University of Twente
University for technical and
social studies
Oxygen vacancy ordering in iron-doped
calcium titanate
Michael van den Bossche
M. Sc. Thesis
Faculty of Science and Technology
Inorganic Materials Science
Enschede, October 2005
Oxygen vacancy ordering in iron-doped
calcium titanate
M. Sc. Thesis
By
Michael van den Bossche
Enschede, October 2005
Graduation committee
Prof. Dr. ing. D.H.A. Blank (chairman)
Dr. H.J.M. Bouwmeester
Dr. S. McIntosh (supervisor)
Dr. B.A. Boukamp
Dr. J. Vente (ECN)
Dr. K. Seshan
Always read stuff that will make you
look good if you die in the middle of it.
P.J. O’Rourke
Oxygen vacancy ordering in iron-doped calcium titanate
Summary
Iron doped calcium titanates were doped with different cations to research and
prevent the (partial) ordering of oxygen vacancies in these materials. Samples
were doped with barium on the A site of the perovskite lattice and with
chromium, nickel and cobalt on the B site. Nickel could not be incorporated into
the perovskite lattice. For the other compositions, the formation of oxygen
vacancies was recorded as a function of temperature and oxygen partial pressure.
A relatively small amount of vacancies was created in the compositions
CaTi0.6Fe0.4O3-δ and Ba0.2Ca0.8Ti0.6Fe0.4O3-δ, when heating to 1273 K in nitrogen
(pO2 = 5·10-6). These compositions contained a lot of partially ordered domains.
In the chromium-doped materials, CaTi0.6Fe0.3Cr0.1O3-δ and CaTi0.5Fe0.4Cr0.1O3-δ,
and for the composition CaTi0.8Fe0.2O3-δ, a large number of oxygen vacancies was
created in the same temperature and oxygen partial pressure range.
Oxygen permeation through membranes made from different compositions was
tested at high temperature. The highest flux of oxygen was measured for a onemillimetre-thick membrane with composition CaTi0.6Fe0.3Cr0.1O3-δ at 1223 K, with
air on the high pO2 side and hydrogen/steam on the low pO2 side. It was
concluded that the presence of Cr3+ ions had a beneficial effect on the oxygen flux
through the membrane.
Oxygen vacancy ordering in iron-doped calcium titanate
Oxygen vacancy ordering in iron-doped calcium titanate
Contents
1
2
3
4
5
6
7
8
9
Introduction .........................................................................................6
1.1
Perovskite materials ......................................................................6
1.2
Calcium titanate ...........................................................................7
1.3
Objectives ...................................................................................7
Theoretical background ..........................................................................8
2.1
Defect chemistry of ABO3 perovskites ..............................................8
2.2
Oxygen transport in ABO3 perovskites..............................................9
2.3
Vacancy ordering ........................................................................ 10
2.3.1 Fully ordered structures............................................................ 10
2.3.2 Partially ordered structures ....................................................... 11
2.3.3 Mechanism for partial ordering .................................................. 11
2.3.4 Structure and ordering of CaTiFeO3 prepared in air ....................... 12
2.3.5 Effect of ordering on oxygen permeation ..................................... 13
2.3.6 Prevention of ordering .............................................................. 13
2.4
Mössbauer Spectroscopy .............................................................. 14
2.4.1 Principle of 57Fe Mössbauer absorption spectroscopy ..................... 14
2.4.2 Data analysis .......................................................................... 16
Experimental ...................................................................................... 18
3.1
Sample preparation and composition ............................................. 18
3.2
X-Ray Diffraction measurements ................................................... 19
3.3
Density measurements ................................................................ 20
3.4
DTA/TGA experiments ................................................................. 20
3.5
TPD/TPR measurements............................................................... 20
3.6
Reduction experiments ................................................................ 22
3.7
Oxygen permeation measurements ............................................... 23
Results .............................................................................................. 25
4.1
X-Ray Diffraction ........................................................................ 25
4.2
Reductions ................................................................................. 29
4.3
Oxygen content experiments ........................................................ 30
4.3.1 Oxygen content in nitrogen gas ................................................. 31
4.3.2 Oxygen content in hydrogen atmosphere .................................... 34
4.4
Oxygen permeation experiments ................................................... 36
Discussion.......................................................................................... 39
5.1
Sample purity............................................................................. 39
5.2
Ordering of oxygen vacancies in XRD patterns ................................ 40
5.3
Oxygen content measurements..................................................... 41
5.4
Oxygen permeation results........................................................... 45
Conclusion ......................................................................................... 48
Dankwoord......................................................................................... 49
Literature........................................................................................... 50
Appendices ........................................................................................ 52
Oxygen vacancy ordering in iron-doped calcium titanate
1 Introduction
Calcium titanate or perovskite (CaTiO3) is the archetype of the group of
perovskite materials. The material was discovered by geologist Gustav Rose, who
named it after the Russian mineralogist von Perovski.
1.1 Perovskite materials
The ideal cubic perovskite structure is shown in
Figure 1-1 for CaTiO3. The general formula of the
perovskite unit cell is ABX3, where A and B are
cations and X is an anion, often oxygen. The A
site is typically occupied by an alkali earth metal
or a transition metal. This position has a
coordination number of twelve (cuboctahedral
coordination). The B site can be occupied by rare
earth metals, transition metals and some other
elements, e.g. Al, Pb and Bi. The B ion is located
in the centre of an octahedron of oxygen ions (six
fold coordination) [1].
Figure 1-1: CaTiO3
lattice cell.
The ideal cubic structure that is depicted in Figure 1-1 can be distorted to an
orthorhombic or a tetragonal structure by tilting of the octahedra. This distortion
can be due to changes in temperature or oxygen partial pressure, and can also be
due to a size mismatch of the A and B cations. From the ionic radii of ions present
on the A and B sites, the Goldschmidt tolerance factor t [2] can be calculated:
t=
rA + rO
2 ⋅ ( rB + rO )
[1.1]
ri is the ionic radius of the ions on the A, B and oxygen (X) sites. When more than
one type of ion is present on a site, the average radius is used. The value of the
ionic radii is taken from Shannon [3], and depends on the atomic element, the
coordination number, the electrical charge, and for some ions, like cobalt and
iron, on the electron configuration [4]: high spin or low spin.
When the tolerance factor is close to unity, the crystal symmetry is predicted to
be cubic. t Values in the range of 0.74 - 1 indicate a high likeliness that the
crystal is distorted.
Perovskite materials are doped on the A and B sites with other cations in order to
tailor the material properties. Through doping, oxygen vacancies can be created
in the lattice, allowing for mixed ionic and electronic conduction (MIEC) in the
perovskite. Mixed conductors find application as cathode material in solid oxide
fuel cells (SOFC) and in membranes for oxygen separation and for chemical
production.
6
Oxygen vacancy ordering in iron-doped calcium titanate
1.2 Calcium titanate
When compared to commonly used MIEC’s such as LSCF or BSCF, calcium
titanate (CaTiO3) has the advantage that it is thermally and chemically very
stable. Whereas BSCF and LSCF decompose at oxygen partial pressures of pO2 =
10-5, calcium titanate remains intact at high temperatures and in atmospheres
with extremely low oxygen partial pressure, e.g. in hydrogen. Because of its
stability, calcium titanate is used as a membrane for selective oxygen transport in
partial oxidations of hydrocarbons, which occur at a pO2 of 10-20 or even lower.
Ionic and electronic conduction in calcium titanate are low compared to LSCF and
BSCF, but can be increased by replacing part of the titanium ions with iron. There
is an optimum in the amount of iron doping, however, since for large amounts of
iron, > 0.2 mole, oxygen vacancies form ordered structures, which decrease the
oxygen ionic conductivity and hence the transport of oxygen through the
material.
1.3 Objectives
The goal of this research is to enhance the oxygen permeation and ionic
conductivity of iron-doped calcium titanate by decreasing the oxygen vacancy
ordering in iron-doped calcium titanate. In order to decrease the amount of
ordering, iron-doped calcium titanate will be doped with additional cations like
barium, chromium, cobalt and nickel.
7
Oxygen vacancy ordering in iron-doped calcium titanate
2 Theoretical background
2.1 Defect chemistry of ABO3 perovskites
The perovskite crystal structure described in the introduction is an ideal structure;
a real crystal contains defects. Every lattice will contain a number of intrinsic
defects, in the case of ABO3 perovskites these are mainly oxygen vacancies, due
to an equilibrium between lattice oxygen and gas phase oxygen, in Kröger-Vink
notation:
2OO ←⎯
→ O2 ( g ) + 2VOii + 4e'
[2.1]
Where OO is oxygen on an oxygen lattice site, O2 (g) is gas phase oxygen,
an oxygen vacancy, and e’ is an electron in the conduction band.
VOii is
At equilibrium, the chemical potentials (µ) of the partaking species are related to
each other:
⎛
⎞
⎜ ∑ν i µi ⎟ = µO2 + 2 µVOii + 4µ e' − 2µ OOx = 0
⎝ i
⎠ p ,T
[2.2]
This can also be written as a mass law with equilibrium constant K:
∑ν µ
i
0
i
= −kbT ln K
[2.3]
i
From these equations it can be concluded that the oxygen deficiency of a pure
perovskite material is dependent on at least three factors: temperature, total
pressure and oxygen partial pressure.
Pure calcium titanate is a poor ionic conductor, since it contains few intrinsic
oxygen vacancies. More, extrinsic vacancies can be introduced by doping the
calcium titanate on the B-site, replacing the Ti4+ with lower valency ions. For
example iron. Iron is present as Fe3+ and Fe4+ when the sample is prepared in air,
but it reduces completely to Fe3+ in at high temperatures and in a low pO2
atmosphere. The substitution of iron for titanium in the calcium titanate lattice
can be described with equation 2.4:
1
2TiTi + Fe2O3 + O2 → 2 FeTi + 2TiO2
2
[2.4]
At elevated temperatures and low oxygen partial pressures, iron is reduced and
oxygen vacancies are created:
2 FeTiX ↔ 2 FeTi' + VOii + O( g )
[2.5]
According to this equation, one mole of iron atoms introduced in the material is
theoretically able to generate half a mole of oxygen vacancies.
8
Oxygen vacancy ordering in iron-doped calcium titanate
2.2 Oxygen transport in ABO3 perovskites
The defect chemistry of ABO3 perovskites makes selective transport of oxygen
through the bulk material possible, provided that a gradient in oxygen chemical
potential is present. Oxygen from the gas phase adsorbs on the surface of the
perovskite material, dissociates into oxygen ions and is subsequently incorporated
into the lattice. When sufficient oxygen vacancies are present in the material, a
percolation pathway is available for the incorporated oxygen: it can hop from one
vacancy to another, from the side with high to the side with low oxygen chemical
potential. At the low pO2 side, the ionic oxygen is oxidized and, after association,
can desorb from the surface as oxygen gas.
Purely ionic conductors, such as 8-mol% yttria-stabilised zirconia (8YSZ) or
gadolinia-doped ceria (GDC), require an external circuit to conduct electrons
produced from the dissociation of oxygen gas.
Perovskites can be good electronic conductors and
transport of oxygen can take place without any
external short-circuiting. Materials that conduct ions as
well as electrons are called mixed ionic electronic
conductors (MIEC). In the materials of interest here,
oxygen ions move through the bulk material and
simultaneously, electrons move in the opposite
direction by means of a small polaron hopping
Figure 2-1: ambipolar
mechanism. The overall transport mechanism is called
diffusion.
ambipolar diffusion (Figure 2-1).
At equilibrium, oxygen and electron fluxes are coupled by the law of
electroneutrality. This means that the amount of oxygen diffusing through the
perovskite does not solely depend on the ionic conductivity and the oxygen
chemical potential gradient, but also on the electronic conductivity. The oxygen
flux cannot be described using a simple diffusion equation, but is described using
the Wagner equation:
µ
2
σ O 2− σ e
1
jO2 = − 2 2 ∫
d µO2
4 F L µ " σ O 2− + σ e
O'
[2.6]
O2
Where jO2 is the flux of oxygen through the perovskite, F is faraday’s constant, L
is the distance over which the oxygen is transported, σ i is the conductivity of
particles i and µO2’ and µO2” are the chemical potentials of oxygen at the high and
low pO2 side respectively. The Wagner equation is derived from the equations for
oxygen and electron flux described by Onsager, the law of electroneutrality and
the equation for chemical equilibrium (2.3). A complete derivation is given by
Wagner [7].
According to equation 2.6 oxygen flux depends on the driving force (the
difference in partial pressure), the dimensions of the material (length L) and the
type of material (ionic and electric conductivities). Conductivities depend on the
mobility and the number of conducting particles. In most perovskites, at standard
oxygen pressures, electrons and oxygen vacancies are present in about equal
amounts, but the electrons are much more mobile, making electronic conductivity
the predominant part of the total conductivity.
9
Oxygen vacancy ordering in iron-doped calcium titanate
2.3 Oxygen vacancy ordering
The introduction of Fe2O3 in calcium titanate influences the crystal structure. This
influence can be seen in the CaTiO3–CaFeO2.5 phase diagram (Figure 2-2). The
diagram was constructed from XRD measurements in reducing atmosphere, all
iron is thus considered to be Fe3+ [8, 9].
Figure 2-2: phase diagram for the (reduced) CaTiO3-CaFeO2.5 system
[8]
.
The crystal structure of pure CaTiO3 is orthorhombic Pbnm (circles in Figure 2-2)
at room temperature, but changes to tetragonal I4/mcm (diamonds) and to cubic
Pm-3m (squares) at higher temperatures. Calcium ferrate has the brownmillerite
Pcmn structure. When calcium titanate is doped with an increasing amount of
iron, the structure changes from orthorhombic to tetragonal to cubic. For larger
amounts of iron, the doped calcium titanate partially and eventually fully orders
2.3.1 Fully ordered structures
For CaTi1-xFexO3-δ with x > 0.5, the material is fully ordered. This means that
throughout the whole material, the crystal structure of the material has changed.
