JOURNAL OF APPLIED PHYSICS 108, 093501 共2010兲
Generation of amber III-nitride based light emitting diodes by indium rich
InGaN quantum dots with InGaN wetting layer and AlN encapsulation
layer
C. B. Soh,1,a兲 W. Liu,1 S. J. Chua,1,2,b兲 S. S. Ang,1 Rayson J. N. Tan,1 and S. Y. Chow1
1
Institute of Materials Research and Engineering, Agency for Science, Technology and Research
(A*STAR), 3 Research Link, Singapore 117602
2
Singapore-MIT Alliance, National University of Singapore, 4 Engineering Drive 3, Singapore 117576
共Received 5 July 2010; accepted 28 August 2010; published online 2 November 2010兲
Indium rich InGaN nanostructures grown by metalorganic chemical vapor deposition were
incorporated in InGaN/GaN quantum wells for long wavelength generation. These results were
achieved by optimizing the growth temperature of the nanostructures, InGaN quantum well, the AlN
capping layer and the GaN barrier layers. Before the growth of nanostructures, a thin InGaN wetting
layer was included to reduce the lattice mismatch as well as to enhance the deposition of
indium-rich InGaN nanostructures These individual quantum wells were each subsequently capped
with an AlN layer which better preserved the In-rich phase in the nanostructures and prevented the
indium interdiffusion between the InGaN/GaN heterojunctions. The AlN capping layer also reduces
the effect of piezeoelectric field in the active layers of the light emitting diodes as seen from the
reduction in the blueshift in the electroluminescence peaks with higher injection currents. The
energy band profile of such a structure is discussed. © 2010 American Institute of Physics.
关doi:10.1063/1.3499618兴
I. INTRODUCTION
In the past decade, there has been a considerable effort in
research and development into white InGaN light emitting
diodes 共LEDs兲 due to it potential saving for lighting applications. Commercial white LEDs make use of blue LEDs
coated with luminescent materials such as yellow emitting
phosphor Y3Al5O12 : Ce3+ 共YAG:Ce兲,1 CdSe-ZnS,2 or capped
CdSe quantum dots 共QDs兲 with Sr3SiO5 : Ce3+, Li+
phosphor.3 However, it is difficult to achieve a high color
rendering index 共of higher than 85%兲 due to their red spectral deficiency.4,5 There is also the drawback due to additional coating step for the light converter. Recently, there
have been reports of phosphor free white LEDs using prestrained InGaN wells,6 combination of quantum wells 共QWs兲
with varying thickness7 and indium phase separated multiple
quantum wells 共MQWs兲 as the active layers.8 However, the
internal quantum efficiency of the long wavelength component was relatively low, leading to drastic degradation in its
emission efficiency after fabrication9 with the issue of quantum confined Stark effect not addressed. To realize superior
device performances, it is desirable to introduce QDs or
nanostructures into the active layers of the devices.10 Self
assembled QDs structures or indium rich InGaN nanostructures have been known to increase the recombination efficiency of the emitting layers.11,12 The QDs generated by the
self assembly process generally have a random distribution
in terms of its size12 which causes the inhomogeneous broadening of the electroluminescence 共EL兲 spectra. Chen et al.13
has reported on the growth of InGaN nanostructures with
homogeneous size using electron beam lithography techa兲
Electronic mail: cb-soh@imre.a-star.edu.sg.
Electronic mail: elecsj@nus.edu.sg.
b兲
0021-8979/2010/108共9兲/093501/5/$30.00
nique to pattern a template. However, this technique is not
feasible for large area patterning of a 2 in. LEDs wafer as it
is time consuming. Furthermore, LEDs have to undergo subsequent p-GaN growth and annealing for Mg activation
which is carried out at a temperature much higher than that
for the growth of indium rich InGaN QDs. This will lead to
unavoidable spectra broadening and blueshift in peak wavelength.
Wang et al.14 has reported on the effect of increasing the
annealing temperature for p-GaN activation in InGaN QDs
based LEDs. With the annealing temperature at 830 ° C, degradation of the QDs takes place through atomic diffusion of
indium into the underlying wetting layer, leading to the formation of a wide InGaN QWs with higher indium contents.
