high resolution x-ray measurements of strain-mediated ... - ICDD

Copyright(c)JCPDS-International Centre for Diffraction Data 2000, Advances in X-ray Analysis,Vol43 201
HIGH RESOLUTION X-RAY MEASUREMENTS OF
STRAIN-MEDIATED DIFFUSION IN CuInSe 2
P. Fons, S. Niki, A. Yamada, M.Uchino, and H. Oyanagi
Electrotechnical Laboratory
1-1-4 Umezono, Tsukuba, Ibaraki JAPAN 305
email: fons@etl.go.jp
tel. 81-298-54-5636 fax. 81-298-54-5615
August 31, 2000
Abstract
CuInSe 2 is a direct bandgap material that assumes the chalcopyrite structure (space group I¯42d). Due
to its large optical absorption coefficient (α ∼ 10 5 cm −1 ),its use in thin film polycrystalline solar cells is
an active research topic. As most of the research effort to date has been applied directly to polycrystalline
film formation for use in solar cell applications and relatively little attention has been paid to fundamental
intrinsic material properties,there is a need to further investigate basic optical and structural properties
of single crystal material. Due to the fact that it is difficult to prepare stoichiometric material using
bulk growth techniques,we have employed molecular beam epitaxy to grow epitaxial layers of CuInSe 2
(a=0.57821 nm, c=1.1619 nm) on both GaAs (a=0.56537 nm) and relaxed,nearly lattice matched,
linear-graded buffer In xGa 1−xAs pseudosubstrates using molecular beam epitaxy. All films were grown
at 450 − 500 ◦ C using elemental sources of In (7N),Cu (7N),and Se (6N) and were prepared under Se
overpressure conditions. High resolution x-ray diffraction using a Ge (220) four crystal monochromator
as well as cross-sectional electron microscopy were used to study the as grown films. X-ray analysis found
that in the case of CuInSe 2 grown on a GaAs substrate interdiffusion of Ga into the relaxed CuInSe 2 layer
occurred,while in the case of the pseudosubstrate,no interdiffusion was observed. Photoluminescencebased
and positron annihilation analysis of point defects in the material will also be presented as well as
substrate annealing effects.
1 Introduction
One of the most promising materials for thin film semiconductor photovoltaics is CuInSe 2 which has a
direct bandgap of 1.04 eV at 1.4 K. The extraordinarily high near bandedge absorption coefficient of
CuInSe 2 (α ∼ 10 5 cm −1 ) has allowed fabrication of solar cells using only one to two microns of active layer
material. CuInSe 2 is a I-III-VI semiconductor that forms in the chalcopyrite structure with lattice constants
a=0.57821 nm and c=1.1619 nm yielding a c/a ratio of 2.0095 (JCPDS powder diffraction data).
While early work on CuInSe 2 began in the late 1970’s, the vast majority of research activities have been
directed towards the fabrication of polycrystalline solar cells. Indeed while single crystals of CuInSe 2 have
been grown, the complications of the phase diagram make it very difficult to control exact composition and
ensure the existence of a single phase of uniform composition. In an attempt to fill in some of the details of
the growth kinetics, structural features, and point defect (optical properties) structure, we have undertaken
the growth of CuInSe 2 using molecular beam epitaxy. The resulting thin films were then examined by highresolution
x-ray diffraction, hall/van der Pauw , photoluminescence, and positron annihilation measurements
among others. Details of the growth procedure are given in the next section. Interestingly, the highest
performance (polycrystalline) CuInSe 2 photovoltaics with efficiencies greater than 18%, have been fabricated
using evaporation.
The chalcopyrite structure belongs to the body-centered tetragonal system and its unit cell is composed of
two sphalerite subcells along the c direction with a reversal of Cu and In atoms occurring between the subcells.
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Copyright(c)JCPDS-International Centre for Diffraction Data 2000, Advances in X-ray Analysis,Vol43 202
A schematic picture of the chalcopyrite structure is shown in Fig. 1 on the left. The c axis is slightly larger
than twice that of the a axis. The chalcopyrite structure can be fully described by the lengths of the a and
c axes and a u parameter which specifies the remaining degree of freedom of the Se in the I¯42d space group.