All fully ordered structures are intergrowths of CaTiO3 perovskite unit cells (Figure
2-3a) and CaFeO2.5 brownmillerite cells (Figure 2-3b). The perovskite cell contains
only octahedral layers (abbrev. O), brownmillerite contains alternating layers of
octahedrons and tetrahedrons (abbrev. T) [10, 11]. The fully ordered structure of x
= 0.5 (Figure 2-3c) can ideally be described as TOOO [12], indicating that for
every layer with tetrahedral coordinated iron atoms, three layers with octahedral
coordination are present, two that contain titanium and one that contains iron
ions. Ideally, because the stacking in the material is likely to be TOOOOTOO at
times. The unit cell is orthorhombic, the length of one of the three axes being
very large, 30.22 Å compared to 5.4 Å for the other two axes. The structure of Fe
=0.66 is ideally a stacking of TOO sequences [13]. The materials in between the
mentioned compositions can be described as a mixture between different ideal
stacking orders.
10
Oxygen vacancy ordering in iron-doped calcium titanate
Figure 2-3: Intergrowth (c) of octahedral (a) and tetrahedral (b) layers
[14]
.
2.3.2 Partially ordered structures
From XRD-data [15], the partially ordered
structures
can
be
interpreted
as
intermediaries between disordered CaTiO3
and fully ordered structures (Figure 2-4).
For an increasing amount of iron (and
decreasing annealing temperature) the
pm-3m cubic perovskite structure (2-4a)
has it peaks broadened and eventually split
up, while new peaks appear (2-4b,c,d) [15].
The result is a completely ordered
structure (2-4e).
Figure 2-4: XRD diagram of iron doped CaTiO3.
(a) Is in the disordered region, (b), (c) and (d)
are partially ordered and contain increasing
amounts of iron. (e) Is totally ordered, with an
iron content of x=0.5 [15].
2.3.3 Mechanism for partial ordering
McCammon et al. have proposed a mechanism for the partial ordering of oxygen
vacancies in iron doped CaTiO3 [16], under reducing atmosphere. The mechanism
is illustrated in Figure 2-5. Starting with pure CaTiO3, Ti4+ ions are substituted by
Fe3+, creating one oxygen vacancy for every two substituted titanium ions, to
maintain electroneutrality. When this vacancy is located near an iron ion, this ion
will have fivefold instead of six fold coordination (2-5, left). For low amounts of
iron, the iron atoms and the oxygen vacancies are placed randomly in the lattice,
creating only pentacoordinated positions in addition to the hexacoordinated ions.
11
Oxygen vacancy ordering in iron-doped calcium titanate
Figure 2-5: Short range ordering of oxygen vacancies
[16]
.
For higher iron contents, two vacancies will be situated together (dimer), and one
of the iron ions will be tetra coordinated (2-5, middle). When more iron oxide is
added, chains of oxygen vacancies are formed. Most of the iron ions next to these
chains will be tetracoordinated. The oxygen vacancies are now partially ordered in
chains, but they do not form a regular stacking of octahedral and tetrahedral
planes yet (2-5, right).
2.3.4 Structure and ordering of CaTiFeO3 prepared in air
Structure and ordering of iron-doped calcium titanate depend strongly on the
oxygen partial pressure of the atmosphere. When the sample is prepared in air, a
significant part of the iron ions is Fe4+ [6]. The crystal structures of CaTi1-xFexO3-δ
prepared in air differ from reduced samples [6, 17]. Pure, disordered CaTiO3 is
orthorhombic (Pnma), and it remains orthorhombic up to 0.4 moles of iron
doping. For higher amounts of iron (x = 0.6) the structure changes to a cubic
(Pm-3m) unit cell (Figure 2-6). TEM and SAED [14] measurements however show
an orthorhombic structure at this point. The cubical structure detected in XRD, is
explained by Canales-Vazques et al. [14] who claim that the orthorhombic ordering
takes place along three perpendicular axes, resulting in an overall (pseudo)cubic
structure.
Figure 2-6: XRD patterns of air synthesized CaTi1-xFexO3-δ
[6]
.
12
Oxygen vacancy ordering in iron-doped calcium titanate
Iwahara et al. [5] also measured on air-treated,
iron doped calcium titanate samples and found
that orthorhombic peaks already disappeared at
x=0.30 (Figure 2-7). At higher amounts of iron,
for 0.30 < x < 0.50, the material is cubic. For
x=0.60 iron, the material consists of a cubic
phase and an unknown secondary phase.
Figure 2-7: XRD patterns of
CaTi1-xFexO3-δ prepared in air
[5]
.
2.3.5 Effect of ordering on oxygen permeation
The (partial) ordering of oxygen vacancies in iron
doped calcium titanate has a significant detrimental
effect on the conductivity of oxygen ions and thus on
the permeation of oxygen through the material. From
literature [5, 6] it can be gathered that ionic conductivity
is at an optimum for an iron doping of x=0.2 moles.
When the calcium titanate is doped with more iron
(>0.2) the ionic conductivity decreases (Figure 2-8),
although more oxygen vacancies are formed in the
material.
The decrease in ionic conductivity is caused by the
partial ordering of the oxygen vacancies in chains and
sheets making oxygen vacancies less mobile.
Figure 2-8: Ionic
conductivity as a function
of T and iron content [5].
2.3.6 Prevention of ordering
In literature it has been shown that ordering may be prevented. One method is
by doping the perovskite on the A-site, an example is the doping of SrCo0.8Fe0.2O3
(SCFO) with barium [18]. From XRD measurements it follows that undoped SCFO
shifts from perovskite to brownmillerite structure at low oxygen partial pressures.
SCFO doped with x > 0.3 moles of barium is stable at low oxygen partial
pressures.
One common explanation for this is the variation in ion size, characterised by the
tolerance factor (equation 1.1). In low oxygen partial pressures, Fe4+ in SCFO will
reduce to Fe3+ (High spin), which is larger than Fe4+. The change in iron ion
radius can be compensated by partly substituting strontium ions with larger
barium ions, keeping the relation of ionic radii of A and B site ions constant.
Another way of preventing ordering is by doping a material on the B-site, like
LaSr2Fe3O9 doped with chromium on the iron site [19]. The structure with the
chromium is more stable, because the chromium, just as titanium, prefers
13
Oxygen vacancy ordering in iron-doped calcium titanate
octahedral coordination, instead of the iron, which is present in both octahedral
and tetrahedral coordination. Doping with chromium thus prevents the formation
of ordered tetrahedral layers. Furthermore, the chromium can still be reduced
from the +4 to the +3 oxidation state, which allows the creation of oxygen
vacancies.
2.4 Mössbauer Spectroscopy
Mössbauer spectroscopy (MS) is a spectroscopic method. This means that the
emission or the absorption of radiation by a sample is measured as a function of
radiation frequency. An emission or absorption peak is due to an energy transition
in the measured material. In the case of Mössbauer spectroscopy γ-radiation is
used and the transitions that occur are due to changes in energy levels of nuclei
in the material. MS only yields useful results for a small number of elements, Fe
and Sn being the most important. From the MS spectrum information can be
deduced about coordination and oxidation states of the element in question. In
this chapter, theory and data analysis of 57Fe MS will be explained briefly. A more
detailed description is available in literature [20, 21].
2.4.1 Principle of 57Fe Mössbauer absorption spectroscopy
Important in Mössbauer spectroscopy is the generation of radiation with a very
small energy distribution. Emission peaks have to be narrow, because the
difference in frequencies between two absorption peaks can be very small.
Radiation source
For 57Fe Mössbauer absorption experiments the source is
radioactive 57Co, often incorporated in a non-magnetic solid
possessing cubic symmetry, like rhodium, to ensure that only one
frequency is emitted by the cobalt source. The cobalt has a
transition to an excited state of 57Fe by means of electron capture
(Figure 2-9). This excited state (spin S = 5/2) decays
predominantly (89 %) to another 57Fe excited state at 14.4 keV (S
= 3/2), which decays to the ground state of 57Fe (S =1/2). The
latter transition is used for 57Fe MS and is called the Mössbauer
transition. The radiation released from this transition can be
absorbed by a similar nucleus (57Fe) in the sample. The principle
that the transition that produces the radiation is the same as the
transition that absorbs the radiation is called resonance
absorption.
Figure 2-9:
57
Co source.
Energy distribution of the radiation
The Mössbauer transition is not a discrete one; emitted photons have a
distribution of energies because of an uncertainty in the life time of the excited
state. From the uncertainty relation the natural line width (Γ) can be calculated:
Γ=
τN
[2.7]
14
Oxygen vacancy ordering in iron-doped calcium titanate
in which τN is the life time of the state emitting the photon. For the Mössbauer
level in 57Fe, Γ amounts to 4.67·10-9 eV. Free nuclei emit radiation with a broader
energy distribution because of two effects:
1. Recoil. Due to mass and momentum conservation rules, an emitted photon has
an energy equal to the energy available from the transition (ET), minus the recoil
energy of the emitter (ER, Em).
2. Thermal broadening. Random moving sources like gas molecules have broader
peaks because of the Doppler effect.
The energy distribution of the emitted photon is the total of these contributions:
E Em = E T − E R , Em + E D = E T −
pγ2
2M
−
pγ P
M
[2.8]
in which pγ is the momentum of the photon, M is the mass of the nucleus, and P is
the momentum of the nucleus in the direction of the emitted photon.
A similar equation applies for the energy of photons absorbed by nuclei but in this
case, the recoil energy Er, abs is positive. Resonance absorption only occurs when
the absorption and emission distributions overlap (hatched area):
Figure 2-10: Resonance absorption and scattering.
Mössbauer effect
Overlap normally occurs for low energy radiation (near UV and visible regions); Er
and ED are small, and the overlap is large. For higher energies, like γ-radiation, Er
and ED values are much higher and the overlap is negligible. Mössbauer found
that for nuclei present in a solid, a fraction of the emission events has effectively
zero Er and ED values. The line width of the radiation is now the natural line width.
This is called the Mössbauer effect; it implies that a source can be created that is
able to emit radiation with a very narrow energy distribution. With such a source,
one is able to accurately measure the transitions that occur in the absorber,
provided that a range of frequencies can be covered.
Measuring a frequency range
To measure a frequency range the source is accelerated to and from the sample.
This induces a Doppler effect that shifts the emission frequency to lower and
higher frequencies. The Doppler effect used here does not cause any peak
broadening, because the movement is in one direction instead of random. The
shift in energy (∆E) can be calculated from the following equation:
∆E =
v
Eγ
c
[2.9]
in which v is the velocity of the source with respect to the absorber.
15
Oxygen vacancy ordering in iron-doped calcium titanate
Absorber (sample)
MS is often performed on a powder; this can be necessary, because the
orientation of the crystals in the absorber influences the intensities of the peaks
in the absorption spectrum; in a powder these crystals are randomly orientated.
There is an optimum in the amount of powder used: more powder allows a more
significant absorption of photons; too much powder causes broadening of the
emission peak, making the measurement inaccurate.
2.4.2 Data analysis
A typical absorption spectrum is depicted in Figure 2-12. Transmittance (counts)
is plotted versus the velocity of the driver coil (in mm s-1). The velocity scale can
be recalculated to a frequency or an energy scale (equation 2.9). The spectrum
consists of a baseline, which relates to 100 % transmittance (no absorption). The
peaks in the spectrum are caused by absorption of radiation by the sample.
Figure 2-12: A typical Mössbauer spectrum.
These peaks contain information about the oxidation state and the coordination of
the iron nuclei in the sample. An isolated iron nucleus has only one degenerated
excited state at 14.4 keV; in an absorption spectrum, only one peak would
appear. However, because of the environment in which the iron nucleus is
present, the excited state has its energy levels split up. There are a number of
influences on the energy levels of the nucleus:
-
Electric fields arising from its own orbital electrons.
Magnetic fields arising from its own orbital electrons.
(A relatively small contribution from) electric fields from remote ions.
An externally applied magnetic field.
In theory, an external electric field would also influence the energy levels of the
nucleus, but it is not possible to create sufficiently strong electric fields.
The magnetic and electric fields mentioned above interact with the magnetic
dipole and the electric quadrupole moments of the nuclei. These interactions are
called hyperfine interactions. This implies that materials with no magnetic dipole
(quantum spin number = 0) and with no electric quadrupole (quantum spin
number < 1) are not interesting for MS measurements.
16
Oxygen vacancy ordering in iron-doped calcium titanate
Interactions with electric fields
The interaction of a nucleus with an electric field brings about
two changes in the spectrum. Firstly, the absorption peak
shifts from zero position (as obtained from calibration with a
reference compound, usually α-Fe foil). This shift is called the
isomer or chemical shift (δ). The second change is the splitting
up of peaks in the spectrum.
The isomer shift is illustrated in Figure 2-13 and is due to a
change in the energy level of the ground state and of the
degenerate 14.4 keV state. The magnitude of the shift
depends on the oxidation state and the coordination of the
element. This influence is shown schematically in Figure 2-14a
(oxidation state) and 2-14b (coordination number).
Figure 2-14a: Isomer shift (mm/s) as a
function of oxidation and spin number.
Figure 2-13:
Isomer shift.
Figure 2-14b: Isomer shift (mm/s) as a
function of coordination number.
The second effect of the electric interaction, the peak splitting, is due to the
excited state getting split up into two separate levels. The energetic difference
between these two levels is very small, but measurable due to the very small
peak width emitted by the source. The isomer shift for a set of split peaks is
calculated from their average shifts.
