Effect of Native Defects on Optical Properties of InxGa1-xN Alloys
S. X. Li and E.E. Haller,
Materials Sciences Division, Lawrence Berkeley National Laboratory, and Department of
Materials Science and Engineering, University of California, Berkeley, California 94720
K. M. Yu, W. Walukiewicz, J. W. Ager III, J. Wu, and W. Shan,
Materials Sciences Division, Lawrence Berkeley National Laboratory, Berkeley,
Hai Lu and William J. Schaff
Department of Electrical and Computer Engineering, Cornell University, Ithaca, New
in InxGa1-xN ternary alloys can be controlled using high energy 4He+ irradiation. The blue
The energy position of the optical absorption edge and the free carrier populations
shift of the absorption edge after irradiation in In-rich material (x > 0.34) is attributed to
the band-filling effect (Burstein-Moss shift) due to the native donors introduced by the
irradiation. In Ga-rich material, optical absorption measurements show that the
irradiation-introduced native defects are inside the bandgap, where they are incorporated
as acceptors. The observed irradiation-produced changes in the optical absorption edge
and the carrier populations in InxGa1-xN are in excellent agreement with the predictions of
the amphoteric defect model.
PACS numbers: 78.66.Fd, 78.40.Fy
Electronic Mail: firstname.lastname@example.org
Since the discovery of its narrow (0.7 eV) bandgap [1, 2], InN has attracted
intense research interest. Interestingly, the actual value of its direct bandgap has been the
subject of some controversy [3-8]. In most cases, measurements of apparent band gaps
larger than 0.7 eV can be attributed to the significant Burstein-Moss shift (conduction
band filling) that occurs at high electron concentrations . Because it has the highest
electron affinity (~5.8 eV) in known semiconductors, InN is extremely susceptible to n-
type doping, either by native donors or by impurities such as oxygen. The relatively small
electron effective mass and high energy position of other conduction band valleys (e.g. L
and X) make the conduction band-filling effect particularly strong in InN. For example,
for an electron concentration in the mid-1020cm-3 range, the absorption edge in InN shifts
to ~1.7 eV .
We have shown previously  that native donors and/or acceptors introduced
through energetic particle irradiation can be used to precisely control the free electron
concentration in In1-xGaxN alloys. Here, we show that the defects have a profound effect
on the optical absorption properties as well. Our results suggest that energetic particle
irradiation provides a simple and convenient way of altering the doping and optical
absorption characteristics of InxGa1-xN alloys in a controllable manner. This technique,
which is applicable to all group-III nitride alloys and could be performed potentially with
a focused ion beam, could open up novel applications of these alloys.
Epitaxial InN and InxGa1-xN thin films (310-2700 nm thick) used in this study
were grown on c-sapphire substrates by molecular beam epitaxy (MBE) with a GaN
(~200 nm thick) buffer layer . The initial free electron concentrations in these
samples ranged from the low 1018 cm-3 to low 1017 cm-3 and the mobility ranged from 7
cm2/V·s (x = 0.24) to above 1500 cm2/Vs (x = 1). A GaN sample (3 µm thick) grown by
metal-organic chemical vapor deposition (MOCVD) with an initial electron concentration
of 7.74×1017 cm-3 and a mobility of 189 cm2/V·s was also used in this study.
The samples were irradiated with 2 MeV 4He+ ions generated by a Van de Graaff
accelerator, with fluences between 1.12×1014 and 2.68×1016 cm-2. In all cases, the particle
penetration depth exceeded the film thickness, assuring homogeneous generation of
defects in the film. Ion channeling spectroscopy showed that the minimum yield χ
increased from 0.04 in an as-grown InN sample to only 0.11 after 4He+ irradiation with a
dose of 1.8×1016 cm-2, indicating that the InN film remains single crystalline. Cross-
sectional transmission electron microscopy of an InN sample subjected to the heaviest
dose of irradiation showed that no additional extended defects were formed, indicating
that point defects are responsible for the observed changes in electrical and optical
properties of the irradiated samples. The optical absorption measurements were
performed at room temperature using a CARY-2390 NIR-VIS-UV spectrophotometer.
