Electro-chemo-mechanical Studies of Perovskite-Structured Mixed
Ionic-Electronic Conducting SrSni-xFex03-x/2+6
by
Chang Sub Kim
B.S. Physics
California Institute of Technology, 2013
SUBMITTED TO THE DEPARTMENT OF MATERIALS SCIENCE AND ENGINEERING IN
PARTIAL FULFILLMENT OF THE REQUIREMENTS FOR THE DEGREE OF
MASTER OF SCIENCE
IN MATERIALS SCIENCE AND ENGINEERING
MASSACHUSETTS INSTITUTE
OF TECHNOLOGY
AT THE
MASSACHUSETTS INSTITUTE OF TECHNOLOGY
OCT 28 2015
SEPTEMBER 2015
LIBRARIES
C 2015 Massachusetts Institute of Technology. All rights reserved.
Signature redacted
Authored by
Chang Sub Kim
Department of Materials Science and Engineering
July 24, 2015
A
Signature redacted
Certified by
Hari L. Tuller
Professor of Ceramics an Electronic aterials
Thesis Supervisor
Accepted by
Signature redacted______
Donald R. Sadoway
Chair, Departmental Committee on Graduate Students
I
2
Electro-chemo-mechanical Studies of Perovskite-Structured Mixed
Ionic-Electronic Conducting SrSnl-xFex03-x/2+6
by
Chang Sub Kim
Submitted to the Department of Materials Science and Engineering
On July 24, 2015
in Partial Fulfillment of the Requirements for the Degree of
Master of Science in Materials Science and Engineering
ABSTRACT
High efficiency and fuel flexibility make solid oxide fuel cells (SOFCs) attractive. However, when
operating at reduced temperatures, there is significant loss in efficiency, of which slow surface reaction
kinetics at the cathode are most responsible. Previously, the mixed ionic and electronic conducting
(MIEC) perovskite-structured SrTixFex03-O2+6 (STF) materials system was identified as a promising
candidate for SOFC cathodes given rapid oxygen surface exchange kinetics. The exchange kinetics were
correlated with the minority electron charge density in STF, which in turn depends on its defect chemistry
and band structure.
In this work, an alternate B site host cation, Sn, was selected to replicate and extend the STF studies, due
to its distinct band structure and higher electron mobility. Oxygen nonstoichiometry and the defect
chemistry of the SrSnxFex03-/2+6 (SSF) system were examined by means of thermogravimetry as a
function of oxygen partial pressure in the temperature range of 973-1273 K. Marginally higher
reducibility was observed compared to corresponding compositions in STF system. The bulk electrical
conductivity was measured in parallel to examine how changes in defect chemistry and electronic band
structure associated with the substitution of Ti by Sn impact carrier density and ultimately electrode
performance. Bulk chemical expansion was measured by dilatometry as a function of oxygen partial
pressure while surface kinetics were examined by means of AC impedance spectroscopy. The electrochemo-mechanical properties of SSF were found not to differ significantly from the corresponding
composition in STF. It is believed that Fe dominates the character of the valence and conduction bands
and thus governs the electronic properties in SSF. Though slightly shifted by Sn's larger size, the defect
equilibria - including the oxygen vacancy concentration - were found to also be largely dominated by Fe
and thus differed only in a limited way from that in STF. Key thermodynamic parameters of
SrSnosFeo.3502.825 6 SSF35 obtained include the reduction enthalpy (4.30 eV) the electronic band gap
(1.72 eV) and the anion Frenkel enthalpy (0.52 eV). Key kinetic parameters include the migration
enthalpy of oxygen vacancies (0.70 eV), the activation energy of area-specific-resistance (1.65 eV) and
the electron (0.0002 0.00005 cm 2 /V-s) and hole (0.0037 0.0015 cm 2 /V-s) mobilities. With the surface
exchange rate nearly identical to the STF35 counterpart, the main advantage of SSF35 as a SOFC
electrode would be its enhanced chemical stability and slower degradation.
Thesis Supervisor: Harry L. Tuller
Title: Professor of Ceramics and Electronic Materials
3
TABLE OF CONTENTS
LIST OF FIGU RES......................................................................................................................................
6
LIST OF TA BLES .......................................................................................................................................
8
ACKN OW LED GEM ENTS .........................................................................................................................
9
CHA PTER 1. IN TRODUCTION ...............................................................................................................
10
1.1
M otivation ..................................................................................................................................
10
1.2
Solid oxide fuel cells (SO FCs).................................................................................................
