materials
Article
Crystallisation Phenomena of In2O3:H Films
Ruslan Muydinov 1,*, Alexander Steigert 2, Markus Wollgarten 3, Paweł Piotr Michałowski 4,
Ulrike Bloeck 3, Andreas Pflug 5, Darja Erfurt 6, Reiner Klenk 6, Stefan Körner 1,
Iver Lauermann 6and Bernd Szyszka 1
1Institute of High-Frequency and Semiconductor System Technologies, Technical University Berlin,
2Institute for Nanospectroscopy, Helmholtz-Zentrum Berlin für Materialien und Energie GmbH,
3Department of Nanoscale Structures and Microscopic Analysis, Helmholtz-Zentrum Berlin für Materialien
[email protected] (U.B.)
4Institute of Electronic Materials Technology, Wolczynska Str. 133, 01919 Warsaw, Poland;
5Fraunhofer Institute for Surface Engineering and Thin Films IST, Bienroder Weg 54e, 38108 Braunschweig,
Germany; [email protected].de
6
PVcomB, Helmholtz-Zentrum Berlin für Materialien und Energie GmbH, Schwarzschildstr. 3, 12489 Berlin,
iver[email protected] (I.L.)
*Correspondence: r[email protected]
Received: 21 November 2018; Accepted: 4 January 2019; Published: 15 January 2019
Abstract:
The crystallisation of sputter-deposited, amorphous In
2
O
3
:H films was investigated.
The influence of deposition and crystallisation parameters onto crystallinity and electron hall mobility
was explored. Significant precipitation of metallic indium was discovered in the crystallised films by
electron energy loss spectroscopy. Melting of metallic indium at ~160
◦
C was suggested to promote
primary crystallisation of the amorphous In
2
O
3
:H films. The presence of hydroxyl was ascribed
to be responsible for the recrystallization and grain growth accompanying the inter-grain In-O-In
bounding. Metallic indium was suggested to provide an excess of free electrons in as-deposited In
2
O
3
and In
2
O
3
:H films. According to the ultraviolet photoelectron spectroscopy, the work function of
In
2
O
3
:H increased during crystallisation from 4 eV to 4.4 eV, which corresponds to the oxidation
process. Furthermore, transparency simultaneously increased in the infraredspectral region. Water
was queried to oxidise metallic indium in UHV at higher temperature as compared to oxygen in
ambient air. Secondary ion mass-spectroscopy results revealed that the former process takes place
mostly within the top ~50 nm. The optical band gap of In
2
O
3
:H increased by about 0.2 eV during
annealing, indicating a doping effect. This was considered as a likely intra-grain phenomenon
caused by both (In
0
)
O••
and (OH
−
)
O•
point defects. The inconsistencies in understanding of In
2
O
3
:H
crystallisation, which existed in the literature so far, were considered and explained by the multiplicity
and disequilibrium of the processes running simultaneously.
Keywords: In2O3:H; thin films; crystallisation; TCO; high mobility
1. Introduction
Hydrogen doped indium oxide (In
2
O
3
:H) films demonstrate an electron Hall mobility (
µe
) of
over 100 cm
2
/Vs [
1
,
2
]. Such outstanding property can be achieved if the In
2
O
3
film is deposited in
an amorphous state and subsequently crystallised at T > 160
◦
C. Its optical transmittance is superior
Materials 2019,12, 266; doi:10.3390/ma12020266 www.mdpi.com/journal/materials
Materials 2019,12, 266 2 of 20
in both ultraviolet (UV) and near-infrared (NIR) spectral ranges when compared to the widely used
transparent conductive oxides (TCOs), such as ZnO:Al or In
2
O
3
:Sn. This results in a performance gain
for Si-based [3–5] and CIGS (Cu-In-Ga-S/Se)-based [6–8] solar cells.
Amorphous In
2
O
3
films are usually obtained by sputter-deposition in the presence of water
vapour without intentional heating [
1
,
2
]. Low crystallisation temperatures make them applicable on
top of a-Si and CIGS-absorbers. Moreover, the crystallisation seems to proceed within a few minutes [
3
].
Thus In2O3:H can be essential for tandem solar cell concepts, e.g., with hybrid perovskites.
There is a number of thorough theoretical and experimental investigations on In
2
O
3
,
its defectiveness and interaction with a gas phase. Nevertheless, some ambiguity with respect
to the origin of the n-type conductivity, the role of the crystallisation and the origin of the high
carrier mobility In
2
O
3
:H films still exists. For instance, there is no satisfactory explanation for the
following phenomenon so far: the crystallisation of In
2
O
3
:H films is accompanied with a Burstein-Moss
shift and a decrease of the charge carrier concentration (N
e
) simultaneously. At the same time, it is
widely accepted in literature that the n-type doping of crystalline In
2
O
3
is related to the hydrogen
incorporation. Another important disagreement concerns the “high mobility”. Exactly the passivation
of dangling bonds on grain-boundaries by hydrogen is supposed to be the most probable reason for
high mobility [
9
]; however, only this intergranular hydrogen disappears primarily during heating in
vacuum [10].
Our recent studies revealed the effect of oxygen addition to the sputter-gas on the In
2
O
3
:H
crystallisation progress [
11
]. However, the question of whether we deal with so-called plasma damage
still remains. This negative effect of the accelerated O
−
ions largely determines electrical properties
and their spatial deviations over the film [12,13].
The inconsistencies noted above represent a lack of understanding of the In
2
O
3
:H material. In this
work we investigated microstructural, compositional, optical and electrical properties of the In
2
O
3
and In
2
O
3
:H films. We also observed how these properties change during In
2
O
3
crystallisation.
Additionally, the discussion section contains a brief review of the relevant findings accumulated so far
in literature. Finally, we attempted to compose a full puzzle from all these data and understand this
material better.
2. Experimental Section
The as-deposited films will be further assigned as In
2
O
3
:H
2
O, because they were found to contain
hydroxyl groups. The post-deposition annealed films we denote by In
2
O
3
:H, as widely accepted.
In
2
O
3
:H
2
O films were obtained in a stationary mode by RF (Radio Frequency, 13.56 MHz) magnetron
sputtering from a round Ø 3
00
ceramic In
2
O
3
(99.99%) target onto unheated, stationary 1
00 ×
1
00
Eagle
XG (Corning, 0.7 mm thick) glass substrates. The substrates were cleaned in a multistage glass-washer
using surfactants and de-ionised water. The target-to-substrate distance was fixed at the optimum
40 mm. The RF power was set at 60 W, which results in 1.3 W/cm
2
power density. The base pressure
of the sputtering system was 5
×
10
−6
Pa. Water vapour was admitted through a needle valve
from a reservoir. Adjusting the valve manually, the p(H
2
O) was stabilised at 2
×
10
−3
Pa prior to the
deposition. The sputtering gas was pure Ar, and the following process pressures were compared: 0.5 Pa
(Ar flow: 20 sccm) and 1.3 Pa (Ar flow: 100 sccm). The films were annealed after deposition under
different p(O
2
) conditions: ~21 kPa (ambient air) or ~10
−6
Pa (ultra-high vacuum). Very thin (~20 nm)
films obtained on 0.3 mm thick Si-wafers were analysed by XPS/UPS without breaking vacuum in the
combined deposition and surface analysis system (CISSY) [
14
]. It includes a standard XPS laboratory
module with a non-monochromatic X-ray source and Mg (Al) anodes providing an excitation energy of
1253.6 eV (1486.6 eV). UPS measurements were done with a standard He lamp, yielding 21.22 eV (He I)
excitation energy. As photoelectron analyser, a VG CLAM 4 with a hemispherical energy filter and an
electron detector based on discrete channeltrons was operated at 20 eV pass energy. A sputtered Au
foil served as a reference for energy calibration based on the Au4f7/2 transition at 84.0 eV.
Materials 2019,12, 266 3 of 20
Si-substrates were also used for the secondary ion mass-spectroscopy (SIMS) investigation so
to provide better analytical response. The RF plasma effectively heats thin substrates, resulting in
uncontrolled crystallisation during deposition (see Figure S1). Therefore we have used DC sputtering to
secure the amorphous state of In
2
O
3
:H
2
O film, which should be thick enough for the meaningful depth
resolved SIMS. Our long-time experience revealed no principle difference in optical and electrical
properties of the RF- and DC-sputtered In
2
O
3
:H films crystallised from a fully amorphous state.
A mid-frequency pulsed (65 kHz,
τ
= 3.2
µ
s) DC magnetron sputtering was conducted in the deposition
system A600V7 (Leybold Optics Dresden GmbH) at 75 mm target-to-substrate distance. The power
density was 2.67 W/cm
2
, total pressure—0.4 Pa, partial pressures of oxygen (O
2
) and water (H
2
O)
during processing were 2.3
×
10
−4
Pa and 1.7
×
10
−4
Pa, respectively. The deposition took place on an
oscillating substrate. In2O3films were deposited under the same conditions but without introducing
water vapour. The annealing took place in UHV at 220 ◦C for 30 min.
All SIMS depth profiles in this work were performed employing the CAMECA SC Ultra
instrument operating under ultra-high vacuum (UHV) of ~4
×
10
−8
Pa. A Cs
+
primary beam with an
impact energy of 3 keV and intensity of 10 nA was scanned over the (150
×
150)
µ
m
2
area, whereas the
analysis was limited to the (50
×
50)
µ
m
2
area. The positive ion detection mode was used and thus
each element was measured as the CsX
+
cluster. Subsequently, a point-to-point normalisation to the
Cs
+
signal was performed. Thus, due to a significant reduction of the matrix effect [
15
–
19
], we were
able to detect adequately the In/O and H/O ratios as well as the annealing losses. The latter were
determined as follows:
LH=1−IHannealed
IHas−deposited
,LO=1−IOannealed
IOas−deposited
(1)
where I—are the intensities normalised to the Cs signal and L
H
and L
O
are the relative annealing losses
of hydrogen and oxygen, respectively. Despite the reduced detection limit for hydrogen in the positive
ion detection mode, its concentration in water containing samples was found to be high enough for a
reliable determination.