Interactions with magnetic fields
A magnetic field removes all the degeneracy of the nuclear
spin levels. If the nucleus had a spin value of I, then the
magnetic field gives rise to 2I+1 energy levels (Zeeman
effect). In the case of 57Fe, the ground state has I = ½ and
the excited state I = 3/2. Therefore, not only the excited state
is split up (in four levels) but also the ground state is split up,
in two levels. (Figure 2-15). In this figure, some of the
possible transitions are omitted. This is caused by selection
rules for magnetic dipolar emission of a photon, which require
that ∆m for the transition is 0 or ±1. (For electric quadrupole
emission, ∆m of 0, ±1 and ±2 is allowed). Therefore, the
transition from an excited state with m = 3/2 to a ground
state with m = –1/2 is not allowed, because ∆m = 2 for that
transition. The resulting MS spectrum will show six peaks, one
for each different energy transition. This is the spectrum
shown earlier in Figure 2-12.
Figure 2-15:
Energy levels for
57
Fe in a magnetic
field.
17
Oxygen vacancy ordering in iron-doped calcium titanate
3 Experimental
In this chapter, the preparation and the composition of iron-doped calcium
titanate materials are described, followed by a description of the methods used
for characterisation of these materials.
3.1 Sample preparation and composition
Sample preparation is based on literature [5]. The composition of the prepared
materials is summarized in Table 3-1.
Appropriate amounts of high purity powders (≥ 99.95%) of BaCO3, CaCO3, TiO2
(rutile), Fe2O3, NiO and Cr2O3 were weighed in a porcelain boat. Cobalt oxide was
prepared by heating Co(NO3)2⋅6 H2O (99%). The product was determined pure
Co3O4 by XRD.
The powders were ball milled in ethanol (99.98% pure) for 22 to 24 hours, using
zirconia milling balls. The weight ratio of balls to powder was 12:1. Milling balls
with a diameter of 1 mm and of 10 mm were used in a weight ratio of 5:1. After
milling, all samples had a pink-reddish colour, save the CaTi0.6Cr0.4O3-δ sample,
which was light green due to the presence of chromium oxide. Next, the milling
balls were separated from the suspension using a sieve. The suspension was
poured into a porcelain bowl heated to 343 K, to evaporate the ethanol. The
resulting powder was ground in a mortar and pestle. The powders were then
calcined in a platinum boat at 1423 – 1573 K (Table 3-1) for sixteen hours in air.
Heating and cooling rates were 2 K min-1. After calcination, all samples had a dark
brown/black colour. At this point, the calcined powder was again ground and ball
milled in ethanol for 22 to 24 hours. After evaporation of the ethanol, the powder
was ground again and sieved through a 200 µm sieve.
To form pellets, 1.25 grams of powder were uni-axially pressed in a 20 mm
diameter mould using a force of 32 kN (1000 MPa). After that, the disc was
packed into a condom (cleaned with ethanol), and isostatically pressed (4000 bar,
five minutes). The obtained discs were sintered in air for sixteen hours at 1673 –
1698 K. The barium-doped sample is an exception to this; the sample deformed
at these temperatures and was therefore sintered at 1573 K. Heating and cooling
rates were 2 K min-1. After sintering, the discs had a black colour.
Table 3-1: Sample compositions.
Sample
Composition
CTF82
CTF64
BCTF64
CTFCr631
CTFCr541
CTCr64
CTFCo631
CTFNi631
CaTi0.8Fe0.2O3-δ
CaTi0.6Fe0.4O3-δ
Ba0.2Ca0.8Ti0.6Fe0.4O3-δ
CaTi0.6Fe0.3Cr0.1O3-δ
CaTi0.5Fe0.4Cr0.1O3-δ
CaTi0.6Cr0.4O3-δ
CaTi0.6Fe0.3Co0.1O3-δ
CaTi0.6Fe0.3Ni0.1O3-δ
Calcining T
(K)
1423
1573
1523
1523
1423
1523
1423
1423
Sintering T
(K)
1698
1673
1573
1673
1673
1673
1673
1673
18
Oxygen vacancy ordering in iron-doped calcium titanate
Oxygen permeation samples
Oxygen permeation membranes were prepared as above, except for the
following. Calcination was performed at 1423 K for all powders, to keep the grain
size relatively small. After calcination, the powder was ball milled for 24 hours,
and the suspension was sieved through a 38 µm sieve. The retentive suspension
was planetary milled for one hour and again sieved through 38 µm. Ethanol was
evaporated from the permeate suspension and pellets pressed as described
above. All samples were sintered at 1573 K, except CaTi0.8Fe0.2O3-δ, which was
sintered at 1623 K, because of its high titanium content, and CaTi0.6Fe0.4O3-δ
(1673 K), since the latter was not phase pure when sintered at 1573 K. Heating
rate was 2 K min-1, cooling rate was 0.5 K min-1. Afterwards, the samples were
polished to a thickness of 1 mm and a diameter of 15 mm. Oxygen permeation
samples are summarized in Table 3-2.
Table 3-2: Sample compositions for oxygen permeation experiments.
Sample
Composition
CTF82
CTF64
CTFCr631
CTFCr541
CaTi0.8Fe0.2O3-δ
CaTi0.6Fe0.4O3-δ
CaTi0.6Fe0.3Cr0.1O3-δ
CaTi0.5Fe0.4Cr0.1O3-δ
Calcining T
(K)
1423
1423
1423
1423
Sintering T
(K)
1623
1673
1573
1573
Mössbauer samples
Mössbauer samples (Table 3-3) were prepared according to the method for the
oxygen permeation samples, the only difference being the sintering step.
Sintering occurred at 1673-1698 K, with heating and cooling rates of 2 K min-1.
Samples were not polished.
Table 3-3: Sample compositions for Mössbauer experiments.
Sample
Composition
CTF82
CTF64
CTFCr631
CTFCr541
CTFCo631
CaTi0.8Fe0.2O3-δ
CaTi0.6Fe0.4O3-δ
CaTi0.6Fe0.3Cr0.1O3-δ
CaTi0.5Fe0.4Cr0.1O3-δ
CaTi0.6Fe0.3Co0.1O3-δ
Calcining T
(K)
1423
1423
1423
1423
1423
Sintering T
(K)
1673
1673
1673
1673
1673
3.2 X-Ray Diffraction measurements
XRD patterns were collected at room temperature using CuKα radiation, step size
of 0.02° and a collecting time of 1.5 s step-1. A collecting time of 17.5 s step-1 was
used for lattice cell parameter measurements. Measurements were made with
variable receiving slit, which was calculated back to fixed slit before Rietveld
fitting of the pattern using equation 3.1:
I fixed =
I variable
sin(θ )
[3.1]
19
Oxygen vacancy ordering in iron-doped calcium titanate
HT-XRD patterns were recorded in nitrogen (100 ppm O2) and air atmosphere
from 573 K to 1473 K at every 100 K on a Philips X’Pert with CuKα radiation, step
size of 0.04° and a collecting time of 2 s step-1. Measurements were performed
with fixed receiving slit.
Structures were fitted using the Rietveld method as implemented by the software
GSAS [22] to obtain lattice cell parameters and XRD-densities for the different
compositions.
3.3 Density measurements
Absolute density of sintered samples was measured by immersion in mercury and
calculated using equation 3.2:
ρ abs =
mair
m
=
V ( mHg / ρ Hg )
[3.2]
in which mair and mHg are the masses of the sample measured in air and in
mercury respectively; ρHg is the density of mercury, being 13.53 g cm-3 [23]. The
theoretical density was obtained from the Rietveld refinement of the XRD data.
Relative density was calculated by dividing the absolute density with the XRD
density.
3.4 DTA/TGA experiments
DTA/TGA measurements were performed using a Setaram Setsys 16/18. Sintered
discs were ground and then annealed for ten hours at 1273 K, cooled at 0.5 K
min-1 in air. The annealing step and slow cooling ensured that samples were at
equilibrium with the oxygen content in the air up to the temperature during
cooling at which kinetics became too slow. Samples were heated in nitrogen gas
to 1573 K at a rate of 10 K min-1. Dwell was two hours.
3.5 TPD/TPR measurements
Temperature Programmed Desorption (TPD) and Temperature Programmed
Reduction (TPR) measurements were used to measure the change in oxygen
content of the samples as a function of temperature and in different atmospheres.
Experimental
The sample, ground and annealed, is put into a quartz glass tube where it is held
in place by quartz glass wool. Different gas mixtures can be flowed through the
tube using Mass Flow Controllers (MFC’s). The oxygen content of the ingoing gas
mixture is checked with an oxygen sensor. During the experiment, the tube is
heated in a furnace and the composition of the gas flow leaving the tube is
analysed using a mass spectrometer. The setup is shown in Figure 3-1.
20
Oxygen vacancy ordering in iron-doped calcium titanate
Figure 3-1: TPD/TPR setup.
For the TPD measurements, 100 ml min-1 of nitrogen was led over 0.5 grams of
sample. Samples were heated at a rate of 10 K min-1 to a temperature of 1273 K;
after ten minutes at this temperature, the sample was cooled down. The mass
spectrometer measures oxygen released by the sample.
TPR measurements were recorded with a flow of 5 to 6 ml min-1 of hydrogen gas
and 35 ml min-1 of argon gas. During the measurement, hydrogen reacts at the
surface of the sample, reducing the metal ions and consuming oxygen to form
water. The mass spectrometer measures the decrease of hydrogen, from which
the change in oxygen content of the samples is calculated. The temperature
program used is the same as for the TPD measurements, except for the dwell:
samples stayed at 1273 K for 90 minutes.
Data analysis
TPD and TPR data can be interpreted qualitatively as well as quantitatively.
Qualitatively, the peak position gives information at what temperature oxygen is
released and metal ions are reduced. Quantitatively, TPD and TPR data can be
used to calculate the change in oxygen content.
In the case of TPD measurements the mass spectrometer records the amounts of
N2 and O2 as an ionic current. When it is assumed that the current is linearly
dependant on the partial pressures of nitrogen and oxygen, the following equation
applies:
pO2 ,released =
nO2 ,released
nO2 + nN2
=
I O2 − I O2 ,baseline
I O2 + I N 2
[3.3]
in which I O2 and I N 2 are the ionic currents measured by the mass spectrometer.
I O2 ,baseline Is the ionic current due to oxygen already present in the nitrogen sweep
gas. The volumetric amount of oxygen can be calculated for a small time interval
when the flow rate φV , N 2 is known:
VO2 (∆t ) = pO2 ⋅ φV , N2 +O2 ⋅ (ti − ti −1 ) = pO2 ⋅ φV , N2 ⋅ (ti − ti −1 )
[3.4]
21
Oxygen vacancy ordering in iron-doped calcium titanate
In the second part of [3.4], the flow rate of oxygen
φV ,O
2
is neglected compared to
the nitrogen flow. To obtain the amount of oxygen in moles the volume of oxygen
θ
is divided by the molar volume Vm . The different small time intervals are summed
up:
nO2 ( ti − t0 ) =
φV , N
I O2 (ti )
t = ti
2
Vmθ
∑I
t = t0
O2
(ti ) + I N2 (ti )
⋅ ( ti − ti −1 )
[3.5]
In this equation nO2 (ti – t0) is the oxygen released in moles for the time interval
between i and 0;
φV , N
2
is the volumetric flow of nitrogen gas in m3 s-1.
In the case of TPR data it is important to have a clear baseline. This baseline
presents 100% of the hydrogen flow rate. During the heating, hydrogen is
consumed in the reaction with oxygen. The H2 ionic current measured at that
point ( I H 2 , meas ) will be lower than that of the baseline. Assuming that the ionic
current is linearly dependant on the amount of H2 present, the flow of H2 that was
consumed in the reaction with oxygen can be calculated:
φV , H
2 , consumed
In which
⎛
I H 2 , meas ⎞
= φV , H 2 ⋅ ⎜ 1 −
⎜ I H ,baseline ⎟⎟
⎝
⎠
2
φV , H
2
[3.6]
is the volumetric flow of hydrogen set by the controller (in m3 s-1).
The amount of hydrogen consumed can be calculated for a certain time interval
(or temperature interval) using the following equation, which is comparable to
equation 3.4:
nH 2 ( t i − t 0 ) =
φV , H
θ
Vm
2
t =ti
I H 2 ,meas (ti )
t = t0
H 2 ,baseline
∑I
(ti )
⋅ ( ti − ti −1 )
[3.7]
The amount of oxygen gas is half the amount of the consumed hydrogen gas,
because one mol of oxygen reacts with two moles of hydrogen.
3.6 Reduction experiments
Sintered discs were ground to a powder and then annealed at 1273 K for ten
hours. Cooling rate was 0.5 K min-1. 0.5 grams of powder were accurately
weighed in a small platinum crucible and conveyed into an alumina-silica tube.
The tube was put into a small furnace and closed airtight with steel flanges. 30 ml
min-1 of 50% H2 in argon was fed through the tube and the furnace heated at 10
K min-1 to 1273 K and cooled to room temperature after twelve to fourteen hours.
The platinum boat was removed from the tube and its mass recorded. The
difference in weight before and after reduction is due to the decrease of oxygen
content of the sample.
22
Oxygen vacancy ordering in iron-doped calcium titanate
3.7 Oxygen permeation measurements
The flux of oxygen through the perovskite samples was measured using oxygen
permeation measurements. The theory behind this transport has already been
explained in the theoretical background chapter.
The polished membranes were measured for diameter, thickness and density.
Afterwards they were cleaned with ethanol and put into the permeation reactor,
which is shown schematically in Figure 3-2. Samples were sealed to the reactor
by means of a glass ring. The quartz glass reactor was heated in a furnace (0.5 K
min-1) to a temperature of 1313 K at which point the glass ring would seal the
membrane. After twelve minutes the reactor was cooled to 1223 K.
Figure 3-2: Oxygen permeation setup. Adapted from
[1]
.
During the measurements the retentate side of the membrane was exposed to air
(φV = 48.9 ml min-1); this is the high pO2 side. The permeate side of the
membrane was exposed to either argon or a mixture of argon, hydrogen and
water; this is the low pO2 side. The gas at the permeate side was analysed before
and after the reactor by means of a Varian CP 4900 micro gas chromatograph. A
cold water trap in front of the GC prevented flooding of the column.
