surface science
Surface Science 372 (1997) 289-299
ELSEVIER
.
!
Surface properties of indium-doped Cd2SnO 4 ceramics studied by
EELS and photoemission spectroscopy
Y. D o u , R.G. Egdell *
Inorganic chemistry Laboratory, South Parks Road, Oxford OX1 3QR, UK
Received 6 M a y 1996; accepted for publication 3 September 1996
Abstract
Ceramic samples of indium-doped dicadmium stannate (Cd2SnO4) have been studied over the doping range 0.5-2.0 at% In by
electron energy loss spectroscopy (EELS) and ultraviolet and X-ray photoemission spectroscopy (UPS and XPS). The samples were
prepared from SnO2 and Cdl_~InxO, which was itself made from CdO and In203. The maximum surface plasmon energy is around
0.450 eV. The carder concentration derived from the plasmon energy is always lower than that estimated from the bulk dopant
concentration. The atomic percentages of both cadmium and tin on the surface before annealing, determined from XPS, are lower
than those in undoped Cd2SnO4, giving evidence for the substitution of In 3+ ions on both Cd 2+ and Sn 4+ sites. XPS shows a
pronounced decrease in the surface concentration of segregated indium after annealing under UHV conditions. He(I) UPS shows a
well defined conduction band feature with a higher intensity than that for tmdoped samples. It is evident that the lower edge of the
valence band shifts towards higher binding energy as the doping level increases.
Keywords: Cadmium oxides; Ceramics; Electron energy loss spectroscopy; Surface segregation; Tin oxides; UV photoelectron
spectroscopy; X-ray photoelectron spectroscopy
1. Introduction
The binary post transition metal oxides SnO2,
and CdO are extrinsic n-type semiconductors. Both SnO2 and In203 have bandgaps in the
near UV (Egap=3.62eV for S n O 2 [ 1 ] ; E g a p =
3.75 eV for In203 [-2]). When doped, these n-type
materials show a unique combination of electrical
and optical properties: they are electrically conductive, but at the same time they are transparent to
visible radiation and reflective to infrared radiation
[3]. The infrared reflectivity is dependent upon
the carrier concentration: carriers can be introduced by doping the parent compound either by
In203
* Corresponding author. E-mail: egdell@vax.ox.ac.uk
creating oxygen vacancies or by substituting pentavalent Sb 5+ for Sn4+ ions or quadrivalent Sn4+
for In 3+ ions. The carrier concentration in turn
determines the plasmon energy. Maximum plasmon energies of 0.59 eV [-4] and 0.63 eV [5] have
been achieved in ceramic samples of Sb-doped
SnO 2 and Sn-doped In203, respectively. Recently,
we found [6] that the maximum plasmon energy
in ceramic CdO doped with trivalent In 3+ ions is
0.72 eV and exceeds that for Sb-doped SnO2 and
Sn-doped In203. However, CdO is not useful as a
transparent conducting material because of its
relatively small bandgap (Eg,p=0.55 eV for CdO
at 300K [7]) which results in excessive light
absorption from about 0.5#m down to the
ultraviolet.
0039-6028/97/$17.00 Copyright © 1997 Elsevier Science B.V. All fights reserved
PII S0039-6028 ( 9 6 ) 01141-7
290
I:. Dou. R.G. Egdell/ Surface Science 372 (1997) 289-299
The ternary oxides, Cd2SnO4, CdSnO 3 and
CdIn20 4 also have attractive electrical and optical
properties. They are also extrinsic semiconductors
with relatively high electrical conductivity, and
their bandgaps (Egap=2.06 eV for Cd2SnO 4 1,8],
Eg~p=2.23 eV for Cdln20 4 [9] and Egap=2.3 eV
for CdSnO3 [10]) are bigger than for CdO itself.