These effect enhances the blueshift in the emission wavelength at a higher drive current.15 Due to the constraint in the
materials growth properties, we have modified the growth
structures for the active regions to preserve the indium rich
nanostructures formation by preventing out-diffusion of indium at the high temperature thermal annealing stage. Monolayers of InGaN wetting layer is first grown to reduce the
strain for the incorporation of self-assembled indium rich
InGaN QDs. This is similar to the use of an InGaAs wetting
layer to grow self assembled arsenide QDs.16 A thin aluminum nitride 共AlN兲 layer is then incorporated to serve as a
capping layer after the growth of the indium rich InGaN
QDs. AlN capping mechanism has been studied in detail for
GaN QDs/AlN system.17,18 This work will focus on the study
of the mechanism for self assembly of indium-rich InGaN
QDs on InGaN wetting layer, followed by capping with a
thin AIN layer. The band structure which leads to generation
of amber emission from these LEDs structures will be proposed.
108, 093501-1
© 2010 American Institute of Physics
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J. Appl. Phys. 108, 093501 共2010兲
II. EXPERIMENTS
The sample used in this study is grown by metalorganic
organic chemical vapor deposition. c-plane sapphire was
used as the substrate with hydrogen as the carrier gas. After
the growth of the 25 nm low temperature 共LT兲-GaN buffer
layer at 525 ° C, a 1.2 ␮m thick undoped GaN, and 1.5 ␮m
thick Si-doped GaN were grown. These layers serve as the
template for the growth of MQWs. Three MQWs samples
were grown on the GaN template and are labeled as A, B,
and C.
In sample A, three periods of InGaN/GaN QWs were
grown with dimensions of 4 nm/12 nm, respectively. Each
well consists of a wetting InGaN layer, 1 nm thick grown
with trimethylindium 共TMIn兲 flow rate at 24 ␮mol/ min and
trimethylgaldium 共TMGa兲 at 12 ␮mol/ min, designated as
␣-layer, followed by 2 nm of InGaN growth with TMIn flow
rate kept at 70 ␮mol/ min and TMGa at 12 ␮mol/ min, designated as ␤-layer. The InGaN well layer, i.e., the ␣-layer
and ␤-layer, were grown at a temperature of 725 ° C. Sample
B, consists of a set of 3 QWs as well. However, after the
growth of the wetting layer, i.e., the ␣-layer, the sample underwent TMIn treatment10,19,20 where TMIn was allowed to
flow under NH3 rich condition for 30 s to form the indium
rich InGaN QDs before the growth of the ␤-layer. For
sample C, growth condition for the QWs was similar to
sample B except for the additional AlN cap layer of
⬃1.2 nm thick grown after deposition of the ␤-InGaN layer.
The three MQWs samples were then capped by a 10 nm GaN
layer. These three MQWs structures were used for subsequent fabrication of LEDs by depositing a 25 nm thick
p-Al0.25Ga0.75N layer and a 200 nm thick p-GaN layer at
870 ° C. An in situ chamber annealing for 20 min in nitrogen
ambient at 825 ° C was carried out for the LEDs sample to
activate Mg in its pGaN layer. The LEDs sample was labeled
as sample X and Y using the MQWs structures of samples B
and C, respectively.