Fourier-transform infrared spectroscopy (FTIR) was performed on a Vertex 70 (Bruker Optics,
Ettlingen, Germany) using the rock solid interferometer, mercury (Hg) IR-source, DLaTGS IR-Detector
and the KBr Beam-splitter. In
2
O
3
:H films were deposited on double-side polished Si-substrates
for this investigation. XRD patterns were recorded using Cu K
α
radiation in different scanning
modes: symmetrical Bragg-Brentano and asymmetrical detector scan. In the in-plane measurement,
both incident and diffracted beams had grazing angles to the sample surface. D8 Discover (Bruker,
Karlsruhe, Germany) and X’Pert MRD Pro (PANalytical, Almelo, The Netherlands) diffractometers
were used for these tasks. LaB
6
powder (660c) was used as a standard for estimation of the crystallite
size. The Hall mobility of charge carriers was measured with an Ecopia HMS-3000 system (Anyang-city,
Gyeonggi-do, Korea) in van der Pauw geometry at room temperature. The proportionality factor was
taken equal to 1. Scanning electron microscopy (SEM) was made using a Hitachi S4100 microscope
(Hitachi High-Technologies Corporation, Tokyo, Japan).
Samples for cross sectional transmission electron microscopy (TEM) were prepared by gluing
the thin films face-to-face, followed by mechanical and ion thinning for electron transparency.
TEM investigations were performed on two microscopes. The Philips CM12 (FEI Company, Hillsboro,
OR, USA) was used at 120 kV accelerating voltage for preliminary investigation. The system Zeiss
LIBRA 200 FE (Carl Zeiss Microscopy, Jena, Germany) operated at 200 kV accelerating voltage was used
for more detailed analysis by electron energy loss spectroscopy (EELS). This microscope is equipped
with an in-column energy filter for energy filtered image acquisition. The set electron energy loss was
varied from 0 to 30 eV.
Materials 2019,12, 266 4 of 20
3. Results
3.1. Crystallinity versus Electron Mobility
The crystallinity of TCO films, i.e., the size of grains and the texture, can determine electron
mobility to a different extent, which depends on the individual material. According to the band
structure calculations accepted for In
2
O
3
, conduction band minima are formed by the 5 s states [
20
–
22
].
Because of their spherical symmetry, one may then expect no significant impact of the coherence
between grains (texture) onto the grain boundary scattering. However, a limiting role of the grain
boundaries themselves (grain size) illustrates the following fact: the undoped single crystalline In
2
O
3
films and In
2
O
3
single crystals demonstrate an electron mobility of more than 100 cm
2
/Vs
[23,24]
whereas the as-prepared polycrystalline films do not reveal
µe
exceeding 50–60 cm
2
/Vs [
25
,
26
].
This fact is basically related to the difference in N
e
, which is known to be much lower in single crystals
(~10
17
cm
−3
) as compared to the polycrystalline materials (10
19
cm
−3
) [
23
]. In turn, high N
e
values
in polycrystalline In
2
O
3
are usually attributed to the so called unintentional doping caused by the
presence of water in ambient air. Additional negative impact on mobility provide ionized impurities.
In this work, we intended to examine to what extent the grain size and texture determine
Hall mobilities in In
2
O
3
:H films. Figure 1presents the X-ray patterns and Hall mobility data for
various as-deposited and annealed films. Here, two series of experiments were done: variation of
the film-thickness (a) and variation of the total pressure during deposition (b). As one may see from
the series (a), all
≥
200 nm thick In
2
O
3
:H
2
O films appear to be X-ray crystalline. Diffraction maxima
are, however, quite broad. As we explained above, this is due to the heating of the surface by the RF
plasma. This effect increases with time due to the NIR absorption of the growing film and at some
point the film becomes crystalline (see Figure S1).
Materials 2019, 12, x FOR PEER REVIEW 4 of 19
boundaries themselves (grain size) illustrates the following fact: the undoped single crystalline In2O3
films and In2O3 single crystals demonstrate an electron mobility of more than 100 cm2/Vs [23,24]
whereas the as-prepared polycrystalline films do not reveal μe exceeding 50–60 cm2/Vs [25,26]. This
fact is basically related to the difference in Ne, which is known to be much lower in single crystals
(~1017 cm−3) as compared to the polycrystalline materials (1019 cm−3) [23]. In turn, high Ne values in
polycrystalline In2O3 are usually attributed to the so called unintentional doping caused by the
presence of water in ambient air. Additional negative impact on mobility provide ionized impurities.
In this work, we intended to examine to what extent the grain size and texture determine Hall
mobilities in In2O3:H films. Figure 1 presents the X-ray patterns and Hall mobility data for various
as-deposited and annealed films. Here, two series of experiments were done: variation of the film-
thickness (a) and variation of the total pressure during deposition (b). As one may see from the series
(a), all ≥200 nm thick In2O3:H2O films appear to be X-ray crystalline. Diffraction maxima are, however,
quite broad. As we explained above, this is due to the heating of the surface by the RF plasma. This
effect increases with time due to the NIR absorption of the growing film and at some point the film
becomes crystalline (see Figure S1).
Post-deposition annealing at 180 °C for 60 min in ambient air results in further crystallisation
and increase of electron mobility. The biggest gain of mobility is demonstrated by thinner, initially
X-ray amorphous, films. All partially crystalline In2O3:H2O films thicker than 300 nm reveal almost
the same gain of mobility after annealing. Analysing XRD patterns in Figure 1a, one can find two
In2O3 phases: initially crystalline and crystallised after annealing. A coexistence of two phases is
especially well visible in the 400 nm thick film. Correlating this with Hall mobility data, we conclude
that exactly the post-deposition crystallisation is crucial for high mobility. Interestingly, the “high-
mobility” phase reveals a smaller lattice constant than the phase, which crystallises during
deposition. The origin of this difference will be discussed below.
A decrease of intensity of diffraction peaks is observed for the initially crystalline In2O3 phase as
a result of annealing. The most probable reason for that is active recrystallization.
Figure 1b demonstrates the influence of the sputtering pressure on the crystallisation behaviour
of X-ray amorphous films. Apparently, this parameter determines the film texture and Hall mobility.
We attribute this effect to the amount of hydroxyl in a film. Indeed, using a fixed leak rate of the
needle valve feeding water vapour, a higher ptot should result in a smaller p(H2O) and, hence, in a
smaller concentration of hydroxyl in In2O3. In turn, the impact of hydroxyl on the crystalline growth
of In2O3 is a known phenomenon, thus, the (400)-orientation can be suppressed if water is present in
a sputtering gas [27]. This is consistent with our results. A higher process pressure itself may also
contribute, as it increases the energy transfer from plasma to substrate, resulting in a higher
deposition temperature [28].
(a) (b)
Figure 1. XRD patterns and Hall-mobility data for the RF sputtered In2O3:H films: (a) variation of
thickness at ptot = 0.5 Pa; (b) variation of the sputtering pressure at a fixed film thickness. In all cases
the as-deposited state (In2O3:H2O) is compared with the annealed one (In2O3:H). The cubic In2O3
powder standard (ICDD Nr. 00-006-0416) is shown in grey bars.
Figure 1.
XRD patterns and Hall-mobility data for the RF sputtered In
2
O
3
:H films: (
a
) variation of
thickness at p
tot
= 0.5 Pa; (
b
) variation of the sputtering pressure at a fixed film thickness. In all cases
the as-deposited state (In
2
O
3
:H
2
O) is compared with the annealed one (In
2
O
3
:H). The cubic In
2
O
3
powder standard (ICDD Nr. 00-006-0416) is shown in grey bars.
Post-deposition annealing at 180
◦
C for 60 min in ambient air results in further crystallisation and
increase of electron mobility. The biggest gain of mobility is demonstrated by thinner, initially X-ray
amorphous, films. All partially crystalline In
2
O
3
:H
2
O films thicker than 300 nm reveal almost the
same gain of mobility after annealing. Analysing XRD patterns in Figure 1a, one can find two In
2
O
3
phases: initially crystalline and crystallised after annealing. A coexistence of two phases is especially
well visible in the 400 nm thick film. Correlating this with Hall mobility data, we conclude that exactly
the post-deposition crystallisation is crucial for high mobility. Interestingly, the “high-mobility” phase
reveals a smaller lattice constant than the phase, which crystallises during deposition. The origin of
this difference will be discussed below.
Materials 2019,12, 266 5 of 20
A decrease of intensity of diffraction peaks is observed for the initially crystalline In
2
O
3
phase as
a result of annealing. The most probable reason for that is active recrystallization.
Figure 1b demonstrates the influence of the sputtering pressure on the crystallisation behaviour
of X-ray amorphous films. Apparently, this parameter determines the film texture and Hall mobility.
We attribute this effect to the amount of hydroxyl in a film. Indeed, using a fixed leak rate of the
needle valve feeding water vapour, a higher p
tot
should result in a smaller p(H
2
O) and, hence, in a
smaller concentration of hydroxyl in In
2
O
3
. In turn, the impact of hydroxyl on the crystalline growth
of In
2
O
3
is a known phenomenon, thus, the (400)-orientation can be suppressed if water is present
in a sputtering gas [
27
]. This is consistent with our results. A higher process pressure itself may also
contribute, as it increases the energy transfer from plasma to substrate, resulting in a higher deposition
temperature [28].