Measurements with argon on the low pO2 side were performed at 1223 K as a
function of pO2 by measuring at different flow rates. Oxygen flux was calculated
from the concentration of oxygen in the permeate gas:
( )
J O2 PO"2 =
φV ⋅ [O2 ]
106 ⋅ AEff
[3.8]
In which JO2 is the flux of oxygen through the membrane in ml min-1 cm-2, φV is
the volumetric flow of argon gas in ml min-1, [O2] is the oxygen concentration in
the permeate gas leaving the reactor in ppm, AEff is the area of the membrane
that is effectively used for oxygen transport and is calculated from the dimensions
of the membrane and the glass seal.
23
Oxygen vacancy ordering in iron-doped calcium titanate
Measurements with a mixture of argon, hydrogen and water were performed by
flowing 20 ml min-1 of gas (50% Ar, 50% H2) through a bubbler. Permeation was
measured at different temperatures: 1223, 1173, 1123 and 1073 K. Oxygen flux
was calculated from the difference in hydrogen concentration in the permeate gas
entering and the gas leaving the reactor:
J O2 (T ) =
(
0.5 ⋅ φVtot ⋅ [ H 2 ]before − [ H 2 ]after
10 ⋅ AEff
6
)
[3.9]
In which φVtot is the total flow rate on the permeate gas side, i.e. the flow rate of
argon and hydrogen gas in ml min-1, [H2]before and [H2]after are the concentrations
of hydrogen in the permeate gas before and after the reactor, in ppm.
24
Oxygen vacancy ordering in iron-doped calcium titanate
4 Results
4.1 X-Ray Diffraction
Figure 4-1 shows XRD patterns for all prepared compositions recorded at room
temperature. The diffraction pattern of BCTF64 shows the typical reflections of a
cubic perovskite unit cell (space group Pm-3m), however the expected reflections
at 2Θ = 23º, 53º and 74º are missing. This is the result of the influence of
barium doping on the structure factor (FHKL) of the unit cell. The same peaks are
also missing in the Rietveld fitting of the BCTF64 cell. Closer observation of the
BCTF64 pattern reveals additional peaks, belonging to an unknown phase. The
XRD pattern for CTF64 can be fitted on a Pm-3m cubic perovskite unit cell, and is
phase pure. The peaks are very broad, especially at the base, most pronounced
at 2Θ = 48º and 60º. The CTF82 pattern shows primarily cubic symmetry as the
CTF64 composition, however additional peaks belonging to an orthorhombic
perovskite unit cell, space group Pnma are present. This is the same space group
as pure CaTiO3 at these temperatures. CTCr64 has the same structure as CTF82.
Figure 4-1a and b: RT-XRD patterns of the prepared materials.
The composition CTFNi631, shown in Figure 4-1b, shows a set of cubic perovskite
peaks, which are very broad from top to base, and an additional set of peaks
belonging to nickel(II)oxide (indicated with arrows). The cobalt-doped material
does not show an additional phase, but the cubic perovskite peaks are very
25
Oxygen vacancy ordering in iron-doped calcium titanate
broad. The samples that contain iron as well as chromium, CTFCr631 and
CTFCr541, can be indexed on a cubic perovskite unit cell with space group Pm3m. Peaks are relatively narrow for these compositions.
Phase attribution
Phases were attributed using the XRD patterns shown in Figure 4-1a and b. The
results are summarized in Table 4-1. Lattice cell parameters were refined using
the highest purity samples, made for Mössbauer experiments. These results were
compared with the lattice cell parameter XRD measurements that were performed
on the BCTF64, CTF64 and CTFCr631 compositions made for oxygen content
measurements. These latter samples contained some additional peaks which were
neglected when calculating the lattice cell parameters. From the cell geometry,
theoretical density was calculated.
Table 4-1: Phase attribution and cell parameters for CTF XRD patterns.
Sample
Phase(s)
CTF82
Cubic pm-3m (87%),
Orthorhombic pnma (13%)
Cubic pm-3m
Cubic pm-3m
Cubic pm-3m
Cubic pm-3m
Orthorhombic pnma
Cubic pm-3m
Cubic pm-3m,
NiO phase
CTF64
BCTF64
CTFCr631
CTFCr541
CTCr64
CTFCo631
CTFNi631
Lattice cell
parameters (Å)
a
3.845
5.456
3.840
3.870
3.840
3.843
5.404
3.85
3.83
-
b
3.845
7.676
3.840
3.870
3.840
3.843
7.604
3.85
3.83
-
c
3.845
5.414
3.840
3.870
3.840
3.843
5.365
3.85
3.83
-
XRD density
(g cm-3)
3.98
3.99
4.42
4.03
4.01
4.16
3.98
-
HT-XRD patterns
High temperature XRD patterns were recorded for CTF64, BCTF64 and CTFCr631.
Patterns were recorded in air and in nitrogen as a function of temperature
between 298 and 1473 K and are shown in Figure 4-2a to f. Platinum peaks due
to the sample holder are indicated with an arrow.
Figure 4-2a shows the diffraction patterns of CTF64, recorded in air, for different
temperatures. No change of phase was detected from room temperature to 1473
K. The same is true for the same sample recorded in nitrogen (Figure 4-2b).
Figure 4-2c shows the XRD patterns recorded for CTFCr631 in air. At high
temperatures (1473 K), the cubic perovskite peaks are broadening. An effect
which is seen best when observing reflections at 2Θ = 69° and 79°. The same
occurs to a larger extent in the XRD patterns recorded in N2 (Figure 4-2d).
BCTF64 also shows peak broadening at higher temperatures, in air (Figure 4-2e)
as well as in nitrogen atmosphere (Figure 4-2f).
26
Oxygen vacancy ordering in iron-doped calcium titanate
27
Oxygen vacancy ordering in iron-doped calcium titanate
Figure 4-2a to f: HT-XRD patterns for: a. CTF64 recorded in air; b. CTF64 in nitrogen; c.
CTFCr631 in air; d. CTFCr631 in nitrogen; e. BCTF64 in air; f. BCTF64 in nitrogen. Platinum
peaks are marked with arrows.
Rietveld [22] refinements were performed on the HT-XRD patterns using the GSAS
package to yield the lattice cell parameter of the unit cell as a function of
temperature. Plots of lattice cell parameters versus temperature are shown in
Figure 4-3a and b. The first shows the expansion of the cubic cell for CTF64; in
the temperature range of 298 to 1473 K, the cell parameter increases from 3.824
Å to 3.889 Å in N2 and to 3.881 Å in air. Figure 4-3b shows the expansion of
CTFCr631. The lattice expands from 3.835 Å to 3.878 Å in air and to 3.882 Å in
N2 .
28
Oxygen vacancy ordering in iron-doped calcium titanate
Figure 4-3: Lattice cell parameter as a function of temperature in air and N2 for a. CTF64,
and b. CTFCr631.
From the lattice cell parameters at different temperatures, the (total) expansion
coefficients were calculated for the temperature range 298 – 1473 K (Table 4-2).
Table 4-2: Expansion coefficients for CTF and CTFCr.
Sample
CTF64
CTF64
CTFCr631
CTFCr63
Atmosphere
Air
N2
Air
N2
Expansion Coeff
(10-6 K-1)
11.7
13.2
9.95
10.4
4.2 Reductions
All materials were reduced in hydrogen to calculate the absolute oxygen content
of the samples. Crucial in this calculation is that the products of the reduction are
known. From other perovskite compositions (the series Ba1-xSrxCo1-yFey) a trend
in the reducibility of metal ions is deduced. For ions on the B-site, the trend is,
from low reducibility (high oxidation state) to high reducibility (low oxidation
state): Ti < Cr < Fe < Co < Ni. Titanium in CTF materials is expected to remain
Ti4+ even after reduction in hydrogen, whereas Co and Ni are expected to reduce
to their metallic form.
Reduction products of iron-containing calcium titanates
The XRD pattern for reduced CaTi0.6Fe0.4O3-δ showed a lot of additional, nonperovskite peaks. A set of peaks could be attributed to iron metal, another set to
calcium oxide. Perovskite peaks were also still present. It follows that part of the
29
Oxygen vacancy ordering in iron-doped calcium titanate
original perovskite is retained - the cells containing the stable titanium ions while another part of the perovskite is destroyed – the iron-containing part. When
it is assumed that titanium does not change oxidation state, so all oxygen
released can be attributed to the iron ions, then iron was initially present as Fe3+
with a small part Fe4+. The reduction of iron to metallic iron was checked by
reducing Fe2O3 under the same circumstances. With the breaking up of the
perovskite lattice, reduction of iron in the sample should be comparable to the
reduction of Fe2O3. The end product of the reduction was indeed iron metal.
Reduction products of chromium-containing calcium titanates
The XRD pattern of the reduced CaTi0.6Cr0.4O3-δ showed no additional peaks: the
material remained as a perovskite. The material did decrease in mass, enough to
decrease the chromium oxidation state by one. Since chromium is more stable
than iron, and since the perovskite lattice was retained, it is most likely that
chromium was reduced from Cr4+ to Cr3+. When reducing Cr2O3, no weight loss
was recorded, indicating that Cr will remain in the 3+ oxidation state in hydrogen
atmosphere below 1273 K. Reduction of Cr in the perovskite lattice might
however not be comparable to reduction of Cr2O3.
Cobalt and barium doping
Cobalt ions are even more reducible than iron ions and are likely to reduce to
metallic cobalt. This was confirmed by TPR data (4.3.2). Barium, as with calcium,
does not reduce but remains in the same oxidation state (2+).
From the reaction products and the loss of mass during the reduction, the initial
oxygen content of the samples is calculated (Table 4-3). With the exact
stoichiometry of the material known, the molecular weight was calculated.
Table 4-3: Oxygen content and molecular weight of CTF-samples.
Sample
Reaction products
O released
(mol O mol-1
CTF)
O content
(mol O)
Mol. weight
(g mol-1)
CTF82
BCTF64
CTF64
CTFCr541
CTFCr631
CTCr64
CTCo631
CaTiO3; Fe; CaO
BaCaTiO3; Fe; CaO
CaTiO3; Fe; CaO
CaTiCr(III)O3; Fe; CaO
CaTiCr(III)O3; Fe; CaO
CaTiCr(III)O3
CaTiO3; Fe; Co; CaO
0.32
0.54
0.61
0.69
0.57
0.22
0.59
2.92 ± 0.02
2.74
2.81 ± 0.02
2.84 ± 0.02
2.92 ± 0.02
3.02 ± 0.02
2.79 ± 0.01
136.20 ± 0.33
154.39
136.07 ± 0.34
136.95 ± 0.33
137.44 ± 0.38
137.91 ± 0.38
135.43 ± 0.22
4.3 Oxygen content experiments
Knowing the initial oxygen contents of the materials, changes in the absolute
oxygen content as a function of temperature and for different atmospheres can
be calculated from TGA, TPD and TPR data. The oxygen content is a direct
measure of the total amount of oxygen vacancies present in the material.
30
Oxygen vacancy ordering in iron-doped calcium titanate
4.3.1 Oxygen content in nitrogen gas
Oxygen release peaks of samples in nitrogen gas are visible in the TPD data
(Figure 4-4). The release peaks are relative in height to each other and
normalized for the moles of sample measured.
In Figure 4-4a, the top picture shows oxygen released by CTFCo631. Due to
problems with the mass spectrometer, substantial noise is present in the data,
but it is still clear that oxygen release is starting at 1025 K. Only one peak is
clearly present, but there is some deviation of the baseline around 800 K. The
second spectrum shows the release of oxygen for the barium doped sample.
Oxygen is released in one peak from 650 to 950 K. The third and the fourth
compositions, CTF64 and CTF82, release oxygen at about the same temperature
interval as the BCTF64 composition. The peak height of CTF82 is significantly
larger than the other peaks though, indicating that more vacancies are formed.
In Figure 4-4b, the first spectrum shows oxygen release for CaTi0.6Cr0.4O3.02.
Oxygen is released over a broad temperature interval, starting at 750 K. Two
additional peaks, on top of the broad release, are present at high temperature.
These are not artefacts or errors in the measurement, but are reproducible. The
second and the third spectra show TPD data for CTFCr631 and CTFCr541. The
latter is noisy for the same reason as the CTFCo631 composition. Both spectra
show two release peaks. The first release starts at 700 K, the same temperature
as the peaks for the BCTF64, CTF64 and CTF82 composition. The second peak
occurs at 1150 K. Separation of such superimposed peaks is difficult; it is not
possible to tell at what temperature the second release starts. The last spectrum
in Figure 4-4b shows the CTF64 composition again, for comparison.
Figure 4-4a and b: Oxygen release as a function of temperature as measured by TPD.
31
Oxygen vacancy ordering in iron-doped calcium titanate
Discussion of TPD data
When comparing all of the TPD results, it is concluded that the release of oxygen
at about 650 – 950 K is due to the reduction of iron, since titanium does not
change its oxidation state and no other reducible ions are present in compositions
CTF64, BCTF64 and CTF82. Part of the iron is initially present as Fe4+, as can be
derived from the oxygen content calculated from the reduction data (4.2). This
part will reduce to Fe3+ starting at 650 K. The sample containing no iron, CTCr64,
does not have a clear peak at this temperature, but it does show a slow release of
oxygen, becoming more clear at temperatures of 900 K. Since no iron is present
in this sample, this peak must be the reduction of chromium, from 4+ to 3+. The
samples containing both chromium and iron show two peaks, the first is the iron
reduction, the second chromium. The cobalt data has no release peak at 650 –
950 K, but the deviation of the baseline mentioned earlier could be due to oxygen
released at 650 – 950 K. The release at 1025 K can be ascribed to the transition
Co3+ -> Co2+.