The carrier concentration in these materials can
also be tuned either, by creating oxygen deficiencies
or by chemical substitution. These ternary oxides,
therefore, stand alongside SnO2 and In20 3 as
potential transparent conducting materials and
they have received extensive study over the past
few years. Dicadmium stannate, Cd2SnO4, is the
most intensively investigated material among these
ternary oxides 1'11 ]. It has a relatively high carrier
mobility and its visible absorption edge moves to
much higher energy with n-type doping owing to
the exceptionally large Moss-Burnstein shift associated with the low carrier effective mass [8,12,13].
This results in an increasing window of transparency in the U.Vregion and increasing reflectance
in the IR region with: increasing electrical
conductivity.
Polycrystalline Cd2SnO4 adopts an orthorhombic structure of the SI:2PbO 4 type 1,~14], ila which
SnO6 octahedra form chains by: sharing edges
along the 1-001] direction and these linear chains
are held together by CdO 7 polyhedra (Fig. 1), The
Sn 4+ cation keeps roughly the same coordination
environment as in the tin oxide SnO2. However,
the Cd 2+ site geometry represents a deviation from
the octahedral symmetry which it holds in cadmium oxide CdO itself. By analogy with the binary
systems discussed above, it is to be expected that
substitution of In into Cd sites or Sb into Sn sites
in Cd2SnO 4 will result in n-type doping.
The effect of n-type doping on the electrical
properties of Cd2SnO 4 thin films 1'15,16] and thick
films 1,17]~ have been extensively studied. There
have also been previous studies of antimony-doped
Cd2SnO4 (Cd2Snl_~SbxO4) ceramics by NMR,
E P R a n d Mrssbauer spectroscopies [18] and by
conductivity measurements~ [ 19]. Recently, we
reported a detailed study of Sb doping in Cd2SnO 4
ceramics by electron energy loss spectroscopy
(EELS) and X-ray and ultraviolet photoelectron
spectroscopy (XPS and UPS) [20]. It was shown
b
..~
C
a
•
0
Cd
0
Fig. 1. The structure of Cd2SnO4, showingthe chains of edgesharing SnO 6 octahedra, O (open circles) and Cd ions
(solid circles).
that EELS combined with UPS and XPS provided
direct information on the effect of doping on the
surface electronic structure.
By analogy with In doping in CdO, substitution
of In into Cd sites in Cd2_~In=SnO4 is also
expected to produce n-type doping and indeed this
has been demonstrated experimentally in thin film
material [21]. However, In in Sn sites will function
as a compensating acceptor. The present paper
reports the application of EELS, XPS and UPS to
study changes in electronic structure of CdzSnO 4
produced by In doping. This in turn allows us to
appraise themerits of Sb and In as n-type dopants.
The analysis of EELS data shows that carrier
concentrations introduced by indium doping are
always lower than those expected from nominal
dopant concentrations. With XPS it is found that
Y. Dou, R.G. Egdell/Surface Science 372 (1997)289-299
the atomic percentages of cadmium and tin on the
surface are lower than those found experimentally
in undoped Cd2SnO4, suggesting that when doped
into Cd2SnO4 the In 3+ ions are distributed between
Cd 2+ and Sn4+ cation sites. In a + is a donor in the
former, whereas it is an acceptor in the latter. It is
also been demonstrated that there is a dramatic
change in the surface segregation of indium after
annealing under UHV conditions. Finally
In-doped Cd2SnO 4 is compared with Sb-doped
Cd2SnO4 and with In-doped CdO.
2. Experimental
Polycrystalline samples of C d 2 S n O 4 containing
0.5, 1.0, 1,5 and 2.0 at% In were prepared by firing
a well-blended mixture of the required quantities
of Cdl_xInxO (0<x<0.02) and SnO2 at 1050°C
for 6 h in air in a recrystallised alumina boat. The
product was slowly cooled to room temperature.
The samples had a yellow colour with a green
shade that deepened slightly as the dopant concentration increased from 0.5 to 2.0 at%. The phase
compositions of the products were monitored by
X-ray powder diffraction using a Philips movingarm diffractometer. The samples were all single
phase with an orthorhombic structure [ 14]. Fig. 2
shows the variation of the orthorhombic unit cell
?
E
o
176.8
!