A. Surface morphology study by atomic force
microscopy „AFM…
In order to simulate the effect of possible degradation of
the QDs during the pGaN activation process, in situ chamber
annealing in N2 ambient was carried out for 5 min for QW
sample, B and C. Figure 1 shows the tapping mode AFM
image of the surface morphology of the QWs sample. Figure
1共a兲 shows that the formation of indium rich InGaN QDs on
the InGaN wetting layer after TMIn treatment. The use of N2
as contrast to H2 as a carrier gas for growth of the QDs
facilitate indium incorporation as it slows down the diffusion
of indium adatoms. The reduction in the atomic diffusion of
indium enhances the Stranski–Krastanow growth of InGaN
QDs. Coupled with the higher continuous flow of NH3 at 18
slm 共slm denotes standard liters per minutes兲 as compared to
conventional InGaN well growth 共12 slm兲, a higher amount
of N interacts and forms bonds with impinging indium at Ga
sites of the InGaN wetting layer. This explains the formation
of indium rich InGaN QDs of size ⬃20 nm in diameter and
⬃1.5 nm in height with a density of ⬃1 ⫻ 1010 cm−2. The
InGaN QDs were then capped with a 1.5 nm InGaN well
FIG. 1. 共Color online兲 AFM image of 共a兲 indium rich InGaN QDs deposited
on the InGaN wetting layer, 共b兲 single QW sample B with clusters of InGaN
QDs formed by layer growth of the InGaN well layer which wets the QDs,
共c兲 single QW sample C with stripelike AlN islands on the InGaN well
formed with addition of AlN capped layer, 共d兲 second QDs layer with higher
density grown on single QWs sample C, 共e兲 sample B with annealing at
820 ° C for 10 min, and 共f兲 sample C with annealing at 820 ° C for 10 min.
layer with the NH3 flow rate reduced to 12 slm and the
temperature increased to 725 ° C to promote layer by layer
growth. Figure 1共b兲 shows the AFM image of sample B after
the growth of the InGaN well layer, ␤ layer, with the embedded InGaN QDs. The InGaN QDs clustered and formed
nanostructures due to layer growth of the InGaN over the
InGaN QDs. In order to mitigate the effect of the degradation
of these indium rich QDs with further deposition and annealing for the pGaN layer at a higher temperature, a 1.2 nm AlN
was deposited at 780 ° C to serve as a capping layer on the
␤-InGaN. Figure 1共c兲 shows that the AlN cap leads to the
formation of stripelike AlN islands on the InGaN well
共␤-layer兲 which consisted of the embedded InGaN QDs.
GaN QDs capped by AlN displayed the unique characteristic
of being insensitive to In out-diffusion.17,18 This is attributed
to the formation of stable Al–N bonds and also to the relatively lower mobility of Al adatoms on the film surface at a
lower chamber temperature of 780 ° C. When AlN is first
deposited, it will first wet the surface of the InGaN nanostructures with its first few monolayers and subsequent AlN
grows preferentially to enlarge the nanostructures. When
AlN is grown to a thickness of 1.2 nm, it formed elongated
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J. Appl. Phys. 108, 093501 共2010兲
FIG. 2. 共Color online兲 HR-STEM image taken along 共21̄1̄0兲 direction for
共a兲 sample B with embedded InGaN QDs in the InGaN well layers, 共b兲
sample B after thermal annealing which illustrated the out-diffusion of QDs
to the surrounding matrix and interdiffusion into the GaN barrier causing the
shrinkage of the QDs dimension, and 共c兲 sample C after thermal annealing
with the QDs better preserved but with some out-diffusion into the InGaN
active well region.
islands to coalesce the InGaN nanostructures. Aluminum
also has a lower vapor pressure as compared to indium and
this can reduce the diffusion length of other group III atoms.
This implies that the growth of the subsequent GaN barrier
and InGaN wetting layer can be affected by the base AlN
layer. There will be a reduction in the diffusion length of In
atoms on the first QW layer leading to the increase in the QD
density as shown in Fig. 1共d兲 compared with Fig. 1共a兲. The
QDs density for the second repeat layer is estimated to be
3 ⫻ 1010 cm−2 as compared to a QDs density of ⬃5
⫻ 109 cm−2 for a single well layer. The results suggest that
there is substantial reduction in strain for the embedded InGaN in the second repeat layer.
After annealing for 10 min in N2 ambient at 825 ° C, it
was observed that there is a change in the surface morphology for sample B and C. In sample B, the nanostructures of
the QW formed through clustering of the QDs experienced
further agglomeration to give rounded islands as shown in
Fig. 1共e兲. As for sample C with AlN cap layer, the elongated
structure further bunched together to form connected islands
with dimension of ⬃120– 220 nm. The observed pits separating the structures in the layer are also significantly wider
and deeper. This enables the formation of stable QW elongated structures as observed in Fig. 1共f兲.