Free electron concentration and mobility were measured at room temperature using a
home-built Hall effect system with a magnetic field of 3000 Gauss.
Figure 1 shows the evolution of the optical absorption spectra of InN and
In0.4Ga0.6N with increasing irradiation dose. In both samples, the absorption coefficient is
on the order of 5x104 cm-1 at 0.5 eV above the absorption onset, which is typical for
direct bandgap semiconductors. With increasing irradiation dose the absorption edges
show a blue shift. More specifically, the absorption edge shifts to higher energy, while
the baseline and the slope above the absorption edge remain unchanged. The shift is
composition-dependent, with smaller shifts found in the samples with higher Ga content.
From as-grown to the highest dose of 4He+ irradiation of 2.68×1016 cm-2 the absorption
edge shifted by 1.05eV in InN, by 0.71 eV in In0.7Ga0.3N (data not shown), but by only
0.15 eV in In0.4Ga0.6N. The shift is also observed to saturate with radiation dose. For
example, as seen in Fig. 1 (a) for InN, the blue shift slows down as the irradiation dose
increases and eventually becomes insensitive to further irradiation at a sufficiently high
dose (typically > 1016 cm-2). The behavior of the more Ga-rich (x < 0.34) material is
different. As illustrated in Fig. 2, irradiation of GaN does not affect the fundamental
absorption edge energy at ~3.4 eV but rather produces a new sub-bandgap absorption
feature at ~2.7 eV. Both the strength and the linewidth of the absorption peak increase
with increasing irradiation dose. Clearly unfilled or partially-filled defect states are
formed inside the bandgap of GaN as a result of the irradiation.
Changes in the electronic properties of the irradiated InxGa1-xN samples, the
details of which have been reported separately , cause the changes in the optical
properties. Despite having similar initial free electron concentrations, the InxGa1-xN
samples with different compositions behave very differently after irradiation. In InxGa1-
xN with x > 0.34 the electron concentration increases with increasing 4He+ irradiation
dose and eventually saturates at high (> 1016 cm-2) doses. On the contrary, in Ga-rich
InxGa1-xN the free electron concentration decreases with increasing irradiation dose. In
the case of GaN, the electron concentration decreased so quickly in the fluence range
used that the sample became too resistive to be measured. Figure 3 shows the free
electron saturation concentration of InxGa1-xN at high dose of irradiation as a function of
Ga fraction . The saturation concentration decreases from 4.1×1020 cm-3 in InN to
6.5×1014 cm-3 in In0.24Ga0.76N.
Both the optical and the electrical properties can be understood using the
amphoteric defect model (ADM) [12, 13]. The model predicts that in all semiconductors
the native defects that are highly localized in nature have a common energy level, which
is termed the Fermi-stabilization Energy (EFS) and is located ~4.9 eV below the vacuum
level. In the inset of Figure 3, the position of the EFS is shown together with the
composition dependence of the valence and conduction band edges (VBE and CBE) of
InxGa1-xN. EFS is located ~0.9 eV above the CBE of InN but ~0.7 eV below the CBE of
GaN (2.7 eV above the VBE). According to the ADM either donor- or acceptor-like
native point defects are formed depending on the relative position of the Fermi level (EF)
to the EFS. In InN the EFS is above the CBE; as a result, donor-like irradiation-induced
point defects form. Therefore, as the irradiation dose increases, the electron concentration
increases until EF approaches EFS. At this point both donor and acceptor-like defects are
formed at similar rates and compensate each other leading to stabilization of EF and
saturation of the electron concentration. As a result, a large increase and then a saturation
in the Burstein-Moss shift of the optical absorption edge is predicted; this is in fact
observed in Fig. 1(a).