The atom positions are as follows. The Cu atoms sit on sites with Wykoff 4a symmetry, namely at (0,0,0)
and (0,1/2,1/4) while the In atoms sit on 4b sites at the positions (1/2,1/2,0) and (1/2,0,1/4). The Se atoms
reside on 8d sites at the positions (u, 1/4, 1/8), (−u, 3/4, 1/8), (3/4, 1/2−u, 3/8), and (1/4, 1/2+u, 3/8). The
special case of u =1/4 corresponds to Se sitting equidistant from Cu and In. In the case of CIS, u =0.224,
therefore the Se atom sits closer to the Cu atom than the In atom. At high temperatures (T ≥ 800 ◦ )the
cations disorder and CuInSe 2 changes to a sphalerite structure δ (space group F¯43m, No. 216).
_
I42d
: 2a Cu
: 2b In
: 8d Se
Chalcopyrite
_
I42m
: 2a Cu (0.8)
: 2b In (0.4)
: 4d In
: 8i Se
Stannite
Figure 1: The chalcopyrite and stannite structures
Unlike ZnSe, the II-VI binary analogue of CuInSe 2 , a variety of compounds have been found to exist in
the CuInSe 2 phase diagram which is usually expressed as a pseudobinary with endpoints Cu 2 Se and In 2 Se 3 .
The significance of this pseudobinary tieline is that represents constant valence stoichiometry. On the Cu
rich side of the phase diagram, CuInSe 2 coexists with the Cu-rich phase Cu 2−x Se. As Cu 2−x Se is metallic in
nature, Cu-rich compositions are avoided in solar cell fabrication due to the formation of Cu 2−x Se shunt paths
along CuInSe 2 grain boundaries which reduce cell performance. Indeed, the high temperature (> 180 ◦ C)
phase of Cu 2−x Se is well known for being a superionic conductor. A schematic representation of the current
understanding of the phase diagram is down in Fig. 2 in which CuInSe 2 is represented as a pseudobinary
with endpoints Cu 2 Se and In 2 Se 3 . On the In-rich side of the phase diagram, there is no clear phase boundary
observed implying a variety of vacancy stabilized chalcopyrites may exist. The structure of CuIn 3 Se 5 has
recently been investigated by convergent beam TEM and Rietveld fitting and has been determined to have
the stannite structure (space group I¯42m as can be seen in Fig. 1. The unit cell of CuIn 3 Se 5 has lattice
constants c=1.1518 nm and a=0.5554 nm with the extinction rule h+k+l=2n. As the stannite space group
I¯42m is not a subgroup of the chalcopyrite point group I¯42d, this implies that vacancy stabilized In-rich
composition CuIn x Se y are not formed by partial order or local cation substitution of CuInSe 2 .
2 Growth
The film growths described in this manuscript can be divided into two parts. The first set of experiments was
dedicated to investigating the effects of heteroepitaxial strain on CuInSe 2 and used both LEC GaAs(001) and
and nearly lattice matched In 0.29 Ga 0.71 As (001) pseudo-substrates. For these early experiments, substrates
were mounted to the Mo sample holders using In-solder. After these early experiments were complete, due
to the sensitivity of CuInSe 2 films to low temperature annealing effects, it was decided to avoid the heating
of samples to melt the In. In lieu of In a Ga-In eutectic was used that is liquid at room temperature. Thus
for the second set of experiments presented here relating to annealing effects, a Ga-In eutectic was used to
mount the samples. Sample temperatures were calibrated using infrared pyrometry.
For the first set of experiments relating to heteroepitaxial strain, all CuInSe 2 samples were grown by
molecular beam epitaxy from solid sources of Cu (7N), In (7N) and Se (6N) at a growth temperature
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Copyright(c)JCPDS-International Centre for Diffraction Data 2000, Advances in X-ray Analysis,Vol43 203
Cu 2 Se-In 2 Se 3 Pseudobinary System
CuInSe 2 CuIn3 Se 5
CuIn 5 Se 8
1000
δ + L
Temperature (¡C)
900
800
700
α
δ
α
+
δ
δ+β
δ+β
β
β
+
γ
γ γ + In 2 Se 3
Cu 2 Se
50 60 70 80 90 In 2 Se 3
Figure 2: The CuInSe 2 pseudobinary phase diagram
of 500 ◦ C as explained above. In these experiments, both sulfur-passivated GaAs (001) substrates and As 2 -
capped, nearly lattice matched In 0.29 Ga 0.71 As (001) pseudo-substrates were affixed to the same Mo substrate
block using In solder. For the second set of experiments relating to annealing effects, sulfur-passivated GaAs
(001) wafers were mounted to the Mo substrate block using a Ga-In eutectic and were grown at 450 ◦ C. In
both experiments, all wafers were (001) oriented and had no intentional miscut.