Q
Tq)
176.4
E
0
>
176.0
l)
O
._~
175,6
0
O
175.2
o
0.0
~I
I
I
I
0.5
1.0
1.5
2.0
bulk In content /
2.5
otom
Fig. 2. Orthorhombic unit cell volumes of indium-doped
Cd2SnO4 as a function of bulk doping level.
291
volume with doping level in indium-doped
An increase in theunit cell volume with
x indicates substitutional incorporation of the In
dopant into the C d 2 a n O 4 host lattice. The
Cdl_xInxO was itself synthesized by heating a
thoroughly blended mixture of stoichiometric
amounts of CdO and In203 at 880°C for 7 days
in air. Phase pure polycrystalline material with the
rocksalt structure of CdO was produced for doping
levels up to 2.3 at%, beyond which small peaks
due to CdIn204 were observed.
The samples were pressed into 10 mm diameter
pellets at 5.5 MPa, and then sintered in air at
900°C for 8 h. XPS and UPS were measured in an
ESCALAB spectrometer with facilities for excitation of photoelectron spectra using unmonochromated A1 Ka or MgK~ X-rays or UV radiation
from a noble-gas discharge lamp. Samples were
mounted on Pt stubs and held in position with Pt
wires. Sample cleaning was achieved by annealing
pellets for 3 h at around 400°C(stub temperature)
in the spectrometer preparation chamber (base
pressure<10-9mbar) with the aid of a watercooled copper workcoil coupled to a 400kHz
radiofrequency generator. After transfer to the
main chamber (base pressure 2 x 10-lo mbar) XPS
were found to be completely free ,of signals due to
carbon or other contamination. The O ls core
peak showed a single component with no evidence
of high binding energy shoulders due to water or
hydroxide contamination. Similarly, He(II) photoemission spectra contained no adsorbate-related
peaks on the high binding energy side, of the main
O 2p valence band, although a strong Cd shallow
core level peak was evident in all spectra. The
nominal analyser resolution~ was set at 100 meV
for He(I) photoemission measurements, and at
400meV for He(II) photoemission and XPS.
Structure due to HeIfi and HeI], lines in He(I)
spectra, and HeIIfi and HeIIy~lines in He(II)
spectra was removed from spectra using an interactive stripping routine which ensured optimal
removal of satellite intensity subject m the constraint that there should be no negative dips in,the
stripped spectra. The position of the Fermi energy
in the spectrometer was established from measurements on a sample of clean polycrystalline Ag foil.
Fermi levels of degenerate conducting stannate
C d 2 a n O 4.
292
Y. Dou, R.G. Egdell/ Surfaee Science 372 (1997) 289-299
samples should equalise with those of the Ag and
indeed a weak Fermi-Dirac-like onset was found
for In-doped Cd2SnO4 samples at the same position as the Ag onset. Binding energies are referred
to this Fermi-Dirac onset. X-ray photoelectron
spectra were stripped of satellites due to A1K~x3,4
radiation again using in-house software. Energyloss spectra were measured in the same spectrometer using an incident electron beam derived from
an electron gun usually used for low-energy
electron-diffraction studies. The analyser resolution
was set at the low value of 20 meV for EELS
measurements, mainly to protect the channeltron
from high incident electron fluxes in the elastic
peak. The experimental resolution in EELS was,
however, limited by the thermal spread of electrons
from the electron gun. Typically, the elastic peak
displayed the asymmetry to high kinetic energy
(negative energy loss) expected for electrons thermionically emitted from a hot filament with a fullwidth at half-maximum height of less than
400 meV.