TGA experiments
The change in mass of the same samples under the same circumstances is
measured with TGA. From this data, oxygen content is calculated as a function of
temperature for all samples. In Figure 4-5a, the oxygen content of CTF82 is
shown as a function of temperature. The initial content was calculated from the
reduction experiments to be 2.92. Between 700 and 950 K this decreases to 2.89.
Theoretically, if all of the iron is in 3+ oxidation state, the oxygen content would
be 2.90. The next composition is CTF64, with an initial oxygen content of 2.81,
which decreases to 2.80 in the temperature range of 700 – 900 K. 2.80 is the
expected value of CTF64 if all iron is present in the 3+ oxidation state. For the
cobalt doped material, the initial oxygen content of 2.79 is reduced in two steps
to a value of 2.75. The first step takes place from 650 to 850 K, to yield an
oxygen content of 2.78. The second step starts at 1000 K.
In Figure 4-5b the change in oxygen content of the samples containing chromium
is shown. The CTCr64 composition shows a steady release of oxygen starting at
800 K, decreasing the oxygen content from 3.02 to 2.90. For all chromium in the
3+ oxidation state, the oxygen content would be 2.80. However, as shown in the
figure, the release is not complete at 1550 K, and failed to stabilize after two
hours at 1573 K. Interesting to note is that around 1000 – 1100 K, there is some
variation in the smooth decline. This variation occurs at the same temperature as
the two peaks in the TPD data, proving yet again that this is not an artefact. The
last two pictures belong to the CTFCr631 and the CTFCr541 compositions. Both
samples contain iron as well as chromium ions and have two reduction steps
which overlap. The point where one step ends and the other begins was
estimated by finding the bending point in the plot. The oxygen content of
CTFCr631 decreases from 2.92 to 2.88 due to iron reduction, and to 2.83 due to
chromium reduction. CTFCr541 decreases from 2.84 to 2.82 to finally stabilize at
2.79. For both these compositions, the final content does not reach the expected
value: CTFCr631 would stabilize at an oxygen content of 2.80 instead of 2.83 and
CTFCr541 would reach 2.75 instead of 2.79 if all of the chromium and iron were
in the +3 oxidation state.
32
Oxygen vacancy ordering in iron-doped calcium titanate
Figure 4-5a and b: Oxygen content of the CTF-samples, calculated from TGA data.
Together with the TGA, DTA experiments were run on the samples (Figure 4-6):
Figure 4-6a and b: DTA results.
33
Oxygen vacancy ordering in iron-doped calcium titanate
All samples in Figure 4-6 show endothermic, which is expected, since the
materials are heated up. Apart from this steady endothermic process, most
compositions do not show peaks in energy uptake or release; nothing that could
indicate total ordering in the sample. In Figure 4-6a, the compositions BCTF64
and CTFCo631 show some additional peaks probably due to a reaction taking
place: BCTF64 was not completely phase pure as determined by XRD, and
indications of Co leaching from the cobalt doped sample were present in the
platinum cup after the experiments.
4.3.2 Oxygen content in hydrogen atmosphere
With TPR experiments, the change in oxygen content in a more reducing
atmosphere was measured. Hydrogen gas was used to reduce the materials. All
samples should yield the same products as obtained with the reduction
measurements, performed in paragraph 4.2. Hydrogen consumption peaks are
shown in Figure 4-7 for all samples measured.
Figure 4-7: TPR measurements.
The first composition, CTF82, shows hydrogen consumption at 600 K and 1250 K.
The second composition, CTF64, only shows one peak at about 1250 K. No
hydrogen consumption is seen at about 600 K. The next two compositions,
CTFCr541 and CTFCr631 both show two peaks, at similar temperatures. The first
34
Oxygen vacancy ordering in iron-doped calcium titanate
peak occurs at about 900 K, the second peak at 1250 K. CTCr64 shows one clear
peak at about 900 K, and also shows a decrease in hydrogen signal at about 600
K, which could be the release of additional oxygen (oxygen content above 3.00).
The cobalt-doped material shows one peak at 600 K and a double peak at high
temperatures. The single peak at 600 K is much larger than the peaks observed
in the other compositions at this temperature.
Discussion of TPR data
Compositions containing only iron and titanium (CTF82, CTF64) show large
hydrogen consumption at 1250 K. The CTF82 sample also shows a small peak at
600 K. Since titanium is stable in hydrogen at these temperatures, iron is
reduced. The small peak in the CTF82 sample is the reduction of Fe4+ to 3+. The
peak at 1250 K is Fe3+ reducing to Fe, since the end product is known from
reduction experiments to be Fe. This large decrease in oxidation state also
explains for the large peak size. The latter peak is also visible in other samples
containing iron. The composition containing both chromium and titanium
(CTCr64) shows one reduction peak. Logically, this is the reduction of Cr4+ to
Cr3+. This reduction peak is also present in the samples containing both iron and
chromium (CTFCr631 and CTFCr541). The composition doped with cobalt shows
one single peak at 600 K and a double peak at 1200 K. The first is explained by
the reduction of cobalt from 3+ to 2+. The double peak is partly the reduction of
Fe3+ and partly the reduction of Co2+ to Fe and Co.
From the hydrogen consumption peaks, the change in oxygen content was
calculated as a function of temperature (Figure 4-8a and b).
Figure 4-8a and b: Change in oxygen content of CTF samples in hydrogen as a function of
temperature, calculated from TPR data.
35
Oxygen vacancy ordering in iron-doped calcium titanate
In Figure 4-8a, the topmost line represents the reduction of CTF82. For this
material, the oxygen content decreases from 2.92 to 2.87 (Fe4+ to Fe3+) and
decreases further to 2.61, at which point all iron is present as iron metal. The
data is quite well in agreement with theory; the oxygen content of CTF82 with all
iron in 3+ oxidation state is calculated to be 2.90 (measured as 2.87) and the
content of CTF82 with all iron present as metallic iron is 2.60 in theory. Note that
at this point, there is no single component with structure CTF82, in fact there is a
perovskite material with composition CaTiO3 and CaO and Fe present. The oxygen
content of CTF64 is reduced from 2.81 to 2.13 (theoretical value is 2.20 for
reduction to metallic iron). The composition containing cobalt reduces in two
steps: in the first step, all cobalt is reduced to 2+, which would theoretically bring
the oxygen content to 2.75. In the TPR data, enough hydrogen is consumed at
this point to decrease the oxygen content to 2.66. The second reduction step is
even further off: instead of the theoretical oxygen content of 2.20 (all iron and
cobalt present as metal), the oxygen content is calculated to be 1.72.
In Figure 4-8b, the reduction of chromium containing samples is shown. The first
composition, CTCr64, reduces from O = 3.02 to 2.80, which is the theoretical
value for CTCr64 with all chromium in 3+ oxidation state. The other two
compositions are also reduced to almost the theoretical values for Fe3+ and Cr3+:
2.41 (theory: 2.40) for CTFCr631 and 2.13 (2.20) for CTFCr541. The biggest
difference for both compositions is in the first reduction step. CTFCr631 reduces
to 2.81 in oxygen content (which would theoretically be 2.85) and CTFCr541
reduces to 2.71 instead of 2.80.
4.4 Oxygen permeation experiments
Oxygen fluxes through the prepared membranes (thickness 1.0 mm, diameter
15.0 mm) were measured at 1173 using different flow rates of argon at the low
pO2 side and are shown in Figure 4-9. The flow rate of the argon gas sets the pO2
at the low pO2 side (pO2”). The partial oxygen pressure at the high pO2 side is
kept constant at 0.21 atm (air). The flux of oxygen is plotted as a function of the
pO2” for membranes with compositions CTF82, CTF64 and CTFCr631.
36
Oxygen vacancy ordering in iron-doped calcium titanate
Figure 4-9: Flux of oxygen through CTF membranes; Argon on low pO2 side.
Oxygen permeation through the same materials was measured as a function of
temperature using a mixture of argon, hydrogen and water on the low pO2 side.
The flow and the composition of the mixture on the low pO2 side were held
constant. The actual pO2 (at the outlet of the reactor) differed due to the
consumption of H2 and the formation of H2O by oxygen permeating through the
membrane. Partial oxygen pressures at the outlet are estimated (Appendix A2) to
vary between log pO2” = -17 to -20 (Figure 4-10).
Figure 4-10: Flux of oxygen through CTF membranes; H2/H2O on low pO2 side.
37
Oxygen vacancy ordering in iron-doped calcium titanate
After the oxygen permeation experiments in hydrogen, samples were cooled
down to room temperature and XRD patterns were recorded on the samples to
identify phase changes in the samples. On the surface of the low pO2 side, the
perovskite cells containing iron had decomposed into iron metal and calcium
oxide. The surface on the high pO2 side had not changed phase.
38
Oxygen vacancy ordering in iron-doped calcium titanate
5 Discussion
5.1 Sample purity
XRD patterns measured for the prepared samples showed that most of the
compositions were phase pure. Samples prepared for Mössbauer experiments
were phase pure or consisted of a mixture of two perovskite phases in the case of
CTF82. This two phase mixture was also found by Figueiredo et al. [6] and
Waerenborgh et al. [17] for CTF82 sintered in air at 1600 K, who found that
orthorhombic peaks diminished with increasing iron content, disappearing
completely for the composition CTF46 which XRD pattern can be indexed on a
cubic perovskite cell. The diffraction pattern of CTF82 was indexed by Figueiredo
[6]
on an orthorhombic Pnma space group. Fitting the structure of CTF82 prepared
in this study on a two-phase mixture of a Pnma cell and a cubic Pm-3m cell yields
a good fit of the XRD data.
The diffraction pattern measured on the CTF64 composition was not in complete
agreement with the same composition measured by Figueiredo et al. [6] and
Waerenborgh et al. [17]. Their sample is still partly orthorhombic according to the
XRD pattern; the material used in this study was pure cubic perovskite.
Orthorhombic peaks did, however, disappear in the CTF46 composition made by
Figueiredo [6]. An explanation for this discrepancy could be that due to the shorter
sintering time used by Figueiredo [6], two hours instead of sixteen. It is suggested
that after short sintering times the resulting perovskite is not homogeneous, but
contains small domains with CaTiO3 orthorhombic symmetry. Another explanation
could be that Figueiredo [6] used a longer collecting time when measuring XRD
patterns, increasing sensitivity to the smaller peaks. This could not be checked
because no XRD measurement parameters were mentioned. However, this
explanation is unlikely because the orthorhombic peaks measured in the CTF64
sample by Figueiredo [6] are as large as in the CTF82 sample; since the CTF82
sample prepared in this study did show orthorhombic peaks, these should also be
visible in the CTF64 sample if there was an orthorhombic phase present. Our
results are supported by the findings of Iwahara et al. [5], who found that for
samples prepared in air, a CTF sample containing 0.25 moles of iron still
contained an orthorhombic phase, but that the CTF73 sample, containing 0.30
moles of iron, was completely cubic.
Samples prepared for oxygen permeation were also phase pure, except for the
composition CaTi0.6Fe0.4O3-δ, sintered at 1573 K, which contained some additional
peaks, the most pronounced occurring at 2θ = 37˚, which is the main reflection
of Fe2O3 hematite phase, indicating that not all of the iron oxide had reacted. This
composition was subsequently sintered at higher temperature, 1673 K, yielding a
phase pure material. Samples for oxygen content experiments were prepared
according to the general preparation method (paragraph 3.1). CTF64, BCTF64
and CTFCr631 compositions were made using powder calcined at 1523 – 1573 K.
Calcining at such high temperature yields large grains, which reduce homogeneity
and reactivity of the powder; the result was that CTF64 and BCTF64 samples
contained some additional peaks. This was most significant for the CTF64
material. After sintering the material a second time, the additional peaks
disappeared, but the broad shape of the peaks remained. TGA experiments were
performed on the CTF64 material sintered once, and on the same material after a
second sintering step. No significant differences were measured, so it is
concluded that the additional phase in CTF64 and probably also the extra phase in
BCTF64 did not influence the outcome of the oxygen content measurements.
39
Oxygen vacancy ordering in iron-doped calcium titanate
The nickel-doped sample contained a significant amount of NiO, indicating that at
least part of the nickel is not incorporated in the perovskite lattice. Incorporation
of nickel in a perovskite usually occurs at high temperatures, but after sintering
at 1743 K, the nickel oxide phase was still present. No experiments were
performed on this sample because of this large NiO phase. The cubic perovskite
peaks present in the XRD pattern could be indexed on a cubic perovskite unit cell
with lattice parameter a = 3.83 Å, but the broad peaks made accurate Rietveld
fitting difficult. The broad peaks are thought to be caused by a distortion of the
cubic perovskite lattice because of cation deficiency on the B site. Since no
calcium oxide was found in the XRD pattern, it is deduced that all calcium oxide
has formed a cubic perovskite, albeit with substantial B-cation deficiency. Since
no molecular weight could be calculated for the nickel-doped material, no
theoretical density was calculated for this composition.
The diffraction pattern of the cobalt-doped sample resembled the pattern of the
nickel-doped sample. Both had very broad cubic perovskite peaks. This
resemblance makes it questionable whether cobalt is completely incorporated in
the lattice. At room temperature, cobalt would be present as Co3O4, but no traces
of this phase are visible in the XRD pattern; however part of the cobalt can have
leached out of the sample, as was the case during the TGA measurement, leaving
the perovskite deficient for cobalt, which would account for the broad peaks.
Furthermore, the sensitivity of XRD is not sufficient to discern small quantities,
less than ~ 3 wt%, of Co3O4 which may be present in the material.
The problem with incorporating nickel and cobalt is mentioned in literature by
Iwahara et al. [5] who tried substituting large amounts of titanium in CaTiO3 with
nickel and cobalt. The group concluded that 0.1 was the maximum amount of
titanium that could be replaced by nickel or cobalt. No comments about peak
shapes are provided.