In both sets of experiments, the growth of flat, uniquely oriented, epitaxial thin films of CuInSe 2 as
well as the presence of the chalcopyrite structure was confirmed in-situ by RHEED. For the first set of
experiments, all samples were grown under Cu-rich (as measured by electron microprobe measurements
(EPMA)) conditions whereas for the second set both In-rich and Cu-rich samples were grown. All samples
were grown under Se overpressure conditions. Further details of the sample preparation procedure and
growth process have been described elsewhere[3].
The pseudo-substrates consisted of a linearly-graded In x Ga 1−x As buffer layer with x varying from x=0
to 0.29 using a concentration gradient of 1% In/40 nm. On top of the linearly-graded buffer a ∼ 800 nm
thick layer of In x Ga 1−x As of constant x=0.29 composition was grown. The pseudo-substrates were grown in
a separate chamber and an As 2 cap was used to protect the surface during transport to the growth system.
Triple-axis X-ray diffraction (XRD) measurements of these substrates indicated that the final layer was
about 90% relaxed. Use of a linearly graded buffer to provide strain relaxation and a reduction in threading
dislocation densities below 10 6 cm −2 has been demonstrated for large mismatch systems such as SiGe/Si[4]
and InGaP/GaAs[5]. The room temperature in-plane lattice mismatch was δa/a ∼ 2.2% and δa/a ∼ 0.2% for
the CuInSe 2 /GaAs and CuInSe 2 /In 0.29 Ga 0.71 As layers, respectively. In addition to the room temperature
strain, an additional complication arises due to the differences in thermal expansion coefficients between
CuInSe 2 and the GaAs substrate. This gives rise to an additional tensile strain (of the relaxed CuInSe 2 layer
upon cooling) on the order of δd/d ∼ 10 −3 for the temperatures used here. It should be mentioned here
that a slight compressive strain is necessary to ensure unique orientation of the tetragonal unit cell (e.g. the
CuInSe 2 c axis aligned to the growth direction).
During the initial states of growth, a sharp change in surface reconstruction from a streaky (1 x 2)
reconstruction with respect to the [1¯10] and [110] directions to a (1 x 3) reconstruction occurred. The
presence of a slight waviness along the [110] direction of the (1 x 3) surface was noted as well. This change
occurred after approximately 30 nm of CuInSe 2 . was deposited. In part, to understand the origin of this
change in surface reconstruction, a series of CuInSe 2 films of various thicknesses were grown on GaAs and
In 0.29 Ga 0.71 As pseudo-substrates. CuInSe 2 films of thicknesses 4.2, 12.6, 40, 63, 126, and 270 nm were
grown. These films were then investigated by high-resolution x-ray diffraction (HRXRD) and transmission
electron microscopy (TEM). The surface topology of several films was examined by atomic force microscopy
(AFM) as well.
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Copyright(c)JCPDS-International Centre for Diffraction Data 2000, Advances in X-ray Analysis,Vol43 204
HRXRD measurements were performed using a MAC Science 18,000 watt rotating anode Cu target. A
Phillips MRD diffractometer with a 4-bounce, symmetric Ge(220) monochrometer was used to select Cu
Kα 1 radiation (λ =0.15405981nm). A 1.2 mm receiving slit was used in front of the detector. In later
experiments a parabolic x-ray mirror and a 15 degree asymmetric (220) Ge four crystal monochromator were
used in conjunction with a triple bounce Ge (220) analyzer resulting in an increase in intensity of about
a factor of twenty. Transmission electron microscopy measurements were performed on a 400keV JEOL
4000FX microscope. The cross sections were prepared in a standard way using a tripod polisher and Ar ion
milling.
Film composition was measured by EPMA using a bulk CuInSe 2 crystal as a standard. Composition
is reported here in terms of as-measured Cu/In ratio. It is important to note that EPMA measures a
weighted composition with a strong near-surface weighting. If the film being measured is not uniform in
composition with depth, this results in a composition value representing the electron beam depth distribution.
In particular, for Cu-rich films as will be explained later, a second phase of Cu 2−x Se in the near-surface
region. Cross-sectional electron microscopy of Cu-rich composition films in conjunction with EPMA voltage
dependent measurements has yielded further information about the composition and shape of these Cu 2−x Se
regions as will be explained below [6]. In contrast, all In-rich samples were uniform in composition with
depth and no other phases were observed. Composition values as determined by EPMA measurements are
presented here in terms of molecularity, m which is defined as the Cu/In ratio. Within the experimental
error, EPMA results indicate that the (Cu+In)/Se ratio is always unity.