Considering the cubic bixbyite structure of In
2
O
3
, one can distinguish two kinds of InO
6
octahedrons which interconnect either by corners or by edges [
29
]. As the octahedrons are directed
along the [111] axes in the lattice, the close packed oxygen layers coincide with the {00l} planes.
This explains why the <00l> oriented films might reveal a smaller electron mobility.
In our further XRD, TEM and SIMS investigations we used the ~150 nm thick films. XRD
measurements (see Figure S2) were performed to estimate the size of coherent scattering in both
lateral (D
lat
) and longitudinal (D
long
) directions using Scherrer’s method [
30
]. In case of lateral size the
asymmetric in-plane XRD measurements at
ψ
= 89.2
◦
were performed. Assuming single strength and
linear dhkl −sin2ψdependence we estimated the residual stress in the films. A complex recalculation
which takes into account the measurement conditions was performed on the basis of a procedure
explained elsewhere [
31
]. The elastic constants for In
2
O
3
were taken from literature [
32
,
33
]. Our results
are collected in Figure S2 and Table 1. Hall measurement data are also presented for the assessment.
Table 1.
Results of the XRD analysis and electron Hall-mobility data for the differently prepared
~150 nm
thick In
2
O
3
films on glass. The Roman numerals are used in the text for easier distinction of
the film states.
States
I. In2O3
As-Deposited
w/o Heating
II. In2O3:H2O
As-Deposited
w/o Heating
III. In2O3:H2O
As-Deposited
Tdep = 160 ◦C
IV. In2O3:H
Tann = 180 ◦C,
Ambient Air
V. In2O3:H
Tann = 230 ◦C,
UHV
Dlong (nm) 35 ±6 - 50 ±15 165 ±6 328 ±5 *
Dlat (nm) 22 ±2 - 32 ±2 230 ±60 205 ±20
Residual Stress (GPa) −2.1 ±0.8 - −1.9 ±0.6 −0.5 ±0.1 −0.6 ±0.2
µe(cm2/V·s) 22 47 41 117 118
Ne(cm−3) 4.58 ×1020 5.85 ×1020 2.03 ×1020 2.24 ×1020 2.57 ×1020
* This value has no real physical meaning, because the film is considerably thinner.
We see that the crystalline state itself does not secure high electron mobility: compare, for instance,
the crystalline state III with the amorphous state II—both films were deposited in the presence of
water vapour. The positive impact of water on
µe
is, however, obvious (compare state I with other
states). This effect is explained in literature by a passivation of dangling bonds with hydrogen that
decreases the electron scattering [
9
]. One can see that such passivation is highly pronounced exactly
in the amorphous state with the utmost amount of dangling bonds (compare state II with I or III).
We observe that
µe
strongly increases with the grain size, compare states III
→
(IV, V). Thus, both
qualities: the degree of crystallinity and the passivation of dangling bonds are important for reaching
a high electron mobility in In2O3.
It should be noted why the concentration of free electrons changes. The presence of water does
not result in a marked change of N
e
(compare states I and II), but high temperature does. This effect
will be discussed below.
One has to notice the presence of residual compressive stress in crystalline films, especially
if crystallisation takes place during deposition (states I and III). The values presented are just an
Materials 2019,12, 266 6 of 20
estimation with a large error, based on the measurements of only two
ψ
values. Two reflexions (222)
and (400) with quite similar Poisson’s ratio were taken into account [33].
It is a fact that in the presence of water the growing In
2
O
3
films contain hydroxyl groups [
9
].
Moreover, the hydroxylation apparently stipulates the amorphous state of In
2
O
3
[
34
]. We have,
however, made some curious observations in our experiments, which cannot be easily explained. Thus,
the In
2
O
3
:H
2
O films (state II in the Table 1) grow predominantly X-ray crystalline on such substrates
like molybdenum-films, i-ZnO, Si-wafer, or CIGS-absorber [
11
]. Moreover, if any additional oxygen is
injected into the sputtering gas, the films becomew crystalline and resistive.
3.2. Presence and Role of Metallic Indium
Since the films appearing as X-ray amorphous could still be nano-crystalline, we undertook TEM
investigations of them. Figure 2shows cross-sectional brightfield TEM images of an In
2
O
3
:H
2
O film
deposited at ptot = 0.5 Pa on glass.
Materials 2019, 12, x FOR PEER REVIEW 6 of 19
3.2. Presence and Role of Metallic Indium
Since the films appearing as X-ray amorphous could still be nano-crystalline, we undertook TEM
investigations of them. Figure 2 shows cross-sectional brightfield TEM images of an In
2
O
3
:H
2
O film
deposited at p
tot
= 0.5 Pa on glass.
(a) (b)
Figure 2. Brightfield TEM images of the as-deposited In
2
O
3
:H
2
O film: (a) first minutes of observation;
(b) after about 20 min under the electron beam.
Several effects can be observed. Initially, the film is amorphous, but it changes under the
electron-beam after about 20 min. A diffraction contrast becomes visible, indicating crystalline
structures. At the same time, droplet-like features at the film/glass interface increase in number and
size. Being aware of the inherent effects of our fabrication procedure, we presume here an effect of
the sample heating by the e-beam. Basically, heating of TEM-lamella due to inelastic interactions with
high energy electrons is a known phenomenon [35]. According to the literature, relatively thick
In
2
O
3
:H films crystallise at 180–200 °C [1–3]. In our special case we deal with much thinner lamella,
where the impact of surface and contact interface is apparently larger. This fact may shift the
crystallisation temperature to lower values. Apart from the crystallisation effect observed, we assume
that the appearance of droplet-like features is related to the presence of metallic indium. Its melting
(T
m
= 156.6 °C) might promote In
2
O
3
:H
2
O crystallisation and results in accumulation of In-droplets.
Actually, the appearance of metallic indium seems to be rather probable in our case as the
metallic phase was found in In
2
O
3
and In
2
O
3
:Sn films by other investigators [36,37]. Consideration of
the In-O phase diagram and general chemistry of indium oxide/hydroxide system also supports this
assumption [34,38].
We applied energy filtered TEM to detect metallic indium, as it exhibits a bulk plasmon with an
energy of about 12 eV [39]. The plasmonic spectrum of In
2
O
3
[40] could not be observed in this study.
The set of TEM images obtained at distinct loss energies is presented in the Supplementary Materials
(Figure S3). Metallic indium appears as bright areas at an electron energy loss of ~12 eV. We also
found that indium congregates either at the film-glass interface or localises in the bulk (see Figure
S4). The latter case is shown in more detail in Figure 3, where the correlation of crystallinity (Figure
3a) and appearance of metallic indium (Figure 3b) can be observed. The top In
2
O
3
layer grows
crystalline due to the impact of RF plasma as we discussed above. The crystalline part consists of
columnar crystallites of 20–40 nm lateral size, which is consistent with the film state III in Table 1.
One can conceivably detect some porosity within this area (see Figure 3a and Figure S3). Metallic
indium particles segregate exactly at the transition region between amorphous and crystalline layers
(Figure 3b). These nanoparticles were found to be crystalline (Figure S5). Obviously, melting of
indium and hence its extraction in a separate phase is associated with indium oxide crystallisation.
The liquid phase is known to support even high quality crystallisation in such methods as liquid
phase epitaxy [41], metal-modulated epitaxy [42], volatile surfactant assisted chemical vapour
deposition [43] and others.
Figure 2.
Brightfield TEM images of the as-deposited In
2
O
3
:H
2
O film: (
a
) first minutes of observation;
(b) after about 20 min under the electron beam.
Several effects can be observed. Initially, the film is amorphous, but it changes under the
electron-beam after about 20 min. A diffraction contrast becomes visible, indicating crystalline
structures. At the same time, droplet-like features at the film/glass interface increase in number
and size. Being aware of the inherent effects of our fabrication procedure, we presume here an effect
of the sample heating by the e-beam. Basically, heating of TEM-lamella due to inelastic interactions
with high energy electrons is a known phenomenon [
35
]. According to the literature, relatively
thick In
2
O
3
:H films crystallise at 180–200
◦
C [
1
–
3
]. In our special case we deal with much thinner
lamella, where the impact of surface and contact interface is apparently larger. This fact may shift the
crystallisation temperature to lower values. Apart from the crystallisation effect observed, we assume
that the appearance of droplet-like features is related to the presence of metallic indium. Its melting
(Tm= 156.6 ◦C) might promote In2O3:H2O crystallisation and results in accumulation of In-droplets.
Actually, the appearance of metallic indium seems to be rather probable in our case as the metallic
phase was found in In
2
O
3
and In
2
O
3
:Sn films by other investigators [
36
,
37
]. Consideration of the
In-O phase diagram and general chemistry of indium oxide/hydroxide system also supports this
assumption [34,38].
We applied energy filtered TEM to detect metallic indium, as it exhibits a bulk plasmon with an
energy of about 12 eV [
39
]. The plasmonic spectrum of In
2
O
3
[
40
] could not be observed in this study.
The set of TEM images obtained at distinct loss energies is presented in the Supplementary Materials
(Figure S3). Metallic indium appears as bright areas at an electron energy loss of ~12 eV. We also
found that indium congregates either at the film-glass interface or localises in the bulk (see Figure S4).