~ m
C
I I I I
0
3. Results and discussion
3.1. Electron energy loss spectroscopy
Electron energy loss spectra were recorded using
a 200 eV exciting electron source. Fig. 3 shows
spectra for samples containing 0.5, 1.0 and 2.0 at%
In taken after UHV annealing. The strongest peak
in each spectrum corresponds to electrons scattered
elastically, with no energy loss. The rather intense
peak next to the elastic peak is the surface plasmon
peak which results from a longitudinal surface
excitation of the conduction electron gas and provides valuable information about the concentration
of conduction electrons close to the surface. The
shoulder to the right of the plasmon peak, which
is clearly visible in the expanded scale, is owing to
multiple loss processes. It can be seen from Fig. 3
that the plasmon peak moves tohigher energy loss
as the doping level increases. Measured plasmon
loss energies for the complete series of doped
samples are given in Table 1. Clearly, after reaching
0.42 eV at a bulk content of 1.0 at% In, EELS
shows only small subsequent increases in plasmon
1
2
loss energy /
,.3
4
eV
Fig. 3. EELS spectra of (A) 0.5, (B) 1.0 and (C) 2.0 at%
In-doped Cd2SnO 4 excited with a 200 eV electron beam after
annealing under UHV conditions. Note the resolved plasmon
loss peak in each case.
energy with bulk In content. This is owing both
to the fact that the carrier concentration falls
increasingly below the nominal In doping level as
the latter increases, and to the increasing effective
mass with increasing occupation of the conduction
band (see below).
In order to study the effect of annealing under
UHV conditions on the plasmon energy we also
recorded the EELS of 2 at% In-doped Cd2SnO 4
before annealing. Fig. 4 compares loss spectra
before and after annealing. It is seen that the
plasmon peak shifts as a result of UHV annealing.
The plasmon peak before annealing appears as a
shoulder and is estimated to be at a loss energy of
around 0.33_+0.05 eV. After UHV annealing the
loss energy increases to 0.43 eV.
At the low beam energy of the present experiments, surface loss will predominate over bulk loss.
293
Y. Dou, R.G. Egdell/Surfaee Science 372 (1997) 289-299
Table 1
Surface plasmon frequencies, carrier concentrations and effective mass in indium-doped Cd2SnO4
In content
(at%)
Plasmon frequency
(eV)
Carrier conc.
(10 TM electrons/cm3)
Doping level
(1018 atoms/cm a)
Effective mass ratio
(m*/me)
0.5
1.0
1.5
2.0
0.375
0.420
0.430
0.450
31
51
56
68
40
79
119
158
0.068
0.084
0.088
0.092
!
out by Nozik [8] and Miyata et al. [23] e(oe)
was found to be 4.0___0.1. Empirically we have
found that six values of m* at different carrier
concentrations tabulated in the earlier work are
best described in the terms of a quadratic variation
with the Fermi wavevector k F, i.e. m* varies linearly
with n 2/a :
!
A
m* = m* + an 2/3,
-
I
I
I
!
0
I
2
3
loss energy
/
eV
Fig. 4. EELS spectra of 2 at% In-doped Cd2SnO4 (A) before
and (B) after annealing under UHV conditions.
The surface excitation condition is given by
R e [ e ( c o ) ] = - i instead of Re[e(co)]=0 [22] for
the bulk excitation condition. The surface plasmon
frequency, o)sp is given by:
o~2p = ne2/[eo {[e(oo) + 1}re*I,
(1)
where n is the concentration of conduction
electrons, e the electronic charge, eo the permittivity
of free space, e(oe) the high-frequency dielectric
constant and m* the effective mass of conduction
electrons at the Fermi level. From the studies of
optical properties of C d 2 S n O 4 thin films carried
(2)
where m* = 0.02765 me (me being the rest mass) and
the constant a is such that a/me = 0.41 x 10-14 cm 2.
Combining Eqs. (1) and (2) it is possible to use
the measured values of the surface plasmon loss
energy to define the carrier concentration. The
values obtained from the analysis of energy loss
spectra are given in Table 1 and compared with
the carrier concentration derived from the bulk In
content in the initial mixtures.
It can be seen from Table 1 that the carrier
concentrations determined from EELS are always
lower than bulk doping levels. This could be caused
by the compensation of In a + substitution for C d 2 +
by cation vacancies. A more plausible explanation
is that In a + ions are substitutionally incorporated
into both C d 2+ and S n 4+ sites. Each In a+ contributes an electron to the conduction band on the
former whereas it accepts an electron on the latter.
