborated by the incorporation of impurities. Although unintentional oxygen levels
were higher in AlGaN relative to GaN layers due to the
larger affinity of oxygen in Al-containing layers, the nearly
identical oxygen levels within AlGaN layers grown under
Ga-rich/low-T, Ga-rich/high-T and N-rich/high-T conditions
point to the largely independent oxygen incorporation among
these regimes 共i.e., oxygen sticking coefficient= 1 for AlGaN兲. This can be explained either by Ga-adlayer mediated
growth for Ga-rich conditions38 or due to substantial thermal
desorption at temperatures of thermal dissociation
共⬎750 ° C兲 even for N-rich growth. Hence, oxygen levels
were only increased under N-rich/low-T conditions, where
both suggested mechanisms were absent.38
Considering the similarities of smooth morphology and
low oxygen levels in AlGaN layers grown under high-T conditions with those grown under standard Ga-rich/low-T conditions has enabled the fabrication of high mobility 2DEGs
and low sheet resistances, even for fairly high Al-content
AlGaN/GaN HEMT structures. With the Al content in the
AlGaN cap layer varying at x共Al兲 ⬃ 0.34– 0.38, and respective sheet carrier densities of ⬎1 ⫻ 1013 cm−2 in the 2DEG,
these parameters were rather high for producing optimized
2DEG mobilities at the given TD density. We therefore suggest that the achieved 2DEG mobilities of ⬎1750 cm2 / V s
共at 300 K兲 were dominated by alloy or interface scattering,
rather than by reaching the mobility limit via charged dislocation scattering. The latter becomes typically much more
dominant in the low sheet carrier density regime, i.e., at
nsh ⬍ 3 ⫻ 1012 cm−2, as calculated for the TD densities of
⬃5 ⫻ 108 cm−2 in our AlGaN/GaN layers.48 According to
these numbers, we would expect significantly higher mobilities for sheet carrier densities adjusted to the low-to-mid
1012 cm−2 region.
V. CONCLUSION
This work addressed several critical issues related to the
PAMBE regrowth of AlGaN/GaN HEMT structures on
共0001兲 GaN templates contaminated with typical residual impurities 共O, C, and Si兲 at the RI. Utilizing an optimized in
situ thermal cleaning procedure consisting of thermal dissociation and regrowth cycles at temperatures in the thermal
decomposition regime of GaN 共i.e., 780 ° C兲, O and C impurities were effectively desorbed from the RI. This cleaning
step preserved an overall very smooth surface morphology
while exposing only shallow surface pits associated with the
typical TDs piercing through the GaN template interface.
Due to the very low vapor pressure of Si impurities, their
tenacity and n-type conductive behavior at the RI could only
be compensated by intentionally doped GaN:C buffer layers
for proper isolation and ultimate reduction of buffer leakage.
Further investigations of AlGaN/GaN device layers grown
on top under the recently established high-T PAMBE growth
conditions yielded several promises for low background impurity with large flexibility in AlGaN growth window. While
both the high-T GaN and AlGaN layers seemed to incorporate fairly low levels of unintentional oxygen 共similar to the
widely used Ga-rich/low-T conditions兲, the AlGaN surface
morphology 共structure and rms roughness兲 was found independent of the 共Al+ Ga兲 / N flux ratio, at least from
N-rich/high-T to intermediate Ga-rich/high-T growth conditions. 300 K 2DEG mobilities and sheet resistances of such
N-rich/high-T-grown AlGaN/GaN heterostructures were
⬎1750 cm2 / V s and 295 ⍀ / 䊐, respectively 共at not yet optimized sheet carrier densities of ⬃1 ⫻ 1013 cm−2兲, with
maximum dc drain currents of ⬃1.1 A / mm at a +1 V bias.
We expect that with further adjustments of the sheet carrier
density, much higher electron mobilities can be achieved toward the limit where charged dislocation scattering will become dominant within these fairly low TDD AlGaN/GaN
heterostructures.
ACKNOWLEDGMENTS
The authors thank Charles Evans Analytical Group and
Dr. T. Mates 共UCSB兲 for SIMS measurements and Dr. David
Storm, U.S. Naval Research Laboratories for helpful discussions. This work was supported by USAFOSR 共Kitt Reinhardt, Program Manager兲 and DOE SSL Project No. DEFC26-06NT42857. Further support was provided by ONR
through the DRIFT MURI 共Dr. Paul Maki, Program Manager兲.
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