The latter case is shown in more detail in Figure 3, where the correlation of crystallinity (Figure 3a) and
appearance of metallic indium (Figure 3b) can be observed. The top In
2
O
3
layer grows crystalline due
to the impact of RF plasma as we discussed above. The crystalline part consists of columnar crystallites
of 20–40 nm lateral size, which is consistent with the film state III in Table 1. One can conceivably detect
Materials 2019,12, 266 7 of 20
some porosity within this area (see Figures 3a and S3). Metallic indium particles segregate exactly at
the transition region between amorphous and crystalline layers (Figure 3b). These nanoparticles were
found to be crystalline (Figure S5). Obviously, melting of indium and hence its extraction in a separate
phase is associated with indium oxide crystallisation. The liquid phase is known to support even high
quality crystallisation in such methods as liquid phase epitaxy [
41
], metal-modulated epitaxy [
42
],
volatile surfactant assisted chemical vapour deposition [43] and others.
Materials 2019, 12, x FOR PEER REVIEW 7 of 19
(a) (b)
Figure 3. Medium magnification bright field image (a) and EFTEM image at an electron energy loss
of 12 eV (b) obtained on an as-deposited In2O3:H2O film at the same location. Image (a) elucidates two
parts of the film: the bottom ~100 nm thick amorphous part and the top ~50 nm thick crystalline part.
The electron diffraction pattern is obtained for the region, marked by a circle. Bright regions in the
figure (b) correspond to metallic indium.
3.3. Optical Properties of In2O3:H2O and In2O3:H
Optical measurements allow determining properties of the continuous matter. If a material is
crystalline, we receive the information from the interior area of grains, whereas the electrical
properties are cumulative. Many optical investigations of In2O3 and In2O3:H films have been
undertaken [2,36,44,45]. According to S. Joseph and S. Berger, the fitting of IR optical transmission
spectra by effective medium approximation reveals that no changes in In2O3 transmittance should be
observed if the volume fraction of metallic indium remains less than 10%. A larger indium excess
results in a markedly lower transmittance as compared to that experimentally observed [36]. The
presence of a metallic phase can be also deduced from the temperature dependence of the resistivity
[46]. This, however, demands an even higher volumetric content of indium for the percolation of
electrical current.
Figure 4 presents fitted optical spectra for the thin films. The fits were obtained with the help of
the RIG-VM software developed at Fraunhofer IST [47]. A Tauc-Lorentz oscillator has been used for
the fundamental absorption and a Drude term for the free electrons. A free electron mass of m* = 0.28
me has been used [23]. As one can see, the optical mobility matches well with the Hall mobility
(compare with Table 1) for both film states. Some deviation is observed for the crystalline state, where
the Drude term gives a somewhat smaller mobility. This can be attributed to the insufficient spectral
range to fit the plasma resonance of free carriers accurately. A remarkable difference in Ne values
should be noted, namely, the electrical measurements gave an approximately double concentration
of free electrons, which are not visible optically. This means that we might observe an additional
inter-grain source of free electrons.
Figure 3.
Medium magnification bright field image (
a
) and EFTEM image at an electron energy loss
of 12 eV (
b
) obtained on an as-deposited In
2
O
3
:H
2
O film at the same location. Image (
a
) elucidates
two parts of the film: the bottom ~100 nm thick amorphous part and the top ~50 nm thick crystalline
part. The electron diffraction pattern is obtained for the region, marked by a circle. Bright regions in
the figure (b) correspond to metallic indium.
3.3. Optical Properties of In2O3:H2O and In2O3:H
Optical measurements allow determining properties of the continuous matter. If a material
is crystalline, we receive the information from the interior area of grains, whereas the electrical
properties are cumulative. Many optical investigations of In
2
O
3
and In
2
O
3
:H films have been
undertaken
[2,36,44,45]
. According to S. Joseph and S. Berger, the fitting of IR optical transmission
spectra by effective medium approximation reveals that no changes in In
2
O
3
transmittance should be
observed if the volume fraction of metallic indium remains less than 10%. A larger indium excess results
in a markedly lower transmittance as compared to that experimentally observed [
36
]. The presence of
a metallic phase can be also deduced from the temperature dependence of the resistivity [
46
]. This,
however, demands an even higher volumetric content of indium for the percolation of electrical current.
Figure 4presents fitted optical spectra for the thin films. The fits were obtained with the help
of the RIG-VM software developed at Fraunhofer IST [
47
]. A Tauc-Lorentz oscillator has been used
for the fundamental absorption and a Drude term for the free electrons. A free electron mass of
m* = 0.28 mehas been used [23]
. As one can see, the optical mobility matches well with the Hall
mobility (compare with Table 1) for both film states. Some deviation is observed for the crystalline state,
where the Drude term gives a somewhat smaller mobility. This can be attributed to the insufficient
spectral range to fit the plasma resonance of free carriers accurately. A remarkable difference in
N
e
values should be noted, namely, the electrical measurements gave an approximately double
concentration of free electrons, which are not visible optically. This means that we might observe an
additional inter-grain source of free electrons.
Materials 2019,12, 266 8 of 20
Materials 2019, 12, x FOR PEER REVIEW 8 of 19
(a) (b)
Figure 4. Fits (dashed lines) of the optical spectra (solid lines) for ~100 nm thick In2O3:H2O (a) and
In2O3:H (b) films on glass. The μe and Ne values identified from the fitted spectra are given.
Corresponding plasma edges, film thicknesses and mean square errors are as follows: (a) λp = 916.918
nm, d = 103.385 nm, MSE = 0.00757124; (b) λp = 1486.62 nm, d = 104.554 nm, MSE = 0.00893229.
Based on the measurements in the UV range, we determined optical band gaps based on the
Tauc-Lorentz model. These values are presented in Table 2 for the material states I, II and V.
Table 2. Results obtained from the optical absorption of variously prepared In2O3 films.
States →
Band Gap/Thickness ↓
I. In2O3
As-Deposited
w/o Heating
II. In2O3:H2O
As-Deposited
w/o Heating
V. In2O3:H
Tann = 230 °C, UHV
Eg (eV)/100 nm
400 nm
3.55
3.50
3.46
3.48
3.68
3.65
According to the literature, there cannot be an indirect gap in pure In2O3 due to the parabolic
nature of the conduction band [22]. Moreover, the minimum band gap at 2.9 eV is symmetrically
forbidden and the first allowed optical transition occurs from the level ~0.8 eV below the valence
band top that gives the commonly observed value of Eg ≈ 3.7 eV. As presented in Table 2, the as-
deposited crystalline In2O3 films (state I) reveal considerably lower Eg as compared to the expected
value for this material. Amorphous films, which probably contain water (state II), demonstrate even
lower values; however, for 400 nm thick films the difference between states I and II is negligible. This
is consistent with the fact of partial crystallisation observed in thicker In2O3:H2O films (see Figures 1
and 3). However, no influence of water in the state II is then noticed. We do not detect the influence
of expected hydrogen doping in the post-crystallised In2O3:H films (state V) as well, because the
optical Eg values observed are very close to the ones known for the pure In2O3.
3.4. Chemical Changes in In2O3
To understand if there is any chemical transformation during crystallisation and which doping
mechanism is realised, we performed further investigations. According to the FTIR analysis (data are
not presented in this paper) none of the OH or adsorbed water (1615–1630 cm−1) features were
observed in crystalline films. This might mean that these species, if they exist, are concentrated
mainly at the grain boundaries. The IR transmittance, however, differs considerably for different film
states. Thus, an addition of water during deposition results in a significant reduction of the IR
transmittance (states I and II are compared). Annealing in vacuum does not significantly change it
(states II → V), whereas annealing in ambient air provides quite a strong increase of the IR
transmittance (states II → IV). Taking into account our TEM/EELS results, we attribute this behaviour
to the presence of metallic indium in the films deposited in the presence of water. Annealing of such
films in ambient air provides more effective oxidation of indium as compared to the annealing in
vacuum.
Figure 4.
Fits (dashed lines) of the optical spectra (solid lines) for ~100 nm thick
In2O3:H2O (a)
and In
2
O
3
:H (
b
) films on glass. The
µe
and N
e
values identified from the fitted spectra are
given. Corresponding plasma edges, film thicknesses and mean square errors are as follows:
(a)λp= 916.918 nm
,d= 103.385 nm, MSE = 0.00757124; (
b
)
λp
= 1486.62 nm, d= 104.554 nm,
MSE = 0.00893229.
Based on the measurements in the UV range, we determined optical band gaps based on the
Tauc-Lorentz model. These values are presented in Table 2for the material states I, II and V.
Table 2. Results obtained from the optical absorption of variously prepared In2O3films.
States →
Band
Gap/Thickness ↓
I. In2O3
As-Deposited
w/o Heating
II. In2O3:H2O
As-Deposited
w/o Heating
V. In2O3:H
Tann = 230 ◦C, UHV
Eg(eV)/100 nm
400 nm 3.55
3.50 3.46
3.48 3.68
3.65
According to the literature, there cannot be an indirect gap in pure In
2
O
3
due to the parabolic
nature of the conduction band [
22
]. Moreover, the minimum band gap at 2.9 eV is symmetrically
forbidden and the first allowed optical transition occurs from the level ~0.8 eV below the valence band
top that gives the commonly observed value of E
g≈
3.7 eV. As presented in Table 2, the as-deposited
crystalline In
2
O
3
films (state I) reveal considerably lower E
g
as compared to the expected value for
this material. Amorphous films, which probably contain water (state II), demonstrate even lower
values; however, for 400 nm thick films the difference between states I and II is negligible. This is
consistent with the fact of partial crystallisation observed in thicker In
2
O
3
:H
2
O films (see Figures 1
and 3). However, no influence of water in the state II is then noticed. We do not detect the influence of
expected hydrogen doping in the post-crystallised In
2
O
3
:H films (state V) as well, because the optical
Egvalues observed are very close to the ones known for the pure In2O3.