3 Surface Morphology
The typical surface morphology of films with molecularities less than unity (m ∼ 0.82) and greater than
unity (m ∼ 1.3) are show in Fig. 3 a and b, respectively. The surface morphology of films with m1 exhibit long undulations aligned along the [1¯10]
directions. Spacing between these undulations was typically 1-2 microns. The m=1.32 film show in Fig. 3b
had a RMS roughness of 6.8 nm while the m=0.82 film had a RMS roughness of 18.0 nm.
These surface undulations are associated with elastic strain relaxation heteroepitaxial strain and have
been seen in other compressively strained heteroepitaxial systems [7]. It is speculated here that the unique
directional orientation of the undulations is related to the 3 x 1 surface reconstruction of CuInSe 2 . T hese
surface undulations may also act as stress centers for the selective nucleation of dislocations and twins as is
discussed later.
For the film with m>1 (Fig. 3b), there are also what appear to be irregularly shaped inclusions present
on the surface. Cross sectional transmission electron microscopy (XTEM) and voltage dependent EPMA
modeling indicate that these are actually wedge shaped inclusions of Cu 2−x Se with x=0.5 that taper to
several tens of nanometers into the bulk of the layer. Rough estimates indicate that the volume fraction of
Cu 2−x Se is on the order of an atomic percent. EPMA analysis typically assumes that sample composition is
uniform with depth and that characteristic x-ray absorption effects can be removed using a standard sample.
Due to the presence of the Cu 2−x Se surface phase in Cu-rich films, the composition values for Cu-rich samples
reported here do not reflect the true composition of the films and are quoted only to indicate the extent of
Cu 2−x Se surface phase present, whereas for the In-rich films EPMA values do represent true composition.
4 Relaxation Dynamics
Fig. 4 shows TwoTheta/Omega scans around the CuInSe 2 (008) reflection for a series of CuInSe 2 films of
differing thickness (m >1) grown on a GaAs substrate. A vertical line also indicates the bulk JCPDS powder
data file two theta value corresponding to the relaxed CuInSe 2 (008) position.
Calculations of equilibrium critical thickness yield a value on the order of 8 nm, but from the figure it
is clear that relaxation has not occurred for films as thick as 40 nm and is only partially complete for films
63 nm thick. Two possible strain relief mechanisms are introduced. As mentioned in Sec. 3, the presence of
undulations implies the elastic relaxation of the interfacial misfit by diffusion of the larger atoms to areas of
lower strain as has been observed in SiGe alloys.
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Copyright(c)JCPDS-International Centre for Diffraction Data 2000,Advances in X-ray Analysis,Vol.43 205
50nm
40
30
20
10
0
a. In-rich film CdIn4.8 b. Cu-rich film Cu/In-1.3
Figure 3: Atomic Force Microscopy Images of as-grown In-rich and Cu-rich CuInSes films.
For the case of CuInSes, there is an additional complication of the Cu2-xSe surface phase which is tensile
strained on the surface. We have grown epitaxial Cuz-,Se on GaAs(OO1) and performed high resolution
x-ray mapping measurements as well as XTEM observations. As the lattice constants of the Cuz-,Se phase
almost certainly change with composition, the Cu2-xSe epitaxial layer was grown under the same growth
conditions as the CuInSez layer with the exception that the In cell shutter was closed. EPMA measurements
indicate a composition of Cui.sSe for both the Cuz-,Se/GaAs epitaxial layer and for the Cus-$e surface
layer on CuInSez. The lattice constants of the film were nominally a= 0.576 and c=O.574 nm although
there is a monoclinic supercell present as well [S]. While cross sectional TEM observations confirmed the
presence of the monoclinic supercell for Cus-,Se/GaAs, similar observations of the Cuz-,Se on top of the
CuInSes epilayer did not exhibit any supercell spots in selective area diffraction. It is speculated here
that small amounts of In from the CuInSez epilayer may inhibit the phase transition to the low temperature
monoclinic phase, causing the crystal structure to remain in the higher temperature (T > 125% cubic phase.
This difference in crystal structure along with the slight variation in composition is sufficient to explain the
differences in observed lattice constants.