B. Transmission Electron Microscopy „TEM… analysis
In order to explore the impact of annealing on Mg activation in the p-GaN layer on the optical properties of the
QDs embedded LEDs, a detailed investigation by TEM was
carried out as illustrated in Fig. 2. In situ chamber annealing
was carried out for LEDs sample, X and Y, at 820 ° C for
FIG. 3. 共Color online兲 PL spectra at 10 K for 共a兲 sample C and 共b兲 sample
B before and after thermal annealing for 10 min in N2 ambient.
activation of Mg in the p-GaN layer of the LEDs. As shown
in Fig. 2共b兲, for LEDs sample X after chamber annealing,
there is a shrinkage in the QDs size with diffusion of indium
atoms out of the QDs to the underlying InGaN wetting layer.
This caused the InGaN well 共␤ layer兲 to form InGaN clusterlike structure as compared to the more uniform well layer
of LEDs sample X before in situ chamber annealing as
shown in Fig. 2共a兲. The QDs appears to form small dotlike
structures and this shrinkage in size can contribute to the
blueshift in emission wavelength for the LEDs. Chuo et al.21
has reported that the density of indium rich dotlike structures
tends to decrease after thermal annealing due to out-diffusion
of indium into the surrounding matrix including the GaN
barrier and the InGaN QWs. With the addition an AlN cap
layer, we observed from TEM image in Fig. 2共c兲 that the
InGaN QDs in the InGaN well are better preserved as there
is minimal out-diffusion of the InGaN nanostructures. The
capping mechanism of GaN QDs by AlN capping layer has
been demonstrated by Coraux et al.17 It is observed that the
capping of InGaN nanostructures in the QWs by AlN carried
out in this work shows a similar characteristic in that its
involves the wetting of InGaN nanostructures by AlN followed by intervalley filling where preferential AlN growth
occurs.
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J. Appl. Phys. 108, 093501 共2010兲
FIG. 4. 共Color online兲 Current dependent EL spectra for 共a兲 LEDs X and 共b兲 LEDs Y with chamber annealing for 20 min in N2 ambient at 820 ° C for Mg
activation. 共c兲 Plot of the light-output power and peak emission wavelength of LEDs as a function of injection current.
C. Photoluminescence „PL… from QDs/InGaN MQWs
TEM images show that with annealing, there is shrinkage in the size of the QDs with out-diffusion of indium from
the In-rich InGaN QDs into the InGaN well and interdiffusion into GaN barrier. With the addition of the AlN capping
layer, the InGaN QDs are better preserved. Rapid thermal
annealing for 10 min at 820 ° C in nitrogen ambient was
carried out for QWs, sample B and C to simulate the effect
of chamber annealing during Mg activation of LEDs.
Figure 3 shows a typical low temperature PL at 10 K for
samples B and C before and after thermal annealing process.
There is a blueshift in the PL emission for sample B whereas
for sample C, a redshift in its PL emission is observed. The
shift in the PL emission can be attributed to following reasons. First, there exists a strong piezoelectric field in InGaN/
GaN QW which is further enhanced in InGaN/GaN QDs
filled QWs 共Ref. 22兲. The strain in the mismatched layers,
reduces the diffusion length of the QDs and enhanced their
incorporation as well as its density. During thermal annealing
at 820 ° C, there is a reduction in the strain of InGaN 共well兲
active region and hence the indium in the QDs starts to diffuse with greater ease. The out-diffusion of indium enhances
the indium content in the QW region 共i.e., the InGaN wetting
and well layer兲. As shown in Fig. 3共a兲, after thermal annealing of sample C 共the QW sample with the AlN cap layer兲, the
PL spectra of the dominant peak associated with the InGaN
wetting and the InGaN well layer are redshifted by 15 nm
due to out-diffusion of indium from the indium-rich QDs.