5.2 Ordering of oxygen vacancies in XRD patterns
No peaks were present in the diffraction patterns that would prove the existence
of a totally ordered structure of iron doped calcium titanate, since this ordered
structure would be present as a different phase in the XRD patterns, as
mentioned in the theoretical background (chapter 2); this was also expected from
the phase diagram provided by McCammon et al. and shown in [16], who showed
that total ordering does not take place for less than 0.5 moles of iron doping.
The broad bases of reflections in the CTF64 composition are a remarkable feature
of this spectrum, and are explained by McCammon et al. [16] to be due to the
partial ordering of oxygen vacancies in orthorhombic domains. These
orthorhombic domains are small and do not show directly in the XRD pattern, but
are the cause of strain on the cubic lattice. Canales-Vazquez et al. [14] used a
similar explanation for the disagreement between XRD data and HRTEM data on
CTF46. This composition has very broad peaks in XRD that can be indexed on a
cubic unit cell, but from HRTEM data it follows that the material is build up from
orthorhombic unit cells.
The HT-XRD measurements also did not yield any information about changes in
phase of the BCTF64, CTF64 and CTFCr631 composition. The oxygen content in
the nitrogen atmosphere during the HT-XRD measurements was probably too
high to make the HT-XRD exactly comparable with TPD and TGA data. In
addition, at higher temperatures no significant changes take place in crystal
40
Oxygen vacancy ordering in iron-doped calcium titanate
structure, from which it can be concluded that total ordering does not take place
in the lattice when iron is reduced at 650 – 950 K, or when chromium is reduced
at 1000 K. This is supported by the phase diagram published by McCammon et al.
[16]
.
5.3 Oxygen content measurements
Reduction experiments
The absolute change in oxygen content can only be calculated when the initial
value of the oxygen content is known. This initial content was measured by
reduction in hydrogen at 1273 K (paragraph 4.2). From the expected stability of
the ions present in the perovskite materials and by using XRD data on the
reduced samples, the reduction products can be deduced. Calcium and titanium
did not reduce at all, which is in agreement with reports in literature [24, 25] that
describe the stability of CaTiO3 at high temperature and under strongly reducing
atmosphere. Iron reduced to iron metal, as shown by XRD data on the reduced
perovskite. The initial amount of oxygen calculated from the reduction is high
enough for part of the iron to have an oxidation number of 4+ (in air). This was
supported by Mössbauer data found in the literature [17], from which the presence
of a small amount of Fe4+ next to mainly Fe3+ is suggested. Chromium was mainly
present as Cr4+ and reduced to Cr3+ while remaining in the perovskite lattice.
Cobalt reduced to cobalt metal, leading to the decomposition of the perovskite
phase.
The errors in the initial oxygen content as measured by reduction experiments
were calculated according to the equation derived in the appendix (equation
A1.3). These errors are shown in Table 4-3 and amounted to 0.02 moles of
oxygen for almost all samples. The error for the cobalt-doped sample was less
because of a larger amount of oxygen released from the sample. No error was
given for the barium-doped sample. This sample was reduced using a lot larger
platinum crucible than the other samples, which caused a large error in the
results, because a relatively small difference in the mass of the platinum sample
holder measured before and after the reduction could cause a big difference in
the calculated initial oxygen content of the sample. This mass change is the result
of contamination of the platinum, since platinum itself is inert in hydrogen gas at
these temperatures. As a result the barium-doped sample was measured to have
an initial oxygen content of 2.74, while it was expected to have an oxygen
content of 2.81, comparable to the CTF64 data. This would be logical, because it
has the same oxygen release peak in the TPD data (650 – 950 K), associated with
Fe4+ reduction. However, with an initial oxygen content of 2.74, Fe would have to
be present as a mixture of Fe3+ and Fe2+, which is not the case. The large error in
the measurement accounts for this too low oxygen content.
TPD measurements
TPD measurements were used to record the temperature at which oxygen is
released by samples in nitrogen atmosphere. TPD results were reproducible. An
indication of the accuracy of the TPD measurement is derived from compositions
CTF82 and CTCr64 that were measured twice to control the reproducibility.
Table 5-1 summarises the temperatures for which the oxygen release peaks
started in these measurements.
41
Oxygen vacancy ordering in iron-doped calcium titanate
Table 5-1: Reproducibility of TPD measurements.
Composition
CTF82
CTCr64
T of onset of O2 release (K)
Difference in T
(K)
Measurement 1 Measurement 2
695
691
4
734
717
17
In the discussion of TPD data (Paragraph 4.3.1), oxygen release peaks are related
to the reduction of ions in the perovskite. In literature little proof can be found for
the attribution of these peaks, except for the reduction of Fe4+ to Fe3+, which was
measured in a TPD-like measurement by Figueiredo [6] and indeed started around
650 – 700 K, the same temperature range that was attributed to the reduction of
Fe4+ to Fe3+ in these experiments. However, the attribution of peaks is compliant
with the type of ions present in the composition (Ti, Cr and/or Fe). Measurements
of compositions with only titanium and chromium (CTCr64) and with only titanium
and iron (CTF64, CTF82) made attribution to the type of ion possible. From the
initial oxygen content, it was possible to attribute the peaks to the change in
formal oxidation state (e.g. 4+ reducing to 3+). The amount of oxygen released
in these oxidation steps also complied with the theoretical oxygen content for
complete reduction steps.
An additional remark can be made about the CTCr64 composition. The TPD
spectrum contained two additional peaks on top of the broad release peak. These
peaks were also present in the duplicate measurement of the same material, and
could be found in the TGA data. It is suggested that an impurity is not the cause
of these peaks as CTCr64 was found to be XRD phase pure and the peaks are
associated with a significant amount of oxygen release. One explanation could be
that the oxygen in the CTCr64 orthorhombic lattice is not chemically equivalent.
Two different sites for oxygen exist in the Pnma lattice [26]. Each peak could
belong to a different site. The slow release, starting at low temperatures, could be
additional oxygen in the lattice getting released. The oxygen content of the
CTCr64 composition was calculated to be 3.02 from the reduction experiments, so
additional oxygen could be present in the lattice. However, the errors calculated
for the oxygen content derived from reduction experiments amounted to 0.02;
the value of 3.02 is still within error range of 3.00.
TGA measurements
When comparing the temperature for which oxygen is released as recorded by
TPD with the temperature of the mass change measured by TGA (Table 5-2), it is
obvious that the changes in mass in the TGA data are due to the release of
oxygen. The release of oxygen and the reduction in weight occur at the same
temperature interval.
From the TGA data, the amount of oxygen released from the sample can be
calculated as moles of oxygen leaving one mole of sample. With the reduction
step known, the percentage of the ion that is reduced can be calculated. In this
way, the numbers of iron and chromium present as 4+ in the sample can be
estimated. It is calculated that for the CTF64 composition, 3 to 4% of iron is Fe4+,
while the remainder is Fe3+. The BCTF64 sample contains 5% of Fe4+ and the
CTF82 sample 22% Fe4+. These results are comparable with Mössbauer data from
literature, in which the amount of Fe4+ and Fe3+ is measured before and after
heating in argon gas (to 1273 K) [17]. Note that a very small amount of Fe4+ is
present as 4+ even after heating to 1273 K. This could be explained by the fact
42
Oxygen vacancy ordering in iron-doped calcium titanate
that Mössbauer measurements were performed on the sample when it is cooled
down to room temperature again. It could have taken up a small amount of
oxygen at room temperature, re-oxidizing some of the Fe3+ in this way. For
comparison, the initial and the remaining amount of Fe4+ in these Mössbauer
samples are shown in Table 5-2. No Mössbauer data was present in the literature
for iron in chromium doped CTF samples.
Table 5-2: Comparison of TPD and TGA data.
Sample
Reduction
step
CTF82
BCTF64
CTF64
CTFCr541
CTFCr541
CTFCr631
CTFCr631
CTCr64
CTFCo631
CTFCo631
+4->+3
+4->+3
+4->+3
+4->+3
+4->+3
+4->+3
+4->+3
+4->+3
+4->+3
+3->+2
(Fe)
(Fe)
(Fe)
(Fe)
(Cr)
(Fe)
(Cr)
(Cr)
(Fe)
(Co)
TGA data
T range
Ions
(K)
red.
(%)
707 – 1026
635 – 950
714 – 1012
764 – 1081
1081 – 1573
808 – 1081
1081 – 1569
893 – 1573
658 – 1009
1009 - 1574
22.10
5.35
3.60
11.55
63.60
28.73
98.20
62.10
4.87
62.40
TPD data
T range
Ions
(K)
red.
(%)
691 – 1095
641 – 1006
660 – 975
705 – 1008
1008 – 1273
669 – 977
977 – 1273
737 – 1273
983 - 1273
28.70
4.95
3.20
7.20
67.20
25.20
87.80
12.95
40.00
Mössb.
Ions
red.
(%)
23 -> 2
7 -> 1
-
TGA quantitative data is also compared with TPD data for all compositions. The
TPD data is in good agreement with the TGA results especially for the
compositions that only contain iron and titanium on the B site. When chromium is
also present, it is difficult to separate the iron reduction peak from the chromium
reduction peak. This is a problem for both the TPD data and for the TGA data, and
causes the TPD determined data to deviate from the TGA data. Deviation is also
caused by the temperature program of the TPD measurements; chromium
reduction was not complete at the maximum temperature, 1273 K, in nitrogen.
This is clearly visible in TPD data since the oxygen intensity does not return to
baseline level at 1273 K. Also in TGA data, at 1550 K the CTCr64 and CTFCr631
composition have not yet stabilised: chromium reduction continues. This explains
why not 100% of chromium is reduced from Cr4+ to Cr3+. This could however also
be partly due to the fact that chromium is partly present as 3+ at room
temperature (in air). This is not very probable though, because Cr4+ to Cr3+
reduction only takes place at relatively high temperatures. The actual amounts of
Fe4+ and Cr4+ that are reduced can only be estimated more accurately with
Mössbauer data on these samples.
Oxygen vacancy formation in nitrogen
In Table 5-3 the number of oxygen vacancies present in the different
compositions is shown before and after reductions of Fe, Cr and Co. The amount
of vacancies is calculated from the difference in oxygen content before and after
the reduction step is measured with TGA. For CTF64 almost no additional oxygen
vacancies are formed; most of the oxygen vacancies that are able to form
because of Fe4+ reduction are already present in the material in air. After the TGA
measurement, at 1550 K, the oxygen content of the CTF64 is 2.80 – no Fe4+ is
present, only Fe3+. From another view, one could say that the CTF64 composition
undergoed little oxidation during the annealing at 1273 K, the oxygen
stoichiometry is constant in the range of 298 K (pO2 = 0.21 atm) to 1273 K (pO2
= 5·10-6). The reason for this may be that the material is not able to pick up any
oxygen: oxygen vacancies are partially ordered in orthorhombic domains. CTF82
43
Oxygen vacancy ordering in iron-doped calcium titanate
and also the chromium doped samples are reduced, oxygen vacancies are
formed, at low temperatures (starting at 650 – 700 K) in nitrogen. This implies
that there is less partial ordering of oxygen vacancies in these materials at pO2
between air and 5·10-6 atm and below 1273 K. The shape of the peaks in XRD
data confirms this: the broad peaks in the CTF64 pattern are explained by partial
oxygen vacancy ordering in orthorhombic domains. Chromium doped CTF
samples have significantly narrower peaks.
It can be concluded from the oxygen content experiments in nitrogen that the
amount of oxygen vacancies available for the transport of oxygen shows a trend;
from high to low: CTFCr631, CTFCr541, CTF82, BCTF64 and CTF64. More
vacancies are created in the CTFCr631 than in the CTF64. In agreement with this,
CTFCr631 is more readily oxidized again in air.
Table 5-3: Number of oxygen vacancies (VO) for one mole of sample.
Sample
Initial [VO]
(mol VO mol-1
CTF)
CTF82
0.08
BCTF64
0.26
CTF64
0.19
CTFCr541
0.16
CTFCr631
0.08
CTCr64
-0.02
CTFCo631
0.21
3+
* Reduction Co to Co2+.
[VO] after Fe4+ ->
Fe3+ reduction
(mol VO mol-1 CTF)
0.11
0.27
0.20
0.18
0.12
-0.02
0.22
[VO] after Cr4+ ->
Cr3+ reduction
(mol VO mol-1 CTF)
0.11
0.27
0.20
0.21
0.17
0.10
0.25*
TPR results
Peaks in TPR data were attributed in the same way as peaks in TPD data
(Paragraph 4.3.1). The products of the TPR experiments are the same as in the
reduction experiments. From the TPR data it follows that iron doped calcium
titanates are stable up to 1100 K in hydrogen gas. The actual temperature for
which the perovskite decomposes could be lower due to slow kinetics. Chromium
doped calcium titanate is more stable than its iron doped counterpart. Even at
1273 K in hydrogen gas, the perovskite is not destroyed. Cobalt doped calcium
titanate is the least stable, the perovskite structure decomposes at 980 K. These
results agree with the expected order of stability as mentioned in the results
(Paragraph 4.2).
The quantitative results from the TPR data are calculated to the percentage of
ions in the material that are reduced. These numbers are in Table 5-4 and
compared to the results from the reduction experiments. Overall, the total oxygen
release calculated from the TPR data is in good agreement with the oxygen
released in the reduction experiments. The smaller peaks in the TPR data,
belonging to the reduction of Fe4+ or Cr4+ to their 3+ oxidation states, are
however relatively too large. When the percentages of Fe4+ and of Cr4+ that are
reduced in TPR are compared to the percentages reduced in TGA data (Table
5-2), TPR data yields an oxygen release that is more than two times too large.