To further elucidate the details of the strain relief mechanism, an area map was taken using a 1.2 mm
receiving slit in the vicinity of the CuInSez asymmetric (1 . 1 . 10) reflection for a 300 nm CuInSez film grown
on a GaAs substrate. The area map can be seen in Fig. 5. The presence of Cus+$Se phase can be clearly
seen in the figure as well as the fact that the Cuz-$Se phase is strained to the CuInSez in-plane lattice
constant. The lattice constants of the Cuz-$Se phase as seen in the figure are a=0.5800 and c=5.704 nm,
consistent with the lattice constants measured for the Cus-$Se film grown epitaxially on a GaAs substrate.
It should be noted here that the Cu2-$e second phase was visible in thinner films of the series as well and
always occurred at approximately the same location.
In addition to elastic strain relief related to the surface undulations, there is intermixing of Ga from
the substrate into the CuInSe2 layer giving rise to elongation of the CuInSez (1 + 1 . 10) reflection. This
is a consequence of interdiffusion of Ga from the substrate into the CuInSes film while Cu diffuses into
the substrate. The interdiffusion has been confirmed by TEM-based EDX observations as well as SIMS
measurements of the sample. Dynamical simulation of the interdiffusion were attempted but the large
mosaic of the sample made simulations difficult. Lattice constant changes lead to an estimate of a maximum
Ga concentration on the order of one percent. From SIMS results the extent of the interdiffusion is estimated
to be several hundred angstroms as can be seen in Fig. 6.
Photoluminescence data on the CuInSes/GaAs films indicated only weak free exciton emission at the
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Copyright(c)JCPDS-International Centre for Diffraction Data 2000, Advances in X-ray Analysis,Vol43 206
Log Intensity
5
4
3
2
1
Strained
CuInSe 2
CuInSe 2
(008)
42 126
400 630
2700
1260
Relaxed
CuInSe 2
GaAs
(004)
0
60 61 62 63 64 65 66
Two Theta/Omega
Figure 4: A series of 008 TwoTheta/Omega Scans for CuInSe 2 films of differing thickness
bandedge and in which a conduction band to donor transition was dominant. To eliminate the possibility
that threading dislocations were the cause of the weak near-bandedge photoluminescence emissions, CuInSe 2
was grown on a In 0.29 Ga 0.71 As linear grade buffer with a room temperature lattice mismatch of 0.2%.
Photoluminescence on these films showed sharp bandedge emission in which the free exciton and its excited
states were dominant and the conduction band to acceptor transition was dramatically reduced[9].
Triple axis measurements of the a and c lattice constants of CuInSe 2 grown on GaAs and the pseudosubstrate
were performed. The a differences in the a lattice constants along the two orthogonal < 110 >
directions were ∆a/a avg ≤ 0.1% for CuInSe 2 on GaAs while a difference of ∆a/a avg ≤ 0.2% was observed for
the case of a CuInSe 2 grown on the pseudosubstrate. The average of the lattice constants along the two orthogonal
directions were a=0.57920 and c=1.16149 nm for CuInSe 2 /GaAs and a=5.7835 and c=1.16175 nm
for the case of CuInSe 2 grown on the pseudosubstrate. This indicates that the CuInSe 2 film grown on the
pseudosubstrate is nearly perfectly relaxed while that grown on GaAs direction is under slight tensile strain.
This is attributed to the factor of two difference in the thermal expansion coefficient of CuInSe 2 over GaAs
in conjunction with the poorer activation of slip during the cooling of the substrate from growth to room
temperature for the CuInSe 2 grown on GaAs as opposed to the layer grown on a pseudosubstrate. This is
in turn a consequence of the higher density of misfit dislocations in the CuInSe 2 grown directly on GaAs.
The complex strain interactions among these dislocations limit slip in the CuInSe 2 /GaAs case relative to
the lower dislocation density of the CuInSe 2 film grown on the pseudosubstrate.
The FWHM of the CIS (008) rocking curves were nearly identical with values 0.117 and 0.106 degrees
for CuInSe 2 on GaAs and the pseudosubstrate, respectively. This is not surprising in light of the reduction
in threading dislocation density as confirmed by cross sectional TEM. This is consistent with recent microdiffraction
measurements of SiGe alloys grown on Si and linearly graded buffers in which it was found that
domains of dislocation free blocks present in the SiGe film grown on the graded buffer were significantly
larger than those of the SiGe film grown directly on the Si substrate regardless of the fact that both films
had similar mosaic structures as measured by conventional x-ray diffraction [10].