3.4. Chemical Changes in In2O3
To understand if there is any chemical transformation during crystallisation and which doping
mechanism is realised, we performed further investigations. According to the FTIR analysis (data
are not presented in this paper) none of the OH or adsorbed water (1615–1630 cm
−1
) features were
observed in crystalline films. This might mean that these species, if they exist, are concentrated mainly
at the grain boundaries. The IR transmittance, however, differs considerably for different film states.
Thus, an addition of water during deposition results in a significant reduction of the IR transmittance
(states I and II are compared). Annealing in vacuum does not significantly change it (states II
→
V),
whereas annealing in ambient air provides quite a strong increase of the IR transmittance (states II
→
IV). Taking into account our TEM/EELS results, we attribute this behaviour to the presence of metallic
indium in the films deposited in the presence of water. Annealing of such films in ambient air provides
more effective oxidation of indium as compared to the annealing in vacuum.
Materials 2019,12, 266 9 of 20
SIMS was used to determine the In/O and H/O ratios across the film. As we pointed out in
experimental section the DC-sputtered~150 nm thick In
2
O
3
films were analysed in this case. Figure 5a
compares the In/O ratio in as-prepared and annealed films deposited with and without water. As one
can see, In
2
O
3
films contain less oxygen and do not change during annealing in vacuum. According
to our visual observations, In
2
O
3
films are usually darker than In
2
O
3
:H films. Therefore, we assume
the In
2
O
3
to be oxygen deficient. On the contrary, very transparent In
2
O
3
:H films seem to possess a
stoichiometry close to x= 3 in In
2
O
x
. Curiously, we observe the same oxygen enrichment profile in the
upper ~50 nm of both: as-deposited and annealed In
2
O
3
films. At the very surface this enrichment
approaches the overall level observed in the annealed In
2
O
3
:H film. As this effect is insensitive
to annealing, we most probably deal with the impact of a short exposure to the ambient air prior
to the measurement. It is observed only in the case of tiny crystalline, sub-stoichiometric films,
which indicates a reaction in the inter grain space.
Materials 2019, 12, x FOR PEER REVIEW 9 of 19
SIMS was used to determine the In/O and H/O ratios across the film. As we pointed out in
experimental section the DC-sputtered~150 nm thick In2O3 films were analysed in this case. Figure 5a
compares the In/O ratio in as-prepared and annealed films deposited with and without water. As one
can see, In2O3 films contain less oxygen and do not change during annealing in vacuum. According
to our visual observations, In2O3 films are usually darker than In2O3:H films. Therefore, we assume
the In2O3 to be oxygen deficient. On the contrary, very transparent In2O3:H films seem to possess a
stoichiometry close to x = 3 in In2Ox. Curiously, we observe the same oxygen enrichment profile in
the upper ~50 nm of both: as-deposited and annealed In2O3 films. At the very surface this enrichment
approaches the overall level observed in the annealed In2O3:H film. As this effect is insensitive to
annealing, we most probably deal with the impact of a short exposure to the ambient air prior to the
measurement. It is observed only in the case of tiny crystalline, sub-stoichiometric films, which
indicates a reaction in the inter grain space.
Water containing, as-deposited films (state II) have the highest oxygen content among the
investigated systems. It decreases after annealing (state V) together with the increase of transparency
in both UV and near IR spectral regions. We attribute this oxygen loss to the release of water, as the
In/O ratio remains unchanged in the water-free In2O3 sample under the same treatment. Furthermore,
active oxygen diffusion in In2O3 starts at temperatures above 600 °C [48].
For better understanding, we represented the measured data in the form of an H/O ratio (Figure
5b) and the percentage loss of hydrogen and oxygen (see formulas (1)) as a result of the annealing
(Figure 5c). Obviously, hydrogen-to-oxygen ratio is higher in as-deposited state. Its depth profile
demonstrates several pronounced maxima, which correspond to the substrate oscillation and passing
by the inlet of water vapour. This result validates our procedure of hydrogen detection, proving the
satisfactory sensitivity, which we can only achieve for the water-containing films. On the other hand
we realise that hydrogen detected in as-deposited films represents most likely just water. A
comparison of LH and LO discloses an interesting effect: the oxygen loss remains stable over the entire
film thickness, whereas hydrogen releases more actively from the top but demonstrates a stable LH in
the depth. We suggest the release of H2O and H2 species from the film being annealed (Figure 5c), as
only their formation in a free volatile form is chemically possible. Water can evaporate in its free form
if it is contained or released from the hydroxyl groups as shown in the reaction (4) below. Hydrogen
would form only in the presence of metallic indium according to the reactions:
In + 3 H2O → In(OH)3 + 3/2 H2 (2)
In + In(OH)3 → In2O3 + 3/2 H2 (3)
The probability of such reactions and some supporting experimental data published will be
considered in the discussion chapter below.
Considering SIMS results, we could not operate with the absolute values since we did not use
any external standard. However, the qualitative suggestions made were based on internal
standards—indium and oxygen. We realised that the crystallised In2O3:H film, which is a high
mobility TCO to be used in various devices, might suffer from the chemical heterogeneity.
(a)
Materials 2019, 12, x FOR PEER REVIEW 10 of 19
(b) (c)
Figure 5. SIMS depth profiles obtained on various ~150 nm thick DC sputtered In2O3 films. Indium to
oxygen concentration ratios (a) were obtained for the film states I, II and V, which correspond to the
Table 1. State Ia represents the UHV-annealed In2O3 film. Hydrogen to oxygen concentration ratio (b)
and the percentage losses of hydrogen and oxygen (c) are compared for the as-deposited and annealed
states of the film, intentionally containing water.
To observe the processes taking place on the film surface, UPS and XPS measurements were
undertaken. Our XPS measurements (spectra are not shown here) reproduced the results obtained
by Hans F. Wardenga, where In2O3:H2O films revealed a shoulder at the O1s emission line at about
532.6 eV binding energy [9]. This was found to correspond to the OH bonds, which disappear after
annealing.
The UPS spectra acquired during a stepwise increase of temperature from the ambient level (~25
°C) up to 230 °C in UHV show the following changes in the In2O3:H2O film (Figure 6). A ~0.4 eV shift
in the secondary electron edge is observed. The secondary electron edge can be used to determine
the work function of a material according to the relation Wf = Eex − EBsec. Thus, we observe here the
Wf change from ~4.0 eV to ~4.4 eV that basically contradicts the doping phenomenon. It is worth
noticing that the work function of the thermally deposited fully oxidised indium oxide is 5.0 eV [49].
This means that indium in the films in question has a lower oxidation state than in the stoichiometric
oxide. The value of Wf is also determined to a large extent by the crystallographic orientation [50]. In
our case the observed increase of a work function is probably caused by the appearance of a distant
order. The significant difference relative to the fully oxidised In2O3 state can be connected to the
presence of metallic indium in the films (our TEM and FTIR data).
(a) (b)
Figure 6. UPS spectra obtained for the ~150 nm thick In2O3:H2O films during stepwise annealing in
UHV without breaking vacuum. Two regions (a,b) of the same spectra are shown. For comparison,
the intensity was normalised and the background was removed. The excitation energy was 21.2 eV.
Figure 5.
SIMS depth profiles obtained on various ~150 nm thick DC sputtered In
2
O
3
films. Indium to
oxygen concentration ratios (
a
) were obtained for the film states I, II and V, which correspond to the
Table 1. State Ia represents the UHV-annealed In
2
O
3
film. Hydrogen to oxygen concentration ratio (
b
)
and the percentage losses of hydrogen and oxygen (
c
) are compared for the as-deposited and annealed
states of the film, intentionally containing water.
Water containing, as-deposited films (state II) have the highest oxygen content among the
investigated systems. It decreases after annealing (state V) together with the increase of transparency
in both UV and near IR spectral regions. We attribute this oxygen loss to the release of water, as the
In/O ratio remains unchanged in the water-free In
2
O
3
sample under the same treatment. Furthermore,
active oxygen diffusion in In2O3starts at temperatures above 600 ◦C [48].
For better understanding, we represented the measured data in the form of an H/O ratio
(Figure 5b) and the percentage loss of hydrogen and oxygen (see formulas (1)) as a result of the
annealing (Figure 5c). Obviously, hydrogen-to-oxygen ratio is higher in as-deposited state. Its depth
profile demonstrates several pronounced maxima, which correspond to the substrate oscillation and
passing by the inlet of water vapour. This result validates our procedure of hydrogen detection,
Materials 2019,12, 266 10 of 20
proving the satisfactory sensitivity, which we can only achieve for the water-containing films. On the
other hand we realise that hydrogen detected in as-deposited films represents most likely just water.
A comparison of L
H
and L
O
discloses an interesting effect: the oxygen loss remains stable over the
entire film thickness, whereas hydrogen releases more actively from the top but demonstrates a stable
L
H
in the depth. We suggest the release of H
2
O and H
2
species from the film being annealed (Figure 5c),
as only their formation in a free volatile form is chemically possible. Water can evaporate in its
free form if it is contained or released from the hydroxyl groups as shown in the reaction (4) below.
Hydrogen would form only in the presence of metallic indium according to the reactions:
In+3H2O→In(OH)3+ 3/2 H2(2)
In + In(OH)3→In2O3+ 3/2 H2(3)
The probability of such reactions and some supporting experimental data published will be
considered in the discussion chapter below.
Considering SIMS results, we could not operate with the absolute values since we did not
use any external standard. However, the qualitative suggestions made were based on internal
standards—indium and oxygen. We realised that the crystallised In
2
O
3
:H film, which is a high
mobility TCO to be used in various devices, might suffer from the chemical heterogeneity.
To observe the processes taking place on the film surface, UPS and XPS measurements were
undertaken. Our XPS measurements (spectra are not shown here) reproduced the results obtained
by Hans F. Wardenga, where In
2
O
3
:H
2
O films revealed a shoulder at the O1s emission line at about
532.6 eV binding energy [
9
]. This was found to correspond to the OH bonds, which disappear
after annealing.