Especially the data for CTFCr541 and CTFCr631 which contain a relative small
amount of chromium is too large: 262% and 200% respectively, instead of
100%. The larger peaks belonging to the reduction of Fe3+ to Fe, are more
accurate, accept in the cobalt-doped sample. The reason for this is most likely the
influence of water, created during the reduction, on the pressure in the ionizing
44
Oxygen vacancy ordering in iron-doped calcium titanate
chamber of the mass spectrometer. Using a smaller amount of sample in the
experiment would likely prevent this.
Table 5-4: Comparison of TPR results with reduction data.
Sample
CTF82
CTF64
CTFCr541
CTFCr631
CTCr64
CTFCo631
Reduction
step
4+->3+ (Fe)
3+->0 (Fe)
Total
3+>0 (Fe)
4+->3+ (Cr)
3+->0 (Fe)
Total
4+->3+ (Cr)
3+->0 (Fe)
Total
4+->3+ (Cr)
3+->2+ (Co)
3+->0 (Fe)
2+->0 (Co)
Total
T range
(K)
TPR Data
O released
(mol O mol-1
CTF)
525 - 817
1153 - 1273
Ions
red.
(%)
740 - 1028
562 - 762
0.05
0.26
0.31
0.66
0.13
0.56
0.69
0.10
0.39
0.49
0.22
0.13
109
260
0.32
0.61
0.68
0.56
0.22
-
978 - 1273
0.94
-
-
1108 - 1273
729 - 966
1152 - 1273
828 - 1033
1152 - 1273
1.07
50
87
Reduction
O released
(mol O mol-1
CTF)
110
262
93
200
87
0.59
5.4 Oxygen permeation results
The flux measured through the CTF materials is in general very low. BSCF and
SCF compositions have a much higher flux than CTF materials, but these
materials lack the stability of iron doped calcium titanates [25]. BSCF and SCF
decompose at oxygen partial pressures of 10-5, making use in syn gas production
impossible because of the very low partial pressures involved in syn gas
production. The material that is most commonly used in the CTF group is the
composition CTF82. Xie et al. [27] measured oxygen fluxes through CTF82
membranes with different thicknesses and for different temperatures. At 1173 K
(50 K below the temperature used in the experiments), with air at the high pO2
side and a pO2” of 4·10-4 atm, a CTF82 membrane with thickness of 0.84 mm,
oxygen flux amounted to 0.245 x 10-8 mol cm-2 s-1. The largest flux was
measured on an 0.84 mm thick membrane at 1373 K, with a pO2” of 4·10-4 atm,
and amounted to 6.37 x 10-8 mol cm-2 s-1.
For the CTF82 membrane tested in this report, the oxygen flux was 1.14 x 10-8
mol cm-2 s-1 at a pO2” of 0.0004 atm, a temperature of 1223 K and a membrane
thickness of 1.0 mm. This is more than four times as high as the oxygen flux Xie
[27]
measured for a thinner membrane, at a temperature 50 K lower, and is of the
some order of the maximum oxygen flux through the CTF82 membrane measured
by Xie [27].
From the measurement of oxygen fluxes through membranes that differed in
thickness only (ranging from 0.84 mm to 2.50 mm), Xie et al. [27] concluded that,
for pO2” values around 10-5 atm and temperatures of 1100 – 1350 K, the
transport of oxygen ions in the bulk of the membranes is the rate-limiting step for
oxygen permeation. These measurements have not been performed on samples
for pO2” values of 10-18 atm, but because of the rather thick membranes and
because of the relatively low ionic and electronic conductivities in these materials,
45
Oxygen vacancy ordering in iron-doped calcium titanate
it is suggested that the fluxes measured with hydrogen atmosphere on the low
pO2 side of the membrane are also bulk limited.
Because of the rather low fluxes measured with oxygen permeation using argon
on the low pO2 side, the measurements are quite inaccurate. Due to a lack of
buffer gas, a small amount of oxygen (and nitrogen) in the argon gas at the inlet
of the reactor has a significant influence on the measured total amount of oxygen
at the outlet of the reactor. Indeed, for measurements in argon, the incoming
argon gas contained a small amount of oxygen and nitrogen. Furthermore the
argon leaving the reactor still contained some nitrogen, probably not from a leak,
but from the initial presence of nitrogen. These peaks were still present even
after flushing the reactor and the tubing with argon for a couple of days. The
small peaks of oxygen in the inlet gas amounted to 80 ppm of oxygen, which is a
significant amount when compared to the oxygen flux through the CTF64
membrane for the highest argon flow rate, 250 ppm. The oxygen flux through the
CTF82 membrane is however significantly higher, and the initial amount of
oxygen present is only a small influence in those measurements.
Another shortcoming of the measurements in argon is the long equilibration time
of the materials. Especially the oxygen flux through the CTFCr631 membrane
continually decreased over four days before it stabilizing. Multiple measurements
for each composition at each flow rate were performed to check the stabilization
of the oxygen flux. CTFCr631 in argon was measured for two different argon flows
due to the long stabilization times and associated sealing issues.
Oxygen permeation performed with argon on the low pO2 side showed that the
composition CTF82 has the highest oxygen flux. CTF64 has the lowest oxygen
flux. This may be attributed to vacancy ordering. Although this composition has
the highest oxygen non-stoichiometry, the vacancies are mostly ordered in
microdomains and do not contribute to oxygen transport. Oxygen transport
through the CTFCr631 material is higher than through the CTF64 composition
because this composition contains less iron and therefore suffers less from partial
ordering than the CTF64 material. CTFCr631 has however a lower oxygen flux
than CTF82. This is partly explained by the fact that the chromium is mostly in 4+
oxidation state. Although chromium is slowly reduced in argon atmosphere
starting at 1073 K, The oxygen flux through the material and the high oxygen
pressure on the air side increase the oxygen pressure at the sweep side, limiting
the reduction of Cr4+ to Cr3+. This is likely the cause of the slow equilibration
observed for this sample. This explains only partly the fact that CTF82 has a
higher flux than CTFCr631. Assuming that Cr is completely in a 4+ oxidation
state, CTFCr631 still has more iron ions than CTF82 that can also contribute to
oxygen transport. Oxygen vacancies in CTFCr631 may still be partially ordered in
argon.
Of greater significance to the focus of this work, materials for syn gas generation,
is the measurement of oxygen flux with a reducing gas on the low pO2 side.
Oxygen permeation performed with hydrogen on the low pO2 side shows that the
target composition CTFCr631 has the highest oxygen flux of all the tested
compositions. For a hydrogen/water mixture, it is assumed that the pO2 is
sufficiently low to allow for a complete reduction of all chromium to 3+, since this
process already starts at 1073 K in argon atmosphere. When all chromium is
reduced, the same number of vacancies should be created in the CTFCr631
material as in the CTF64 material. However the flux of oxygen for the CTFCr631
material is significantly higher than for the CTF64 material, which can be
explained by the mobility of oxygen vacancies. In CTF64, especially under
reducing circumstances, oxygen vacancies are partially ordered in microdomains,
46
Oxygen vacancy ordering in iron-doped calcium titanate
which reduces their mobility. Oxygen vacancies in CTF82 and in CTFCr631 are
less ordered than for the CTF64 material.
It is remarkable that in argon, CTFCr631 has a significantly lower oxygen flux
than CTF82, while in hydrogen, the flux through CTFCr631 is larger. This can not
be due completely to the creation of oxygen vacancies when chromium is
reduced. CTFCr631 should still have a higher oxygen flux in argon than CTF82 if
the same amount of partial vacancy ordering exists in the two materials. A logical
explanation would be that with the reduction of Cr4+ to Cr3+, the presence of
chromium reduces the amount of partial ordering in the structure.
A prominent feature of the oxygen flux in the hydrogen measurements is the fact
that all measured compositions seem to have a bending in the curve of oxygen
flux versus 1000/T, at the same temperature (1125 K). It is not possible to draw
a linear best fit-line through the oxygen fluxes for each temperature, staying
within the error bars. It is also not possible to draw a line through the oxygen
fluxes for the three lower temperatures, only the highest three temperatures can
be captured in a line. This bending takes place at roughly the same temperature
(1150 K) for which the iron in the material reduces from 3+ to metal. That is,
there is an increase in the amount of oxygen vacancies and therefore an
anticipated increase in ambipolar conductivity at this temperature due to the
reduction of iron. The pO2 gradient across the membrane prevents the complete
reduction of the membrane. At the low pO2 surface of the membrane, the
perovskite cells containing iron are however completely destroyed, judging from
XRD data.
47
Oxygen vacancy ordering in iron-doped calcium titanate
6 Conclusion
The report describes the preparation and characterization of iron-doped calcium
titanate materials doped with additional ions on the A and the B site of the
perovskite lattice. The goal of this additional doping was to decrease the amount
of (partial) ordering of oxygen vacancies in the lattice in order to increase the
oxygen ionic conductivity and the transport of oxygen in the material.
It is concluded that regarding the properties researched in this report, doping on
the A site does not affect the properties of the perovskite much. Barium doping
had no effect on the oxygen content and the creation of oxygen vacancies in the
lattice.
Doping on the B site was performed with nickel, cobalt and chromium ions. It is
concluded that it is practically impossible to incorporate nickel into the perovskite
lattice. The introduction of cobalt ions seems to be possible, but it decreases the
stability of the calcium titanate lattice a lot and it is doubtful whether it is possible
to incorporate large amounts of cobalt in the calcium titanate perovskite. Doping
with chromium proved to be well possible. Large amounts of chromium, up to at
least 0.4 moles can be incorporated in the lattice. Chromium is chemically very
similar to titanium in these perovskites, as long as it is present as Cr4+.
The detection of ordering in these materials is difficult, especially where partial
ordering is concerned. Apart from some minor features (peak broadening), no
indications can be found in (HT-)XRD patterns. Oxygen content measurements
reveal information about the creation of oxygen vacancies in the material, but
actual information about ordering is not provided by these measurements. It is
concluded that only Mössbauer spectrometry and neutron scattering can yield
accurate information on partial ordering of oxygen vacancies in these materials.
However, although the amount of partial ordering is not directly measured, it has
an effect on oxygen permeation through these materials. From measurements of
oxygen fluxes through membranes made out of different compositions, it is
concluded that under strongly reducing atmospheres (pO2 = 10-20) oxygen flux
through a composition with chromium doping (CTFCr631) is significantly higher
than through the undoped calcium titanate, CTF64, and slightly better than the
best and most used composition in this group of materials, CTF82. It was
deduced that the reduction of partial ordering because of the chromium plays a
role in this flux improvement.
From the observation that oxygen flux through the CTFCr631 membrane for
experiments with argon on the low pO2 side is lower than for the CTF82
composition, it was concluded that when chromium is present in 4+ oxidation
state it is much comparable to titanium ions, and is not actively reducing
ordering. It is the Cr3+, present in strongly reducing atmosphere that decreases
partial ordering in these structures.
48
Oxygen vacancy ordering in iron-doped calcium titanate
7 Dankwoord
Bij deze wil ik de mensen bedanken die dit verslag in het bijzonder en mijn
afstuderen in het algemeen mogelijk gemaakt hebben. In de eerste plaats gaat
mijn dank uit naar de afstudeercommissie: Dave, bedankt voor de gezelligheid en
de bijzonder goede muziek tijdens het racletten. Henny, dank je wel voor de
nuttige tips en discussies – die zich niet beperkten tot het wetenschappelijke vlak,
maar zich ook uitstrekten tot distorted gitaarsolo’s en dergelijke. Steve, thank
you for all your advice and guidance, I am looking forward to working with you
again; which will be quite soon actually! Kulathuiyer, Jaap en Bernard: dank jullie
wel voor alle nuttige suggesties en voor het doorlezen van het verslag.
Verder wil ik die mensen bedanken die geholpen hebben bij het praktische werk:
Attila, Gerrit, Henk en Wika bedankt voor jullie behulpzaamheid (en geduld) bij
het wegwijs raken op de labzaal; Cindy, Mieke voor het verrichten van de
DTA/TGA metingen; Herman voor zijn deskundigheid op röntgendiffractie gebied,
en Joop voor het polijsten van de zuurstof permeatie membranen en de leuke
gesprekken over Cadzand.
I want to thank all my co-workers for creating a great atmosphere to work in.
Without you guys (and girls) the ten months would have been quite dull.
In het bijzonder wil ik de mensen in het afstudeerhok bedanken dat ze mijn
‘weirde’ gedrag hebben getolereerd – sorry, I’m Dutch, hihi!
Dank gaat uit naar mijn flatgenoten – vooral diegenen die al zeven jaar mijn
‘flitsende’ gitaarspel hebben moeten aanhoren – voor het verzorgen van de
nodige ontspanning.
Tot slot wil ik mijn ouders en zus bedanken voor alle steun die ik heb mogen
ontvangen gedurende mijn gehele studieperiode – zowel financieel als
emotioneel.
En mocht ik JOU niet bedankt hebben, dan komt dat niet doordat ik dat niet
wilde, maar omdat ik het domweg vergeten ben. Ruim zeven jaren studeren doet
dat met je.
49
Oxygen vacancy ordering in iron-doped calcium titanate
8 Literature
1.
F.H.B. Mertins, ‘Perovskite-type ceramic membranes’, 2005, Febodruk, Enschede.
2.
B.E. Watts, ‘Preparation of Lead Titanate, PZT and PZN’, in encyclopedia of
materials science and technology.
3.
R.D. Shannon, ‘Revised Effective Ionic Radii and Systematic Studies of
Interatomic Distances in Halides and Chalcogenides’, Acta Crystallogr. A. 32,
1976, pp. 751-767.
4.
D.F. Shriver, P.W. Atkins, C.H. Langford, Inorganic Chemistry, 2nd edition, 1994,
Oxford University Press, Oxford.
5.
H. Iwahara, T. Esaka, T. Mangahara, ‘Mixed conduction and oxygen permeation in
the substituted oxides for CaTiO3’, Journal of applied electrochemistry 18, 1998,
pp. 173-177.
6.