Fig. 7 shows two-dimensional x-ray reciprocal maps of the asymmetric chalcopyrite (309) reflection for
CuInSe 2 grown on GaAs and a pseudosubstrate. The maps were taken using a 1.2 mm receiving slit. As the
chalcopyrite peak arises from the difference in scattering between the two cations, it is absent in the GaAs
layer allowing a more detailed examination of interdiffusion in the pseudosubstrate grown film. As can be
seen in the figure, there is a clear asymmetry in the (309) reflection from the CuInSe 2 grown on the GaAs
substrate, whereas in for the layer grown on the pseudosubstrate no asymmetry can be seen.
The presence of interdiffusion for the more highly strained case may be related to differences in threading
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Copyright(c)JCPDS-International Centre for Diffraction Data 2000, Advances in X-ray Analysis,Vol43 207
GaAs
5.55
Q[001] ( -1 )
5.5
5.45
5.4
Cu 2-x
Se
CuInSe 2
z
> 50.
< 50.
< 45.
< 40.
< 35.
< 30.
< 25.
< 20.
< 15.
< 10.
< 5.
5.35
-1.65 -1.6 -1.55 -1.5 -1.45
Q[110] ( -1 )
Figure 5: X-Ray reciprocal space map near the (1 · 1 · 10) CuInSe 2 reflection
Number
2.0
1.5
1.0
speculated
interface
location
Se
In
Cu
Ga
As
0.5
0.0
0
200
400 600 800
Depth (channels)
1000
1200
Figure 6: Normalized SIMS results for a 800 nm CuInSe 2 film grown on GaAs at 500 ◦ C
Q[001] (Å -1 )
5.1
5.05
5
4.95
4.9
4.85
4.8
4.75
Intensity
> 64.
< 64.
< 32.
< 16.
< 8.
< 4.
< 2.
< 1.
A)
Ga
Interdiffusion
5
4.95
4.9
4.85
4.8
4.75
4.7
Intensity
> 9.75
< 9.75
< 9.08
< 8.41
< 7.74
< 7.08
< 6.41
< 5.74
< 5.08
< 4.41
< 3.74
< 3.08
< 2.41
< 1.74
< 1.07
< 0.41
3.1 3.15 3.2 3.25 3.3 3.35 3.4 3.45
3.1 3.2 3.3 3.4
Q[100] (Å -1 ) Q[100] (Å -1 )
B)
Figure 7: X-ray reciprocal space maps of the chalcopyrite (309) reflection for a) CuInSe 2 /GaAs and
b)CuInSe 2 /In 0.29 Ga 0.71 As pseudosubstrates
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dislocation densities between the two systems, but more work must be done to elucidate the mechanism for
the interdiffusion.
5 Annealing Experiments
Typical (hole) resistivity for Cu-rich CuInSe 2 is typically on the order of 2x10 −2 Ω − cm due to the presence
of the conductive Cu 2−x Se phase. Upon annealing of a Cu-rich, 800 nm thick sample (m=1.32) for 30
minutes at 540 ◦ C in an Ar ambient, it was found that the resistivity of the sample had increased over three
orders of magnitude to 300 Ω − cm. AFM observations also indicated that the inclusions on the surface of
the CuInSe 2 had vanished leaving holes in their place. This resulted in an increase in RMS surface roughness
from 6.85 to 8.34 nm for the as-grown and 540 ◦ C annealed samples, respectively. The (hall) mobility also
showed a dramatic increase from a value in the low teens to over 100cm 2 /V − s, a value significantly higher
than bulk stoichiometric CuInSe 2 . This implied that the Cu-rich surface layer had diffused into the GaAs
substrate. As one of the levels of Cu in GaAs is a low activation energy acceptor, a likely explanation is that
the Cu-component diffused through the CuInSe 2 matrix into the substrate.