The UPS spectra acquired during a stepwise increase of temperature from the ambient level
(~25 ◦C)
up to 230
◦
C in UHV show the following changes in the In
2
O
3
:H
2
O film (Figure 6). A ~0.4 eV
shift in the secondary electron edge is observed. The secondary electron edge can be used to determine
the work function of a material according to the relation W
f
=E
ex −
EB
sec
. Thus, we observe here
the Wfchange from ~4.0 eV to ~4.4 eV that basically contradicts the doping phenomenon. It is worth
noticing that the work function of the thermally deposited fully oxidised indium oxide is 5.0 eV [
49
].
This means that indium in the films in question has a lower oxidation state than in the stoichiometric
oxide. The value of W
f
is also determined to a large extent by the crystallographic orientation [
50
].
In our case the observed increase of a work function is probably caused by the appearance of a distant
order. The significant difference relative to the fully oxidised In
2
O
3
state can be connected to the
presence of metallic indium in the films (our TEM and FTIR data).
Materials 2019, 12, x FOR PEER REVIEW 10 of 19
(b) (c)
Figure 5. SIMS depth profiles obtained on various ~150 nm thick DC sputtered In2O3 films. Indium to
oxygen concentration ratios (a) were obtained for the film states I, II and V, which correspond to the
Table 1. State Ia represents the UHV-annealed In2O3 film. Hydrogen to oxygen concentration ratio (b)
and the percentage losses of hydrogen and oxygen (c) are compared for the as-deposited and annealed
states of the film, intentionally containing water.
To observe the processes taking place on the film surface, UPS and XPS measurements were
undertaken. Our XPS measurements (spectra are not shown here) reproduced the results obtained
by Hans F. Wardenga, where In2O3:H2O films revealed a shoulder at the O1s emission line at about
532.6 eV binding energy [9]. This was found to correspond to the OH bonds, which disappear after
annealing.
The UPS spectra acquired during a stepwise increase of temperature from the ambient level (~25
°C) up to 230 °C in UHV show the following changes in the In2O3:H2O film (Figure 6). A ~0.4 eV shift
in the secondary electron edge is observed. The secondary electron edge can be used to determine
the work function of a material according to the relation Wf = Eex − EBsec. Thus, we observe here the
Wf change from ~4.0 eV to ~4.4 eV that basically contradicts the doping phenomenon. It is worth
noticing that the work function of the thermally deposited fully oxidised indium oxide is 5.0 eV [49].
This means that indium in the films in question has a lower oxidation state than in the stoichiometric
oxide. The value of Wf is also determined to a large extent by the crystallographic orientation [50]. In
our case the observed increase of a work function is probably caused by the appearance of a distant
order. The significant difference relative to the fully oxidised In2O3 state can be connected to the
presence of metallic indium in the films (our TEM and FTIR data).
(a) (b)
Figure 6. UPS spectra obtained for the ~150 nm thick In2O3:H2O films during stepwise annealing in
UHV without breaking vacuum. Two regions (a,b) of the same spectra are shown. For comparison,
the intensity was normalised and the background was removed. The excitation energy was 21.2 eV.
Figure 6.
UPS spectra obtained for the ~150 nm thick In
2
O
3
:H
2
O films during stepwise annealing in
UHV without breaking vacuum. Two regions (
a
,
b
) of the same spectra are shown. For comparison,
the intensity was normalised and the background was removed. The excitation energy was 21.2 eV.
Materials 2019,12, 266 11 of 20
According to the Figure 6b, vacuum annealing causes a shift of the valence band edge by ~0.2 eV.
All observed changes are be depicted on the energy diagram, where the UPS data are used to fix
the E
F
and E
VB
levels (Figure 7). Here we used a caption E
g, min
(minimum) for the fundamental
band gap, which is known to be 2.9 eV for the pure In
2
O
3
[
23
]. If we admit any doping in our films,
it can be even smaller due to the band gap narrowing phenomenon [
51
,
52
]. The optical E
g, opt
values
obtained in this study were placed in accordance with the principle described above [
22
]. These energy
diagrams show that the Fermi level is very close to the conduction band in both materials: amorphous
and crystalline. If we admit the same fundamental band gap width for both states, then the latter
would be a non-degenerate semiconductor that should not be the case for such high concentration
of free electrons. Since we observed a Burstein-Moss shift as a result of the annealing, the In
2
O
3
:H
likely remains degenerate due to the band gap shrinkage. Major changes happen with a level of the
allowed optical transition inside the valence band, namely, it shifts markedly downwards. This can be
attributed to the effect of crystallisation as the valence band contains fully occupied 2pand 2soxygen
states and empty 4d-indium states [22].
Materials 2019, 12, x FOR PEER REVIEW 11 of 19
According to the Figure 6b, vacuum annealing causes a shift of the valence band edge by ~0.2
eV. All observed changes are be depicted on the energy diagram, where the UPS data are used to fix
the E
F
and E
VB
levels (Figure 7). Here we used a caption E
g, min
(minimum) for the fundamental band
gap, which is known to be 2.9 eV for the pure In
2
O
3
[23]. If we admit any doping in our films, it can
be even smaller due to the band gap narrowing phenomenon [51,52]. The optical E
g, opt
values
obtained in this study were placed in accordance with the principle described above [22]. These
energy diagrams show that the Fermi level is very close to the conduction band in both materials:
amorphous and crystalline. If we admit the same fundamental band gap width for both states, then
the latter would be a non-degenerate semiconductor that should not be the case for such high
concentration of free electrons. Since we observed a Burstein-Moss shift as a result of the annealing,
the In
2
O
3
:H likely remains degenerate due to the band gap shrinkage. Major changes happen with a
level of the allowed optical transition inside the valence band, namely, it shifts markedly downwards.
This can be attributed to the effect of crystallisation as the valence band contains fully occupied 2p
and 2s oxygen states and empty 4d-indium states [22].
Figure 7. Energy diagram for indium oxide of the states II and V. The scheme is created on the basis
of optical (UV-Vis) and UPS data.
The observations made require a more detailed discussion of the In
2
O
3
chemistry and possible
origin of doping.
4. Discussion
To understand the results obtained in this work, we have to review some basic properties of
indium, In
2
O
3
and In(OH)
3
.
4.1. Appearance of Metallic Indium in In
2
O
3
The electro-chemical potential of metallic indium is φ
0
= −0.3382 V [53]. This means that under
normal conditions, metallic indium should not reduce protons in an aqueous solution to molecular
hydrogen. Nevertheless, the potential is not too high and both reactions, reduction of metal and
reduction of hydrogen, may proceed simultaneously on a competitive basis.
According to the In-O phase diagram, there is no detectable phase of oxygen non-stoichiometry
[38]. If any In-excess is provided (≥0.02 at.%), there is a mixture of two phases: In
2
O
3
and metallic In,
which is solid below 156.634 °C and liquid above this temperature. Indium (III) oxide is
thermodynamically very stable over a wide range of T and p(O
2
). According to the Ellingham
diagram, one needs a p(O
2
) of about 10
−100
atm. in order to reduce it to metallic indium at room
Figure 7.
Energy diagram for indium oxide of the states II and V. The scheme is created on the basis of
optical (UV-Vis) and UPS data.
The observations made require a more detailed discussion of the In
2
O
3
chemistry and possible
origin of doping.
4. Discussion
To understand the results obtained in this work, we have to review some basic properties of
indium, In2O3and In(OH)3.
4.1. Appearance of Metallic Indium in In2O3
The electro-chemical potential of metallic indium is
ϕ0
=
−
0.3382 V [
53
]. This means that under
normal conditions, metallic indium should not reduce protons in an aqueous solution to molecular
hydrogen. Nevertheless, the potential is not too high and both reactions, reduction of metal and
reduction of hydrogen, may proceed simultaneously on a competitive basis.
According to the In-O phase diagram, there is no detectable phase of oxygen
non-stoichiometry [
38
]. If any In-excess is provided (
≥
0.02 at.%), there is a mixture of two phases:
In
2
O
3
and metallic In, which is solid below 156.634
◦
C and liquid above this temperature. Indium
Materials 2019,12, 266 12 of 20
(III) oxide is thermodynamically very stable over a wide range of Tand p(O
2
). According to the
Ellingham diagram, one needs a p(O
2
) of about 10
−100
atm. in order to reduce it to metallic indium
at room temperature. The equilibrium oxygen partial pressure at 200
◦
C is about 10
−55
atm. At the
same time, metallic indium remains stable in air and starts to oxidise visibly only after melting.
The oxidation of indium in a liquid form proceeds about five time faster as compared to the solid [
54
].
Aside from the main oxidation state +3, indium may have also +2 and +1 in combination with oxygen.
A formal oxidation state +2 is most probably a mixture of diamagnetic +1 (5s
2
) and +3 (5s
0
) forms,
as no experimental evidence of magnetism in reduced indium oxides was detected. The theoretical
investigation of hypothetical neutral, molecular In-O clusters with different In/O ratios reveals their
high instability in an ionic environment [
55
]. The HOMO–LUMO gap was found to depend on the
metal-to-oxygen ratio in the cluster. Oxidation is likely unfavourable when the In/O ratio is larger
than 1, as both vertical and adiabatic electron affinities are negative for In
2
O. In oxygen-deficient In
2
O
3
films, metallic indium may form as a result of oversaturation by cooling down after deposition at
elevated temperature [36]. In this case, indium precipitates according to HRTEM and EELS in a form
of 5–30 nm nano-particles independently of an indium excess. It is known that intermediate oxides
disproportionate in contact with water, resulting in In
2
O
3
and metallic indium [
34
]. These data confirm
that metallic indium readily forms if any lack of oxygen and/or water is provided.