With increasing Ga content in InxGa1-xN alloys, the CBE moves towards EFS,
resulting in smaller electron saturation concentrations and absorption edge shifts. For x <
0.34 the CBE moves above EFS and acceptors become the dominant irradiation-induced
defects in n-type samples. For Ga-rich material, the acceptor defects compensate n-type
conductivity until the Fermi energy stabilizes at EFS. In the case of GaN, where the
compensation is the most effective, the sample eventually became semi-insulating after
irradiation. As expected, the decreasing electron concentration in Ga-rich In1-xGaxN does
not have a significant effect on the absorption edge. However, the formation of the native
defect states, which are predicted by the ADM to occur at 2.7 eV above the VBE in GaN
(inset of Fig. 3), is observed clearly by optical absorption in Fig. 2.
To demonstrate the agreement between the optical and electronic properties of
irradiated InN, the absorption spectra are numerically analyzed to obtain the Fermi
energy. To account for broadening effects a Gaussian function was convoluted with the
energy dependent absorption coefficient for a direct gap,
exp1 ) '(
where α0(E') is the ideal absorption of InN and ∆ is the Gaussian broadening parameter.
The best fits, which are plotted in Fig. 1(a) as solid lines, are obtained by adjusting ∆ and
the Fermi energy (EF). The corresponding EF values are labeled by arrows; the ∆ values
do not vary significantly and lie consistently between 0.21 - 0.23 eV. The EF values of
InN and InxGa1-xN alloys derived from absorption spectra show excellent agreement with
the electron concentrations measured by room-temperature Hall effect, from a different
set of irradiated samples. The consistency and repeatability suggest that the irradiation is
a dependable method to control the doping and optical properties of InxGa1-xN alloys.
In principle, the same irradiation effect, which is ultimately the displacement of
lattice atoms, can be achieved in all group III-nitride alloys by using an equivalent dose
of any other energetic particles such as electrons, protons, and other ions. By focusing the
particle beam, one can alter the doping or optical properties of group III-nitride alloys
with desired patterns down to the nanometer scale. This offers a range of possible
applications of this technique for fabrication of highly conducting spatially confined
In conclusion, we have shown that optical absorption properties of InxGa1-xN
alloys can be controlled by high energy particle irradiation. As predicted by the
amphoteric defect model, native point defects introduced by irradiation are incorporated
as donors in In-rich InxGa1-xN (x > 0.34) but as acceptors in Ga-rich InxGa1-xN (x < 0.34).
The native donors increase the electron concentration and cause a blue shift of the
absorption edge in In-rich In1-xGaxN while the native acceptors lower the electron
concentration and form defect states in side the bandgap of Ga-rich InxGa1-xN, as
observed by optical absorption.
We thank Milton Yeh of Blue Photonics Inc. for providing the GaN samples and
Z. Liliental-Weber for the TEM results. This work is supported by the Director, Office of
Science, Office of Basic Energy Sciences, Division of Materials Sciences and
Engineering, of the U.S. Department of Energy under Contract No. DE-AC03-
76SF00098. The work at Cornell University is supported by ONR under contract NO.
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Fig. 1: The evolution of InN and In0.4Ga0.6N absorption spectra with increasing doses of 2
MeV 4He+ irradiation. In Fig. 1 (a), the experimental results of InN are shown as data
points and the numerical fit of the optical absorption spectra (see text) are shown as solid
lines. The Fermi energies obtained from the fitting are marked by the arrows.
Fig. 2: The absorption spectra of GaN with increasing doses of 2 MeV 4He+ irradiation.
Fig. 3: The saturation concentration of InxGa1-xN (1 ≥ x ≥ 0.24) as a function of Ga
fraction . The data points are experimental results while the solid line is a numerical
fit. In the inset, the band offset diagram of In1-xGaxN is shown with the energy level of
Fig. 1 of 3, the following 2 figures are the individual Fig.1 (a) and (b)
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Fig. 2 of 3.
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Saturation Conc. (cm-3)
1-X (Ga fraction)
1-X (Ga fraction)
E (eV, reference to vacuum level)
InxGa1-xN Band Offset
Fig. 3 of 3
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