For sample B without the AlN cap layer, out-diffusion of the
In-rich phase into the surrounding matrix and the interdiffusion between the InGaN/GaN heterojunctions would occur
simultaneously. With longer duration of annealing, indium
diffusion into the GaN barrier, leading to a lower indium
FIG. 5. Schematic representation of the energy band profile of strained
AlN/InGaN QDs/InGaN/GaN QW system.
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093501-5
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Soh et al.
content in the well. This accounts for the blueshift in the
dominant peak position for the well and the wetting layers
for sample B as illustrated in Fig. 3共b兲.
D. EL measurement
Figures 4共a兲 and 4共c兲 shows the injection current dependence of the EL spectra for LEDs X and Y with MQWs of
sample B and C, respectively. The EL spectra for LED X in
Fig. 4共a兲 shows a yellow emitting LED with a emission band
ranging from 500 to 680 nm and a dominant peak at
⬃562 nm at 70 mA. The peak intensity saturated with an
injection current of 70 mA. The dominant peak wavelength
shifts from 582 to 562 nm with the increase in injection
current from 30 to 70 mA as illustrated in Fig. 4共b兲. On the
otherhand, for LEDs Y, its EL spectra shown in Fig. 4共c兲
have an emission band ranging from 550 to 750 nm with a
shoulder emission 共at 500 nm兲 attributed to the InGaN wetting layer. Unlike LED X, there is a no prominent shoulder
peak is observed other than the slight broadening at its lower
wavelength spectrum. The minimal out-diffusion of the indium rich QDs enables the composition and thus the emission wavelength of InGaN wetting layer to be retained during the p-Mg activation step. The EL spectrum with injection
current of 80 mA onwards shows a splitting of the dominant
peak into peak A and B. Peak A is likely due to InGaN well
while Peak B is due to the high indium QDs content nanostructures leading to its higher emission wavelength. As
shown in Fig. 4共d兲, saturation of its peak emission was not
encountered. However, there is slight shift in its dominant
peak from 668 to 658 nm as injection current increased from
30 to 70 mA. The minimal shift also suggests that the effect
of piezeoelectric field in the QWs layer is smaller in LED Y.
Thus the AlN capping layer has not only prevented outdiffusion of indium in the QDs but also reduces the strain
between the InGaN well and GaN barrier layer for the second and the third periods of the QW growth.
E. Energy band profile
Figure 5 gives the schematic representation for the polarizations and the energy band diagram of the strained InGaN wetting layer/InGaN QDs/InGaN well/AlN/GaN QWs
structure. In a polar Ga-face 共0001兲 material, negative and
positive spontaneous polarization 共Psp兲 charges are induced
at the top and bottom interfaces of the wurtzite material system. 共Al, InGaN, Ga兲N is a polar material and the polar
共0001兲 axis is parallel to the QW growth direction. As the
InGaN layer is compressively strained, piezoelectric polarization charges are induced in opposite direction to the Psp
charges.23 On the otherhand, for the sample with AlN capping layer, it is under tensile strain as the Al atomic size is
much smaller than Ga and In. This will assist in lowering the
piezoelectric polarization charge accumulation at the interface of the InGaN well. This accounts for the much smaller
blueshift in the EL emission 共by 10 nm兲 for the amber LED,
Y, with the AlN cap layer as compared to yellow LED, X,
共by 20 nm兲 as the injection current increases from 30 to 70
mA.
III. CONCLUSIONS
In summary, we report on the generation of amber emitting LEDs with minimal shift in its EL spectrum under
higher injection current. This is attributed to implementation
of an AlN capping layer on top of the InGaN well with
embedded indium rich QDs. The AlN cap layer enables the
LEDs to undergo an annealing process with a longer duration
and at a higher temperature. The growth of the AlN cap layer
involves the initial wetting of the InGaN nanostructures with
AlN followed by intervalley filling of a uniform AlN layer.
TEM image revealed that the InGaN QDs in the InGaN well
with the addition of a stable thin AlN layer are better preserved as there is minimal out-diffusion of indium the InGaN
nanostructures. In addition, the AlN cap layer assists in lowering the piezoelectric polarization charge accumulation at
the interface of the InGaN well with GaN barrier layer.
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