F.M. Figueiredo, J.C. Waerenborgh et al., ‘On the relationships between structure,
oxygen stoichiometry and ionic conductivity of CaTi1-xFexO3-δ’, Solid State Ionics
156, 2003, pp. 371-381.
7.
C. Wagner, ‘Equations for transport in solid oxides and sulfides of transition
metals’, Prog. Solid State Chem. 10 (1975) pp. 3-16.
8.
A.I. Becerro, C.A. McCammon et al., ‘Displacive phase transitions and
spontaneous strains in oxygen deficient CaFexTi1-xO3-x/2 perovskites (0 ≤ x ≤
0.40)’, J. Phys Condens. Matter 12, 2000, pp. 3661-3670.
9.
A.I. Becerro, C.A. McCammon et al., ‘Oxygen-vacancy ordering in CaTiO3CaFeO2.5 perovskites: from isolated defects to infinite sheets’, Phase Transitions
69, 1999, pp. 133-146.
10. P. Berategui, S.-G. Eriksson, S. Hull, ‘A neutron diffraction study of the
temperature dependence of Ca2Fe2O5’, Materials Research Bulletin 34(2), 1999,
pp. 303-314.
11. J.-C. Grenier et al., ‘Etude par Diffraction X et Microscopie Electronique du
Systëme CaTiO3-Ca2Fe2O5’, Journal of Solid State Chemistry 20, 1977, pp. 365379.
12. S. Hovmöller et al., ‘Structure Determination of Ca4Fe2Ti2O11 by Electron
Microscopy and Crystallographic Image Processing’, Journal of Solid State
Chemistry 77, 1988, pp. 316-321.
13. J. Rodriguez-Carvajal, M. Vallet-Regi, J.M. Gonzalez Calbet, ‘Perovskite threefold
superlattices: a structure determination of the A3M3O8 phase’, Materials Research
Bulletin 24, 1989, p. 423-30.
14. J. Canales-Vazquez et al., ‘Microdomain texture and microstructures of Fe4+containing CaTi0.4Fe0.6O3-δ ‘, J. Solid State Chem. 177, 2004, p. 3105-3113.
15. A.I. Becerro et al., ‘The transition from short-range to long-range ordering of
oxygen vacancies in CaFexTi(1-x)O3-x/2 perovskites’, Phys. Chem. Chem. Phys. 2000
(2), p. 3933-3941.
16. C.A. McCammon, A.I. Becerro et al., ‘Short-range ordering of oxygen vacancies in
CaFexTi1-xO3-x/2, perovskites (0 < x < 0.4)’, J. Phys.: Condens. Matter 12, 2000,
pp. 2969-2984.
50
Oxygen vacancy ordering in iron-doped calcium titanate
17. J.C. Waerenborgh et al., ‘Fe4+ content and ordering of anion vacancies in partially
reduced AfexTi1-xO3-y (A = Ca, Sr; x ≤ 0.6) perovskites. An 57Fe Mössbauer
spectroscopy study’, J. Phys.: Condens. Matter 13, 2001, pp. 8171-8187.
18. Z. Shao et al., ‘Ba effect in doped Sr(Co0.8Fe0.2)O3-δ on the phase structure and
oxygen permeation properties of the dense ceramic membranes’, Separation and
Purification Technology 25, 2001, pp. 419-429.
19. V.L. Kozhevnikov et al, ‘Disordering and Mixed Conductivity in the Solid Solution
LaSr2Fe3-yCryO8+δ’, Chem. Mater. 16, 2004, pp. 5014-5020.
20. A. G. Maddock, ‘Mössbauer Spectroscopy, Principles and Applications of the
Techniques’, 1997, Horwood Publishing Limited, Chichester.
21. D. Scholten, ‘Surface Modification of Yttria Stabilized Zirconia by Ion
Implantation’, 1987, Febodruk, Enschede.
22. A.C. Larsen, R. B. von Dreele, General Structure Analysis System (GSAS) Report
LAUR B6-748, Los Alamos National Laboratory, Los Alamos, NM, 1985.
23. The Physical and Theoretical Chemistry Laboratory Oxford University, ‘Safety
(MSDS) Data for mercury’, Available:
http://ptcl.chem.ox.ac.uk/MSDS/ME/mercury.html
24. F.M. Figueiredo et al., ‘Surface enhanced oxygen permeation in CaTi1-xFexO3δ ceramic membranes’, Journal of Membrane Science 236, 2004, pp. 73-80.
25. T. Esaka et al., ‘Electrical Conduction in CaTi1-xFexO3-δ under Low Oxygen Pressure
and its Applications for Hydrogen Production’, Solid State Ionics 40/41, 1990, pp.
544-547.
26. Cho et al., ‘Synthesis and crystal structure refinement of (1-x)Ca Ti O3 (x)
(La1/3Nd1/3) Ti O3’, RIST Research Papers 12, 1998, pp. 116-123.
27. S. Xie et al., ‘Mixed oxygen ionic and electronic conduction in CaFe0.2Ti0.8O3-δ a
combined oxygen permeation and electrical conductivity study’, Solid State Ionics
118, 1999, pp. 23-28.
51
Oxygen vacancy ordering in iron-doped calcium titanate
9 Appendices
A.1 Error calculations
A.2 Oxygen partial pressure in hydrogen/water mixture
52
Oxygen vacancy ordering in iron-doped calcium titanate
A.1 Error calculation
In this chapter errors are estimated for the oxygen content calculated from
reduction experiments (A.1.1) and for the density measurements (A.1.2).
A.1.1 Error calculation for reductions experiments
To what extent a perovskite is reduced by hydrogen is derived from the release of
oxygen during the reduction. The release of oxygen is calculated by measuring
the weight of the sample before and after the reduction. The complete equation
for calculating the amount of oxygen released in moles of O per mole of
perovskite sample is given by equation A1.1:
nOreleased
⎡( m
⎤
sample + Pt − mPt )before − ( msample + Pt − mPt ) after / M O
⎣
⎦
=
⎡⎣ msample ,after ⎤⎦ / M sample ,after
[A1.1]
In which msample+Pt indicates the mass of the sample and the platinum boat in
which the sample is reduced. Subscripts before and after indicate the mass
measured before and after the reduction respectively. MO and Msample,after are the
molecular weight of oxygen and the sample after reduction, respectively.
From equation A1.1, the relative error can be calculated as follows:
∆nORe lease
nORe lease
=
∆msample + Pt ,after + ∆mPt , after + ∆msample + Pt ,before + ∆mPt ,before
(m
sample + Pt
− mPt )
before
− ( msample + Pt − mPt )
+
after
∆msample + Pt , after + ∆mPt , after
msample, after
[A1.2]
In this equation the molar masses have been omitted because they are constants.
The error in the mass of the platinum boat is estimated to be equal or smaller to
the difference between the mass measured at the end of one reduction
experiment and the mass at the start of the next; the largest difference between
these two masses is 0.0005 g.
The error in the mass of sample+Pt was estimated to be 0.0001 g, since this is
the accuracy of the balance used.
When these estimates are inserted in equation A1.2, the following equation is
obtained, which allows a quick calculation of the error per reduction experiment:
∆nORe lease
nORe lease
=
(m
sample + Pt
− mPt )
0.0012
− ( msample + Pt − mPt )
before
+
after
0.0006
msample, after
[A1.3]
All reductions had an error of about 0.02 moles of O, on a total of about 0.4 –0.6
moles of O released.
53
Oxygen vacancy ordering in iron-doped calcium titanate
A.1.2 Density measurements
Measurements of the absolute density were performed by immersion of a sample
in mercury (A.1.2.1). More interesting is the relative density, which was obtained
by comparing absolute density and XRD-density (A.1.2.2).
A.1.2.1 Absolute density
Absolute density is calculated using equation A1.4:
ρ abs =
mair
m
=
V ( mHg / ρ Hg )
[A1.4]
The relative error in the density can be calculated as follows:
∆ρ abs
ρ abs
=
∆mair ∆mHg ∆ρ Hg
+
+
ρ Hg
mair
mHg
[A1.5]
in which the error in the density of mercury can be neglected (< 0.005). All
masses were measured using a balance with an accuracy of 0.01 g, so ∆mair =
0.01 g. ∆mHg consists of two errors, due to the method of measuring: the sample
is immersed in mercury using a metal ‘cage’. This cage has to be calibrated first,
without the sample. After that, the cage containing the sample has to be
immersed into the mercury to the same level as during the calibration. An error is
made by the balance when measuring the sample weight (= 0.01 g), and a
second error is caused by a difference in the ‘depth of immersion’ between
measurement and calibration. The latter error is estimated to be 0.01 g, since the
change in mass when changing the emersion of the cage was minimal. In
conclusion, the error in the sample mass measured in mercury is about 0.02 g.
The actual error can thus be calculated using:
∆ρ abs
ρ abs
=
0.01 0.02
+
mair mHg
[A1.6]
Since most samples were small discs the mass is not very high and the absolute
error in the density can be as high as 0.11 g cm-3 (= 3%). The density of a
number of discs was measured twice. Both measured densities are equal within
error range, the largest error being 0.05 g cm-3.
A.1.2.2 Relative density
The relative density was calculated by dividing the absolute density by the XRDdensity and multiplying with 100%. XRD-density was calculated from the lattice
parameters obtained by X-ray diffraction. For a cubic, orthorhombic and
tetragonal cell, the XRD density can be calculated using equation A1.7:
54
Oxygen vacancy ordering in iron-doped calcium titanate
ρ XRD =
M sample
Vunitcell
cubic
⎯⎯⎯
→
M sample
In which a, b and c are the parameters of the unit cell.
density is then calculated as follows:
∆ρ XRD
ρ XRD
=
∆M sample
M sample
+
[A1.7]
a ⋅b ⋅c
The error in the XRD
∆a ∆b ∆c
+
+
a
b
c
[A1.8]
The molecular weight of the sample is not completely constant. It is dependant on
the amount of oxygen that is present in the perovskite. This amount is calculated
from the reduction experiments. The error in the oxygen content for these
experiments is calculated using equation A1.3 and amounted to 0.02 moles of
atomic oxygen for all samples measured. The error in the molecular weight of the
samples is therefore estimated to be 0.02 * 16.00 = 0.32 g.
The lattice parameters a, b and c are estimated using the GSAS software. The
basic fitting of the lattice parameter is quite accurate; the fine-tuning of the X-ray
diffraction pattern changes the lattice parameter only in the third decimal
number. Therefore, the error in the lattice parameter is estimated to be 0.005 Å.
Equation A1.9 can be used to estimate the error in the relative density:
∆ρ rel
ρ rel
⎛ ∆ρ
∆ρ XRD ⎞
= ⎜ abs +
⎟ ⋅100% =
ρ XRD ⎠
⎝ ρ abs
⎛ 0.01 0.02
0.32
0.0005 0.0005 0.0005 ⎞
+
+
+
+
+
⎜⎜
⎟⎟ ⋅100%
m
m
M
a
b
c
air
Hg
sample
⎝
⎠
[A1.9]
This error is a relative error. The absolute error (in %) can be obtained by
multiplying the relative error with ρrel.
55
Oxygen vacancy ordering in iron-doped calcium titanate
A.2 Oxygen
mixture
partial
pressure
In this chapter, the oxygen partial pressure,
in
hydrogen/water
PO2 , in a hydrogen/water (and
argon) mixture will be calculated.
Hydrogen/water equilibrium
The oxygen concentration in a hydrogen/water mixture is set by the equilibrium
between hydrogen, oxygen and water:
2 H 2 ( g ) + O2 ( g )
2 H 2O ( g )
[A2.1]
Since this is an equilibrium, there is an equilibrium constant describing the
reaction thermodynamically:
PH2 O
H 2O ]
[
K c ( p, T ) =
=
2
[ H 2 ] ⋅ [O2 ] PH2 ⋅ PO
2
2
2
= KP
[A2.2]
2
in which KC and KP are the equilibrium constants calculated by concentration and
by partial pressure, respectively. KC = KP is only valid for gas mixtures. The
oxygen partial pressure can be calculated when the equilibrium constant and the
other partial pressures are known. The partial pressure of hydrogen in a
hydrogen/water/argon mixture is measured by the gas chromatograph. The
equilibrium constant and the water partial pressure have to be calculated.
Calculating the equilibrium constant
The equilibrium constant KC is a thermo dynamical constant and it can be
calculated from the change in the Gibbs free energy of reaction:
Kc = e
−
∆Gr
RT
[A2.3]
The change in Gibbs free energy is calculated from the enthalpy change and the
entropy change of the reaction (at standard pressure and T = 298 K):
∆Gr = ∆H θr − T ∆S θr
[A2.4]
in which ∆Hr = -4.84▪105 J mol-1 and ∆Sr = -89 J mol-1 K-1. KC is then calculated
using the temperature of the reactor in Kelvin.
56
Oxygen vacancy ordering in iron-doped calcium titanate
Calculating the amount of water vapour
The water in the H2/H2O mixture originates from a bubbler, put in front of the
actual reactor. The bubbler is kept at room temperature. The liquid water in the
bubbler will form an equilibrium with a phase of gaseous water. This gaseous
water is swept away by the hydrogen flow. The partial pressure of the water
vapour at the inlet of the reactor can be calculated using the Antoine equation:
B ⎞
⎛
⎜ A−
⎟
C +T ⎠
PH 2O (T ) = 10⎝
[A2.5]
in which A, B and C are constants (usable within a certain temperature range)
and T is the temperature in Kelvin at which the equilibrium is set (i.e. room
temperature). PH 2O (298 K) = 0.031 bar.
Table A2-1: Antoine constants for water.
Constant
A
B
C
Value
6.20963
2354.731
7.559
The amount of water in the hydrogen/water mixture at the outlet of the reactor is
larger than the amount at the inlet because of the creation of water in the
reactor. The amount of water on the outlet is calculated by adding the amount of
water created (which is equal to the amount of hydrogen consumed) to the
amount of water at the inlet.
57