To understand the chemical impact of water during sputtering, let us briefly consider the plasma
chemistry of water. Basically, low total pressure and especially plasma excitation change the chemical
activity of water. Upon photo-ionisation, water vapour becomes a weakly ionised plasma consisting of
electrons and H
2
O
+
[
56
]. In the highest state of excitation, the plasma consists of e
−
, H
+
, and O(
n+
).
In a general case of RF-sputtering from the ceramic target, mostly M
+
and MO
+
charged species are
observed [
57
]. In the case of In
2
O
3
, the RF plasma should contain these species in a ratio M
+
/MO
+
of
more than 30. Since this ratio depends on the M-O binding energy, we took the value known for Fe
2
O
3
implying that the M-O binding energy is quite close for Fe
2
O
3
(
∆
G
◦298
=
−
732.1 kJ/mol) and In
2
O
3
(
∆
G
◦298
=
−
809.3 kJ/mol) [
58
,
59
]. When argon is used as a sputtering gas, almost no O
+
, but mostly
neutral oxygen is observed [
58
,
60
]. The RF-plasma above the ZnO target has a similar content; however,
oxygen species generated during DC sputtering of ZnO are O
−
, O
2−
, and O. The content of negatively
charged oxygen species increases exponentially with the reduction of the total pressure [
61
]. The main
difference between the RF and the DC process lies in the concentration of electrons, which is much
higher in the first case. Thus, in our process, we likely deal with an intermediate oxidation state of
indium in the absence of strong oxidants in the plasma, which finally yields metallic indium species in
a film.
4.2. Water Containing In2O3
Despite a chemical impact as hydroxylation, water stipulates the amorphous state of as-deposited
films. It is known that indium (III) hydroxide tends to remain jelly or even forms a colloid
in aqueous solutions rather than precipitating in a crystalline form. The main reasons for that
are the donor-acceptor interaction, typical for metals having free 3dorbitals, and the hydrogen
bonding in hydroxides. As it is known for the most investigated analogue—aluminium hydroxide,
such parameters as concentration, temperature and pH determine the hydrolysis, peptisation,
aging and, finally, crystallisation [
62
]. Basically, al least three processes are coupled with water
release and formation of many networking chemical bonds:
=In-OH + HO-In= →=In-O-In= + H2O (4)
According to the thermal gravimetric analysis, the crystalline In(OH)
3
transforms into In
2
O
3
with
water elimination, starting very slowly from T
≥
200
◦
C and becomes fast at about 230
◦
C in an inert
1 bar atmosphere [
63
]. As per Le Chatelier’s principle, the dissociation in vacuum likely proceeds at
lower temperature.
Materials 2019,12, 266 13 of 20
It is known that gaseous hydrogen can also be successfully applied as a hydrogenation agent
yielding high-mobility In
2
O
3
films [
64
]. In this case, the films were obtained by RF sputtering in
an amorphous state as well and crystallised by post-deposition annealing. Thorough investigation
of oxygen and hydrogen desorption from the In
2
O
3
powders with different surface areas serves
us with the following observations [
10
]. Surface hydrogen starts to desorb in a high vacuum
(p= 5 ×10−7mbar)
already at temperatures somewhat below 100
◦
C. Desorption of stronger bound
hydrogen starts at ~150
◦
C. Water (6 mbar in 1 bar He) becomes an active re-oxidation agent at
temperatures higher than about 250
◦
C, whereas dry oxygen (1 bar) actively re-oxidises the surface
starting from >150 ◦C.
Thus we realise that hydroxylation of In
2
O
3
is most likely the reason for the amorphous state
of as-deposited films. This effect can be achieved via sputtering in the presence of either hydrogen
or water. Hydrogen acts even more reproducibly [
64
], since it probably delivers just a necessary
hydroxylation without any water excess. However, we still need to understand the desorption of
chemically different hydrogen. Additionally, the effect of In
2
O
3
reduction in hydrogen on its electrical
conductivity should be considered.
4.3. Doping and Conductivity of In2O3
It is known that oxygen deficiency in In
2
O
3
causes higher conductivity [
10
,
46
,
65
]. According to
the impedance measurements performed on In
2
O
3
polycrystalline samples, their reduction in dry
hydrogen results in a slow resistance decline, starting already at room temperature. The resistance falls
sharply at a temperature somewhat below 100
◦
C and further decreases much slower up to its minimum
at about 250
◦
C [
10
]. In the presence of water, the same dependency is observed, but the temperatures
are about 50
◦
C higher. This change is reversible and matches the hydrogen adsorption/desorption
data; however, the resistance of the re-oxidised samples was found to be
4–5 orders
of magnitude
lower than the initial one. It might point to the inter-grain changes, e.g., In
2
O
3
reduction, which is
then encapsulated by the fully oxidised shell.
We already showed above that the concentration of electrons and hence electron mobility in
In
2
O
3
films were often found to be determined by the size of crystallites [
23
]. In the literature
this effect is attributed to the so-called unintentional doping, which is supposedly caused by the
inter-grain diffusion of water from ambient air [
66
,
67
]. The mechanism of such doping, however,
remains questionable for us.
Theoretical studies of this matter demonstrate quite discrepant conclusions. Some basic
description of the defect chemistry in In
2
O
3
was done, using solid state chemistry [
68
]. However,
most modern investigations being aimed to justify which point defects exist in the material are
performed using density functional theory (DFT). Thus, according to J. Liu, who used the GGA + U
formalism, the most stable point defects in In
2
O
3
crystals are oxygen vacancies of the anti-Frenkel
type (V
O••
-O
i00
) [
69
]. According to the LDA and LDA+U functional calculations, the formation
energy of V
In000
was found to be very low in n-type In
2
O
3
[
70
]. Considering hydrogen doped In
2
O
3
,
S. Limpijumnong with co-workers suggested H
i•
and H
O•
as the main donor defects in In
2
O
3
:H
rather than V
O••
[
71
]. A combination of theory with the muon rotation/relaxation spectroscopy
revealed that the charge neutrality level (CNL) for hydrogen in In
2
O
3
lies above the conductive band
minimum (CBM), thus providing a shallow donor level with an activation energy of 47
±
6 meV [
66
].
This study also stated that hydrogen often becomes an unintentional donor in many polycrystalline
oxides. According to T. Tomita et al., who used first-principles molecular orbital calculations, the
interstitial indium ions (In
i•
) are the native donors in In
2
O
3
[
72
]. These defects may only coexist with
V
O
(no charge was noticed in the original work), which facilitate the emergence of indium donors as
shallow states.
Considering the penetration of hydrogen into indium oxide, we have to take into account
the classical approach of experimentally obtained ionic radii. Oxygen ions (O
2−
) with tetrahedral
coordination, like in the In
2
O
3
bixbyte structure, have an effective ionic radius of 1.38 Å [
73
]. The OH
−
Materials 2019,12, 266 14 of 20
group would have an even smaller (1.35 Å) ionic radius in this coordination, since the proton is
actually a pristine positive nucleus, which is drawn in to the negative electron shell of oxygen, thus
making the Coulomb repulsion between neighbouring oxygen ions smaller. On the basis of this simple
consideration it is hard to imagine that the hydrogen proton or even the neutral H atom, having a Bohr
radius a
0≈
0.53 Å, can replace oxygen in its site. Thus we recognise that the hydroxyl (OH
−
)
O•
is the
most probable hydrogen containing species in In
2
O
3
. It is known, however, that In
+3
also shows the
atomic absorption of hydrogen [
74
]. Hydrogen was also found to be readily adsorbed by an indium
rich InP surface [
75
]. A large thermodynamic driving force for the neutral covalent binding between
hydrogen and solid state indium dimers was identified.
The existence of interstitial indium ions has also some restrictions. Thus, In
+3
in octahedral
coordination possesses, an ionic radius of 0.8 Å is rather large to squeeze into the cavities of the
bixbyite lattice. A lower indium oxidation state, e.g., +1 (In
i•
), corresponds to an even larger ionic
radius. The intermediate oxidation states are furthermore electrochemically unstable (see above).
However, ferromagnetism was observed in oxygen deficient InO
x
films annealed in UHV at
600 ◦C [65]
.
This effect was found to be accompanied by the In-In clustering and formation of highly defective
glassy regions in crystalline In
2
O
3
[
65
]. According to Preissler and Bierwagen, the existence of doubly
ionized donors best describes the ionized impurity scattering in unintentionally doped In
2
O
3
[
23
].
Attributing this circumstance to the existence of an indium excess, it is not unlikely to suggest such a
defect as In
O••
, which is the non-oxidised indium at the oxygen site. It means that we might basically
have InIn×−InO•• −InIn×clusters with an effective In+2 oxidation state.
As for oxygen vacancies, the main disagreement in literature concerns their energy level,
which represents either deep [
76
] or shallow states [
23
,
67
]. A practical way to discover which
point defects provide conductivity in oxide materials is to measure the conductivity or better the
N
e
dependence on p(O
2
). The main restriction on that is the requirement of an equilibrium, which
for metal oxides means quite high temperatures, far beyond the typical 200–250
◦
C for In
2
O
3
:H.
So the measurements performed at 800
◦
C discovered the
σ∝
p(O
2
)
1/10
dependence for In
2
O
3
.