Substitution of In 3 + ions for both C d 2 + and S n 4 +
cations produces a material with a chemical formula C d 2 _ x I n x + y S n l _ y O 4. When x > y
there
should be net n-type doping, but with a lower
carrier concentration than that when only C d 2+
sites are doped as in Cd2_xInxSnO 4. Preparation
of samples from C d l , x I n x O a n d S I l O 2 w a s
intended to avoid the problem of anti-site substitution and replacement of Sn by In possibly implies
the presence of extraneous phases below the limit
of X-ray detectability.
294
]7. Dou, R.G. Egdell/ Surface Science 372 (1997) 289-299
We are unable to establish the m o d e of substitution of In 3 + from consideration of the increase of
crystal unit cell volume. This is because replacing
either Cd 2+ or Sn 4+ ions with In 3+ could cause
an expansion of the cell. For the former it is owing
to the occupation of antibonding conduction band
states with the free electrons introduced by chemical doping [20], and for the latter it is the consequence of the bigger size of In 3÷ (r =0.94 ,~) relative
to Sn 4+ (r=0.83 A) [24]. However, the evidence
from XPS favours occupation of both Cd and Sn
sites, as discussed below.
3.21 X-rayphotoemission spectroscopy
A1 K s excited X-ray photoemission spectra in
the region of Cd 3d, In 3d and Sn 3d core levels
for In-doped Cd2SnO4 containing 0.5, 1.0 and
2.0 a t % In before annealing are shown in Fig. 5.
i
i
i
i
i
A
Cd:3ds/2
Sn:3ds/2
i Cd:3d3/2
.
Clearly the intensity of In 3d peaks increases as
the bulk In content increases, but is much higher
than that expected from the nominal In concentration, indicating that there is a pronounced segregation of In to the surface. It is also noticeable
that the surface In concentration reaches a high
value after initial doping, but then shows only a
very slight increase as the dopant concentration
further increases. Fig. 6 shows the corresponding
spectra after U H V annealing. Clearly the In 3d
intensity is much reduced after the annealing
process.
To quantify the cation distributions probed by
XPS, the areas of Cd 3d5/2 In 3d~/2 and Sn 3d5/2
peaks were divided by atomic sensitivity factors.
The resulting cation percentages are plotted in
Fig. 7. For undoped material one expects n ( C d ) =
66.7% and n(Sn)= 33.3%: for air-annealed, as pres e n t e d material we found n ( C d ) = 6 6 . 0 % and
Cd:3d6/2
Sn:3ds/2 A
I Cd:3d:3/:z
• Sn:3d3/2
ISn:3d~/2
B
B
I
390
I
I
420 450
I
I
480 510
binding energy / eV
Fig. 5. A1 Kc~ X-ray photoemission spectra of (A) 0.5, (B) 1.0
and (C) 2.0 at% In-doped Cd2SnO 4 in the Cd 3d, In 3d and
Sn 3d region before the vacuum annealing. Structure due to
satellite radiation has been subtracted from the spectra.
I
390
I
|
420 450
I
480
I
510
binding energy / eV
Fig. 6. AI K~ X-ray photoemission spectra of (A) 0.5, (B) 1.0
and (C) 2.0 at% In-doped Cd2SnO 4 in the Cd 3d, In 3d and
Sn 3d region after vacuum annealing. Structure due to satellite
radiation has been subtracted from the spectra.
Y. Dou, R.G. Egdell/Surface Science 372 (1997)289-299
80
A
nil- In"
I~i Cd[
70 . . . . . . . . . . . . . T____m____s_,_. . . . . . . . . . _]
60
0
40 . . . . . . . . . . . .