Q[001] ( -1 ) Q[001] ( -1 )
3.4
3.35
3.3
3.25
3.2
3.4
3.35
3.3
3.25
3.2
3.4
A) 3.35 B)
3.3
3.25
3.2
1.5 1.525 1.55 1.575 1.6 1.625 1.5 1.525 1.55 1.575 1.6 1.625
3.375
C) 3.35 D)
3.325
3.3
3.275
3.25
3.225
3.2
1.5 1.525 1.55 1.575 1.6 1.625
Q[110] ( -1 )
1.5 1.52 1.54 1.56 1.58 1.6 1.62
Q[110] ( -1 )
Figure 8: (113) (low incidence) x-ray reciprocal space maps for a) In-rich as-grown, b) In-rich 540C Ar
annealed, c) Cu-rich as-grown, and d) Cu-rich 540C Ar annealed CuInSe 2 samples
From photoluminescence as well as positron annihilation experiments, it was was found that the dominant
residual point defect in Cu-rich CuInSe 2 is Cu-vacancies V Cu [11]. This implies that there is a large
concentration of V Cu available to assist mass diffusion. It should be noted here that the Cu 2−x Se surface
phase is a well know superionic conductor whose conduction mechanism lies principally with the very high
mobility of Cu atoms which move through a cubic symmetry shell provided by the Se atoms. It is surmised
here that the Cu atoms migrate through the V Cu rich CuInSe 2 layer into the substrate to form a p-doped,
but still high resistive layer on the GaAs side of the interface. As Se evaporates readily in the 540 ◦ Cambient
used, the effect of annealing on Cu-rich CuInSe 2 films is to induce a dramatic increase in resistivity with the
vanishing of the metallic Cu 2−x Se surface phase. At the same time Hall measurements indicate a dramatic
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rise in mobility. The change in mobility is a reflection of the doping of the higher mobility undoped GaAs
substrate. It is well known that the diffusion lengths at the annealing temperature of 540 ◦ Caresuchthat
Cu will diffuse throughout the substrate.
Glancing incidence asymmetric diffraction experiments were performed on the Cu-rich (m=1.32) sample
described above as well as an In-rich (m=0.82) sample using a triple axis configuration and are shown in
Fig. 8. This glancing incidence geometry has the advantage that the penetration depth of the sample due to
absorption and extinction effects is on the order of a micron, resulting in an increase in surface and interface
sensitivity. It is clear from the figure that while there is no apparent interdiffusion for the as grown samples
(A In-rich and B Cu-rich), after annealing for 30 minutes at 540 ◦ C in an Ar ambient, there is significant
interdiffusion only for sample D, the Cu-rich sample. In contrast, there is no apparent change for the In rich
sample post annealing (Fig.8D). Annealing was also performed at the intermediate temperature (not shown)
of 410 ◦ C (30 minutes, Ar ambient) for both In-rich and Cu-rich samples. No change in the {113} reciprocal
maps nor in sample conductivity or mobility was observed.
Figure 9: High resolution TEM cross sections of the CuInSe 2 /GaAs interface for a) Cu-rich and b) In-rich
films
Lateral correlation lengths were calculated from asymmetric {113} reflections by trigonometrically deconvolving
the effects of mosaic tilt as has been described by Fewster [12]. For In-rich and Cu-rich films,
the correlation lengths were on the order of 80 and 60 nm, respectively. After annealing, the lateral correlation
was observed to decrease slightly to 70 and 50 nm for the Cu-rich and In-rich films, respectively.
While the decrease in correlation length for the Cu-rich film may be related to interdiffusion, the decrease
in correlation for the In-rich film is more likely associated with the emergence of the stannite phase. This is
further corroborated by the fact that [001] lattice normal scans showed a factor of two increase in {00 · 2n}
reflections before and after annealing.
A possible explanations for the difference in annealing behavior lies in the point defect composition of
the Cu-rich phase and its resulting effects on interdiffusion. Typical phase image XTEM cross sections of
both Cu-rich and In-rich CuInSe 2 are shown in Fig.9. It is clear from the figure that there is significantly
more contrast at the GaAs/CuInSe 2 interface for the Cu-rich film. The areas of lighter contrast evident in
Fig.9a, did not change with tilt implying that they were probably voids. The In-rich interface seen in Fig.9b,
in contrast, has no such void-like regions and the interface is relatively abrupt.
6 Twinning
Concomitant with a change from Cu-rich to In-rich morphology is a dramatic increase in twin concentration.
The existence of these growth twins is undesirable from a device point of view and thus it is of interest
to study the conditions under which they form. To examine the dependence on twinning more closely, a
series of four CuInSe 2 layers of differing Cu/In ratio were grown on GaAs(001) substrates and were then
examined using pole figure analysis. The chalcopyrite (103) poles was chosen for study as they eliminate the
possibility of interference from the substrate and also provide information about unit cell orientation (e.g.
the tetragonal c axis). In the kinematic approximation the chalcopyrite reflections arise from the difference
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Copyright(c)JCPDS-International Centre for Diffraction Data 2000, Advances in X-ray Analysis,Vol43 210
in scattering from the two different cations or (f Cu − f In ) 2 . As in the sphalerite case there is only one type
of cation this reflection is not allowed.
Growth twins in CuInSe 2 are 180 ◦ rotation twins about {112} poles. In total, (in the upper half plane)
there are 24 different crystallographic directions into which the original four {103} poles can transform. By
the action of a single rotational twinning operation the original four {103} reflections are transformed from
a position which makes approximately a 33 ◦ angle with the (001) surface normal to eight different positions
making approximately a 50 ◦ angle with respect to the surface normal.