Authors attributed this dependence to the (In
i•••
-O
i00
)
•
cluster formation [
77
]. In other work Hall
measurements are presented for In
2
O
3
films obtained at different p(O
2
) by RF sputtering without
intentional heating [
78
]. Conductivity was found to be rather constant (~3
×
10
3Ω−1
cm
−1
) at low
oxygen partial pressure (<8
×
10
−4
Pa). When p(O
2
) increases, conductivity sharply drops over
some orders of magnitude that is mostly caused by a decrease of N
e
from ~10
−19
to 10
−16
cm
−3
(at p(O2)≈10−3Pa)
. The authors suggested oxygen vacancies as the major donor defects. In this case
the N
e
should have revealed the slope
∝
p(O
2
)
−1/6
for the very deficient oxide and
∝
p(O
2
)
−1/4
for
the almost stoichiometric one. From the data presented in this paper one can derive a N
e∝
p(O
2
)
−9
correlation, which cannot be explained by the defect chemistry. This can be rather easier attributed to
the presence of a metallic indium phase. There is another important point supporting this hypothesis.
The work function, measured for metallic indium, varies in the range from ~3.9 to ~4 eV, depending
on temperature [
79
]. This value is very close to the one measured for In
2
O
3
:H
2
O and slightly smaller
as compared to the one measured for In
2
O
3
:H films (see Figure 7). This means that free electrons can
easily be injected from the In0outer shell into the conduction band of both oxides.
We would like to point also at the very interesting effect of photo-induced change in reactively
DC-sputtered amorphous In
2
O
3
films: an exposure of
≤
100 nm thick films to UV light (h
ν≥
3.0 eV)
resulted in a stable increase of conductivity by
×
10
8
reaching
σ≥
10
3Ω−1
cm
−1
[
80
]. Simultaneously,
the absorption coefficient increases by up to a factor of 10
3
for hv < 1.5 eV and the absorption edge
shifts by +0.1 eV. A Drude approximation of the optical absorption in the near IR region gives
Ne= 1.5 ×1020 cm−3
that agrees with the Hall data. These data actually represent the pure effect of
In
2
O
3
reduction without hydrogenation/hydroxylation impact. They reproduce to some extent our
results; however, the Burstein-Moss shift observed in presence of hydrogen is about 0.1 larger.
Materials 2019,12, 266 15 of 20
4.4. High-Mobility of In2O3:H
After T. Koida, the high mobility In
2
O
3
is widely accepted to be doped by hydrogen. He also
stated that the doubly ionised impurities were exchanged by singly ionised ones during the annealing
process that results in about twofold reduction of N
e
[
44
]. The in-situ Hall measurements performed
by H. F. Wardenga et al. during annealing of as-deposited In
2
O
3
:H films in vacuum have allowed
underlining the following stages [
9
]. The first stage elapsing at about 160
◦
C is accompanied with
a slight decrease of
µe
, which occurs, as supposed, due to the phonon scattering being expected for
the degenerated semiconductors. Within this stage N
e
remains settled. During further heating from
160
◦
C up to 200
◦
C, the N
e
increases and
µe
remains unchanged. At T > 250
◦
C, N
e
starts declining
fast and a strong increase of
µe
takes place. The authors suggest that the driving force of the rising
N
e
is crystallisation followed by the grain growth. In turn, the depletion (decrease of N
e
) at grain
boundaries is to be the reason of a measurable depletion in a material with small grains. Crystallisation
and grain growth are superimposed by the decomposition of In(OH)
3
. According to the authors,
the release of oxygen is responsible for the drop in carrier concentration and the grain boundaries are
being saturated by hydrogen, closing dangling bonds.
We may not fully agree with this explanation mainly because of the known fact that hydrogen
disappears first from the inter-grain space. We believe therefore that the dangling bonds existing at
grain boundaries are most probably eliminated by the reaction (4). Moreover, it is known that the
undoped single crystalline In
2
O
3
reveals electron mobility exceeding 200 cm
2
/Vs that is restricted
by ~270 cm
2
/Vs due to the phonon scattering [
23
]. On the other hand the unintentionally doped
polycrystalline samples can demonstrate similarly high
µe
as the hydrogen doped ones [
44
,
81
].
According to T. Koida the effective mass in In
2
O
3
:H seems to depend on N
e
mostly, rather than
on crystallinity [
44
]. Many of the experimental data collected for various In
2
O
3
based systems reveal
a plateau on
µe
=f(N
e
) dependency exactly around N
e
~ 10
20
cm
−3
[
23
]. This phenomenon is also
associated with a large spread of mobilities indicating additional scattering due to imperfections in the
crystal for the samples with µe< 130 cm2/Vs.
5. Conclusions
To conclude, we observed that the free charge carriers in both In
2
O
3
and In
2
O
3
:H films can
appear due to the presence of In
0
. We suggest that metallic indium is present in as-deposited In
2
O
3
or
In
2
O
3
:H
2
O films in a much, up to the atomic level, dispersed state. The presence of water or hydrogen
during In
2
O
3
deposition at low temperature secures the amorphous state of the film. Hydroxylation
of In
2
O
3
is probably the main reason for that. Crystallisation of such films starts at ~160
◦
C when
the excess of indium agglomerates, releasing in a separate nano-crystalline phase due to the melting.
Melted indium species may vanish during annealing in two ways: either via evaporation and oxidation
by water in UHV or via oxidation by oxygen in air. Thus, the concentration of free electrons in In
2
O
3
matrix is reduced and the near IR transparency increases. Both processes, however, do not provide
high mobility. The laterally extended growth of crystallites happens when water is released as a result
of the hydroxide
→
oxide transformation. Growing crystallites interconnect at grain boundaries by the
In-O-In bonds. These factors both provide high electron mobility exceeding 100cm
2
/Vs. According
to our experimental observation, annealing in air demands lower temperature (~180
◦
C) to provide
high mobility as compared to the annealing in UHV (>220
◦
C). We attribute this to the higher water
content in the former case. Crystallisation of the In
2
O
3
:H
2
O system is accompanied with the doping
of In
2
O
3
. We expect that the “unintentional” doping differs from the “hydrogen” doping as follows.
In the first case a spontaneous injection of free charge carriers from the dispersed In
0
metallic species
concentrated in the inter-grain defect-rich spaces takes place. In the second case, we likely deal with
oxidation of the intra-grain In0defects trapped during crystallisation by Schottky vacancies:
InO•• +1
/2O2+ VIn000 →InIn×+ OO×+e0(5)
Materials 2019,12, 266 16 of 20
InO•• + H2O+VIn000 →InIn×+ OHO•+1
/2H2+ 2e0(6)
Reactions (5) and (6) describe oxidation by oxygen and water, respectively. High temperature
makes water an oxidizing agent, whereas low pressure facilitates hydrogen removal. According to our
SIMS results, gaseous hydrogen is removed from the film mostly from the top ~50 nm layer. Metallic
indium can also accumulate hydrogen in the bulk of the film. As we saw, OH
O•
defects most probably
also exist in In
2
O
3
:H films but their formation in the absence of In
0
is not associated with any redox
reaction, so in that case they do not donate electrons.
Supplementary Materials:
The following are available online at http://www.mdpi.com/1996-1944/12/2/266/s1,
Figure S1: SEM cross-section of the RF sputtered (p
tot
= 0.5 Pa) 500 nm In
2
O
3
:H
2
O film obtained on Si-substrate,
Figure S2: XRD patterns for ~150 nm In
2
O
3
films deposited (RF-sputtering, p
tot
= 0.5 Pa) on glass: comparison of
crystallisation conditions. Roman numerals correspond to the film states discussed in the text. Diffraction patterns
were acquired using detector-scanning at grazing incidence in the out-of-plane (a) and in-plane (b) modes, Figure S3:
TEM images obtained using energy filter. The set energy is marked on each image, Figure S4: Cross-sectional TEM
images acquired with electrons having 12 eV energy loss on the sample. The bright areas correspond therefore to
metallic indium in as-deposited In
2
O
3
:H
2
O film. Two types of indium segregation: on the film/glass interface (a)
and within the bulk of the film (b) are observed, Figure S5: TEM image of metallic indium nanoparticles released
in In2O3:H2O matrix. The lattice fringe contrast observed reveals their crystalline state.
Author Contributions:
R.M. planned and organised some (SIMS, TEM/EELS) measurements, developed XRD,
UPS, TEM and optical data, developed ideas and written the paper. A.S. initiated the work, deposited rf-films
and performed the XPS, UPS and XRD measurements. M.W. performed the TEM/EELS investigation and
provided support with the corresponding text. P.P.M. performed the SIMS measurements and provided support
with the text. U.B. prepared the samples for the TEM investigation and performed preliminary TEM analysis.
A.P. developed the optical model for optical data. D.E. deposited the dc-films. S.K. took part in many discussions
and provided support with his experience. R.K., I.L. and B.S. are the leaders of the groups involved; they provided
support with their expertise as well.
Funding:
This work was funded by the German Federal Ministry for Economic Affairs and Energy under contract
number 0325762 (TCO4CIGS). We acknowledge support by the German Research Foundation and the Open
Access Publication Funds of TU Berlin.
Acknowledgments:
The authors express special thanks to their colleagues from the Bruker AXS application lab
in Karlsruhe, namely, to Fernando Rinaldi and to Wolfram Pitschke for the measurements on high-end XRD
equipment. We thank Christoph Genzel (Department for Microstructure and Residual Stress Analysis, BESSY
II, HZB) for professional support in XRD data processing. Carola Klimm (Institute for Silicon Photovoltaics,
HZB) is gratefully acknowledged for the SEM investigation which was very helpful for this work in general but
went beyond this manuscript. We acknowledge Johanna Reck (Optotransmitter-Umweltschutz-Technologie e.V.)
for FTIR measurements and thorough discussion of the optical data. The authors also thank Daniel Abou-Ras
(Department for Nanoscaled Structures and Microscope Analysis, HZB) for fruitful discussions and for making
helpful remarks.
Conflicts of Interest: There are no conflict of interests to declare.
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