30
~ 20
o
~
"~
"o
"~"
"-
80
70
60
50
B
~
40
30
20
10
00.0
0.5
1.0
1.5
2.0
bulk in content / • atom
2.5
Fig. 7. Atomicpercentages of indium,tin and cadmiumon the
surface, derived from XPS, for In-doped Cd2SnO4 (A) before
and (B) after annealing in vacuum. The data for undoped
Cd2SnO4 are representedby dashedlines in (A).
n(Sn)=34.0% [21], in good agreement with the
"ideal" value. For the 0.5 at% In-doped sample
both n(Cd) and n(Sn) are lower than for undoped
sample, suggesting that In occupies both Cd and
Sn sites although substitution into Cd sites predominates. Additionally, it is clear that there is a
dramatic segregation of In to the surface, with
n(In)~10%, a factor of 30 higher than expected
from the bulk doping level. The segregation of In
in this system is highly reminiscent of segregation
of Sb in doped SnO 2 [4] and of Sn in In20 a E5].
In these systems the high concentration of surface
dopant has been attributed to occupation of cation
sites in the topmost ionic layer by dopant ions in
the (N--2) oxidation state, N here being the group
oxidation state. The (N--2) ions carry an electron
lone pair whose energy is lowered by the sp
hybridisation allowed a t surface sites where there
295
are non-centro symmetric electric fields. In this
model the high concentration of surface dopant
atoms is not accompanied by a high concentration
of conduction electrons because the ( N - 2 ) ions
are not donor centres. This situation certainly
pertains in the present system because as we have
seen the surface carrier concentration probed by
EELS is actually less eventhan the nominal doping
level. Thus, the high surface concentration of In
found in the present work is attributed to In + in
surface Cd sites.
Progressive In doping beyond 0.5 at% leads to
only small subsequent increases in the surface In
concentration suggesting that even at the lowest
dopant level the surface is effectively saturated
with a monolayer of In + .
The effects of UHV annealing revealed in Fig. 6
and quantified in Fig. 7 are highly surprising: the
In concentration drops dramatically as compared
with the as-presented samples. This behaviour contrasts with that found in In-doped CdO [ 6 ] where
the surface In concentration increases after UHV
annealing. The decrease in the In concentration in
the present study is accompanied by a small but
significant decrease in the Cd: Sn ratio. This behaviour is also found in undoped CdzSnO4 and reflects
the volatility of CdO. After UHV annealing the
outmost ionic layer is therefore depleted of Cd,
leaving more Sn cation sites on the surface. As
with SnO2 itself, there will be surface reduction
from Sn4+ to Snz + after UHV annealing, However,
the 5s-5p hybrid state associated with Snz+ must
be at lower energy than the corresponding In +
state. Thus there is no thermodynamic driving
force for accommodation of In + in surface Sn2+
sites and the In intensity in XPS drops
dramatically.
3.3. Ultraviolet photoemission spectroscopy
Ultraviolet photoelectron spectra excited with
He(I)~ (hv =21.22 eV) and He(II)0~ (hv=40~81 eV)
radiation were recorded. Fig. 8 shows He(I) spectra
for 1 and 2 at% In doped Cd2SnO4 taken after
annealing under UHV conditions compared with
the spectrum for undoped material [20]. Fig. 9
gives the He(II)spectrum of 2 at% In-doped
Cd2SnO 4 after UHV annealing. The He(I) spectra
296
Y. Dou, R. G. Egdell/ Surface Science 372 (1997) 289-299
i
i
,
i
,
I
i
i
i
i
i
Cd:4d
I
-10
-5
I
0
5
I
I
10
t5
20
binding energy/eV
02p
c I( C s
!
-I0
-5
Fig. 9. He(II) photoemission spectra of 2.0 at% In-doped
Cd2SnO 4 after annealing under UHV conditions Structure due
to O 2p valence electrons and shallow core level Cd 4d
electrons are labelled. Binding energies are given relative to the
Fermi energy of a silver sample stub. Structure due to satellite
radiation has been subtracted from tlie spectrum.
I
I
I
I
I
0
5
10
15
20
25
binding energy / eV
Fig. 8. He(I) photoemission spectra of (A) undoped, (B) 1.0
and (C)2.0 at% In-doped CdzSnO4 taken after UHV annealing.