A
_
(1 1 2)
B
_
(1 1 2)
(1 0 3) T
(0 1 3)
(0 1 3) T
(1 0 3)
_
(0 1 3)
_
(1 0 3)
_
(0 1 3) T
_
(1 0 3)
_ _
(1 1 2)
(1 1 2)
(1 0 3) _ _
T (1 0 3) (0 1 3) (0 1 3) T
_
(1 1 2)
(1 __
1 2)
_
_
_
(1 1 2) T
(1 1 2) T
_
(0 1 3) (1 0 3) _
(0 13) T (1 0 3) T
_
(1 1 2)
_
(1 1 2)
Figure 10: {103} pole figures for a) Cu-rich (m=1.19) and b) In-rich (m=0.85) CuInSe 2 grown on GaAs
Pole figures for two of the samples (Cu-rich m=1.19 and In-rich m=0.85) are shown in Fig.10. It is
clear that there is a sharp drop in twin concentration in going from In-rich to Cu-rich compositions. Taking
the ratio of the integrated intensity of the original {103} poles to that of those in twinning positions (after
correcting for geometrical effects) gives a rough estimate of the twinned volume. For the Cu-rich film shown
in Fig.10a, the volume of twins is ≤ 1 ± 0.3 volume percent. In addition, the twins were to exclusively form
around the (112) poles. On the other hand, for the In-rich films, an example of which is shown in Fig.10b,
twins form about all four {112} poles with a twinned volume of ∼ 3 ± 0.3 volume percent. This anisotropy
of twinning is observed for all samples.
High resolution TEM confirms that the twinning occurs predominantly in the form of microtwinning with
characteristic lengths of about 20 nm. TEM observations also showed that for Cu-rich material, twins are
mostly confined to near the interface with a decreasing concentration towards the surface, presumably due to
recombination of bounding partial dislocations on alternative {112} glide planes. For In-rich films, however,
twins were found to extend throughout the sample volume. As weak stannite reflections (h+k+l=2n) are
observed for In-rich films, the low stacking fault energy of In-rich films may be a consequence of the coexistence
of the vacancy stabilized stannite structure along with the chalcopyrite structure. Indeed, local density
functional calculations have predicted in GaAs that dislocation energies can be influenced by impurities[13].
It should also be noted that for the {112} planes (as opposed to the {1¯12} planes) are terminated with a
cation surface, further suggesting the link between point defects (divacancies) and the anisotropy in twinning.
Recent experiments have also indicated that the introduction of oxygen during growth of In-rich films
dramatically reduces the number of growth twins about all {1¯12} poles.
Photoluminescence and positron annihilation experiments indicate that for In-rich films the predominant
(optically active) point defect is a divacancy which consistent of a copper vacancy and a selenium vacancy on
adjacent sites. By the introduction of a small amount of oxygen during growth, it is possible to significant
reduce the number of Se vacancies. With the introduction of oxygen, the dominant defect observed by
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Copyright(c)JCPDS-International Centre for Diffraction Data 2000, Advances in X-ray Analysis,Vol43 211
photoluminescence is the conduction band copper vacancy transition, the same as for the case of Cu-rich
material. It is likely that this change in point defect type is direction related to the reduction in twin
concentration upon the introduction of oxygen during growth. Currently, further studies are under way to
clarify this interrelationship.
7 Conclusion
Several crystallographic features of Cu-rich and In-rich CuInSe 2 have been summarized. There is apparently
extremely fast diffusion in Cu-rich CuInSe 2 , but comparatively little interdiffusion in In-rich CuInSe 2 .This
fact is reflected in that optimal (large) grain growth in polycrystalline CuInSe 2 films used for photovoltaics
are typically grown in at least two stages, the earlier of which is inevitably Cu-rich while the final composition
is adjusted to become slightly In-rich so as to eliminate the possibility of Cu 2−x Se induced shunt paths which
act to reduce the open circuit voltage of photovoltaics made from the layers. This larger grain growth is a
direct consequence of the higher diffusion lengths in Cu-rich CuInSe 2 over In-rich CuInSe 2 .
That points defects in CuInSe 2 as a function of stoichiometry play an important role in determining
optical and electrical properties is well established, but still only partially understood due to the complex
dynamics of the system. The apparent importance of point defects and hence composition in diffusion is yet
another reflection of this.
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