Structure due to O 2p valence electrons, shallow core level
Cd 4d electrons and secondary electron emission (S) is labelled
in (C). Binding energies are given relative to the Fermi energy
of a silver sample stub. Structure due to satellite radiation has
been subtracted from the spectra.
are in each case dominated by the oxygen O 2p
valence band, the onset of which varies from about
2.80 to 3.22 eV below the Fermi energy as the In
doping level increases. The Cd 4d shallow core
level peak is at about 11.0 eV and is superimposed
on the secondary electron background which is
beyond 15,0 eV. N o major effects of indium doping
on the valence band structure can be noted from
the H e ( I ) spectra. The H e ( I I ) spectra show a better
definition of the Cd 4d core level peak since the
secondary electron tail onset is at higher energy
than in the H e ( I ) spectrum. The binding energy of
the Cd 4d level is 11.7 eV relative to EF or 8.5 eV
relative to the valence band edge. The overall O
2p bandwidth is found to be about 8.5 eV.
A very weak peak close to the Fermi energy in
the H e ( I ) spectra can b e observed on an expanded
scale. This peak arises from occupancy of the
conduction band. The conduction band must be
constructed [18] from both Sn 5s and Cd 5s
atomic orbitals which have very low one-electron
ionisation cross-sections in comparison with the O
2p states at h v = 2 1 . 2 e V (a(Cd 5 s ) = 3 . 2 5 x 1 0 -2
Mb; a(Sn 5s)=3.25 x 10 -2 Mb; o-(O 2p)=2.67 M b
[26]). Coupled with the relatively low doping
levels, this accounts for the low intensity of the
c o n d u c t i o n band structure. Nevertheless, after
removing the Satellites resulting from H e i r (hv =
23.09 eV) and HeI~ (hv =23.75 e V ) r a d i a t i o n from
raw spectra the conduction band structure can be
observed with adequate signal-to-noise to define
the conduction band shape and width. It can be
seen from Fig. 8 that the intensity of the conduction
band is higher for doped samples than for the
undoped sample. In t h e latter, the carriers are
produced by oxygen vacancies. The ratios of the
m a x i m u m intensities of the conduction bands to
those of t h e valence band are given in Table 2
together with that for undoped sample. The values
for the doped samples increase slowly with In
doping, as expected from a free electron model
297
Y. Dou, R. G. Egdell/ Surface Science 372 (1997)289-299
Table 2
The intensity ratios between conduction and valence bands
and the maxima of O 2p valence bands relative to the Fermi
energy EF in He(I) UPS for In-doped Cd2SnO4; the data for
undoped Cd2SnO4 [20] are also listed for comparison
In content
(at%)
Intensity ratio
between conduction and
valence band
maxima (x 104)
Maximum O
2p valence
band relative to EF
(eV)
0.0
0.5
1.0
1.5
2.0
4.6
8.4
8.7
8.9
9.5
4.30
4.54
4.58
4.67
4.76
where the Fermi level intensity should vary with
(cartier concentration) ~/3. This idealised behaviour
is, however, modified by the effective mass variation described in Eq. (2).
An interesting feature in H e ( I ) U P S is that the
m a x i m u m of the O 2p valence band moves gradually towards the higher b i n d i n g energy as the
bulk doping level increases. This parallels a shift
of the low binding energy edge of the valence band
to higher binding energy. The m a x i m u m of the O
2p valence band measured from the H e ( I ) spectra
for each of the doped samples is also listed in
Table 2, together with the data for u n d o p e d sample
for comparison.
The shift is analogous to the so called M o s s Burnstein shift in optical spectroscopy, where it is
found that a shift of the fundamental optical
absorption edge towards shorter wavelengths
(higher energy) accompanies increasing carrier
density. This p h e n o m e n o n has been observed with
thin films of Cd2SnO 4 [8,12,27], C d O [28], In203
[29,30] and In203-doped SnO2 [31] and is discussed in detail in Refs. [2,13,32].
Fig. 10 shows the narrow scan of H e ( I ) spectra
of In-doped and undoped Cd2SnO 4. The low binding energy edge of the valence band is at a binding
energy of 2.8 eV for undoped CdESnO 4 and 3.2 eV
for 2 a t % In-doped Cd2SnO4. The shift of the edge
of the valence band is 0.4eV for the doped sample
relative to the u n d o p e d one. The shift m a y be
attributed to an increase ~ the Fermi energy within
the conduction band resulting from the higher
cartier concentration. Unfortunately, we can not
/
AZ
B
x250
I
-2
-1
I
I
I
I
I
0
1
2
3
4
binding energy / eV
5
Fig. 10. He(I) photoemission spectra of the narrow scan of
(A) undoped and (B) 2 at% In-doped CdzSnO4 after UHV
annealing. The Fermi edge and the lower edge of the valence
band are indicated by a dashed and black line, respectively,in
each case. Binding energies are given relative to the Fermi
energy of a silver sample stub. Structure due to satellite
radiation has been subtracted from the spectra.
compare the two conduction band widths directly
because of the difficulty in locating the exact
position of the high binding energy edge of the
conduction band for the undoped sample, as seen
in Fig. 10. This is owing to the relatively lower
carrier concentration and a large number of bandgap states which obscure the high binding energy
edge of the conduction band. However, we can
make a reasonable estimate of the expected conduction band widths using a free-electron model
[20,33,34]. In this model the Fermi energy, EF,
defined relative to the b o t t o m of the conduction
band, equals the conduction band width and can
be derived from the effective mass, m*, and the
carrier concentration, n, through the following
equation:
EF = (h2/8rczm *)(3zc2n)2/3,
(3)
298
Y. Dou, ~ G. Egdell/Surface Science 372 (1997) 289-299
where h is Planck's constant. For undoped
Cd2SnO4, n and m* can be obtained by assuming
that the plasmon energy is not more than 0.3 eV
[-20]. The conduction band width determined in
this way is 0.39 eV for undoped CdzSnO 4 and
0.64 eV for 2 at% In-doped Cd2SnO4, giving an
approximate shift of the valence band edge of
0.25 eV. The discrepancy between this estimate and
the observed shift of 0.4 eV is probably due to the
uncertainty of the plasmon energy and the over
simplified model for the conduction band. Actually,
the conduction band shape does not conform
exactly to a simple expression such as that above
owing to non-parabolicity and the variation of m*
within the band.
interstitial Cd. It thus seems probable that the
change in plasmon energy for In-doped Cd2SnO4
again reflects the distribution of In between Cd
and Sn sites. Substitution of In into Sn sites is
basically producing p-type doping (which compensates n-type substitution for Cd), but p-type doping
of oxides requires highly oxidising conditions and
therefore annealing in UHV should favour a migration of In to Cd sites.
To summarise, In-doped Cd2SnO4 is a complex
system which shows qualitatively different behaviour to both Sb-doped Cd2SnO 4 and In-doped
CdO.
References
4. Concluding remarks
The carrier concentration determined from
EELS for In-doped CdzSnO 4 is always lower than
bulk doping level. This situation is different from
that in Sb-doped Cd2SnO 4 [-20], where it is found
that the carrier concentration is always slightly
greater than the bulk dopant concentration. The
result from our own study indicates [6] that no
significant compensation of electrons by Cd z+
cation vacancies occurs in indium-doped CdO.
This leads us to consider that In 3+ ions are substituted for both Cd 2+ and Sn4+ ions. XPS demonstrates that the atomic percentages of both
cadmium and tin are lower in In-doped Cd2SnO 4
than in undoped Cd2SnO 4. This gives the evidence
for the above consideration, although the data
relate to an as-presented sample.
From Fig. 4 and the discussion in Section 3.1 it
is known that high temperature vacuum annealing
increases the plasmon energy of 2.0 at% In-doped
Cd2SnO 4 by 0.12 eV, which is consistent with an
increase in the carrier concentration of 5.0 x 1019
atoms crn -3. These carriers are unlikely to be
introduced by oxygen vacancies produced at high
temperature under vacuum because EELS data for
Sb-doped CdzSnO4 show no difference between
plasmon energies for samples before and after
annealing. In fact in In-doped CdO the surface
plasmon energies show a small decrease as a result
of UHV annealing, apparently due to loss of donor
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