12
1.2.1
Principles of operation and characteristics .....................................................................
12
1.2.2
Technical challenges of SO FCs......................................................................................
14
1.2.3
SOFC Cathodes ..................................................................................................................
15
1.3
M ixed ionic and electronic conducting (M IEC) oxides ..........................................................
16
1.3.1
M odel system STF..............................................................................................................
17
1.3.2
M odel system SSF ..............................................................................................................
18
Objectives of the research...........................................................................................................
19
CHA PTER 2. EX PERIM EN TS .................................................................................................................
20
2.1 Sam ple Preparation ..........................................................................................................................
20
1.4
2.1.1
Powder synthesis ................................................................................................................
20
2.1.2
Bulk sam ples ......................................................................................................................
20
2.1.3
Deposition of STF and SSF thin films by PLD ...............................................................
20
2.1.4
Deposition of Au or Pt current collectors........................................................................
21
2.2
Physical Characterization .......................................................................................................
22
2.2.1
X-ray diffraction .................................................................................................................
22
2.2.2
Density m easurem ent .....................................................................................................
22
2.3 Therm ogravim etry............................................................................................................................
22
2.4 Sim ultaneous conductivity and chem ical expansion m easurem ents .............................................
23
2.4.1
Bulk sam ple configuration...............................................................................................
23
2.4.2
M easurem ent setup .............................................................................................................
23
2.4.3
M easurem ent conditions.................................................................................................
24
2.5 Thin film EIS m easurem ents............................................................................................................
24
2.5.1
Cell configuration ...............................................................................................................
24
2.5.2
M easurem ent conditions.................................................................................................
25
CHA PTER 3. RESULTS............................................................................................................................
3.1 Physical/chem ical characterization of SSF powder and thin film s ...............................................
26
26
3.1.1
SSF powders.......................................................................................................................
26
3.1.2
SSF thin film s .....................................................................................................................
26
3.2 Therm ogravim etric analysis.............................................................................................................
29
4
3.3 Bulk conductivity.............................................................................................................................
31
3.4 Bulk chem ical expansion .................................................................................................................
33
3.5 EIS results of SSF thin film electrodes ............................................................................................
35
CHA PTER 4. DISCU SSION .....................................................................................................................
37
4.1 Defect chem ical model.....................................................................................................................
37
4.1.1
Defect chem istry of STF and SSF ...................................................................................
37
4.1.2
M odel defect diagram .....................................................................................................
40
4.1.3
Defect diagrams and thermodynamic parameters of STF35 and SSF35 ........................
41
4.2 Conductivity m odel..........................................................................................................................
44
4.3 Chem ical expansion .........................................................................................................................
49
4.4 Investigation of surface exchange kinetics of STF/SSF thin film s ...............................................
50
4.5 Problem s w ith SSF5 and SSF50 ..................................................................................................
50
4.5.1
Phase decom position of SSF5 .........................................................................................
50
4.5.2
Degradation of SSF50 thin film electrode ......................................................................
51
CHA PTER 5. CON CLU SION ...................................................................................................................
52
5.1 Sum m ary ..........................................................................................................................................
52
5.2 Recom m endations for Future W ork.............................................................................................
53
REFERENCES ...........................................................................................................................................
5
55
LIST OF FIGURES
Figure 1. A reaction diagram of a SOFC running in fuel cell (top) and electrolysis (bottom) modes. ...... 11
Figure 2. Renewable energy cycle diagram. SOFCs running in electrolysis mode to generate fuel with
excess energy and run in fuel cell mode to convert back to electrical power..........................................
14
Figure 3. Area specific resistance of a SOFC, its components - cathode, electrolyte, anode - and charge
carrier diffusion. Reprinted from a reference [6].....................................................................................
15
Figure 4. Sketches of the three reaction paths of the oxygen reduction and incorporation reaction and
some possible rate-determining steps. Modifications of the paths (e.g., adsorption of a molecular rather
than an atomic species or diffusion along the cathode/electrolyte interface) and a combination of electrode
and electrolyte surface paths (adsorption on cathode and surface diffusion onto the electrolyte surface) are
also possible. Reprinted from a reference [7].........................................................................................
16
Figure 5. Ec-EF (left axis) and activation energy of RSTF (or k) measured by EIS (right axis) as a function
of Fe composition at 600'C in air. Reprinted from a reference [12].......................................................
18
Figure 6. Variation of the total conductivity vs. oxygen partial pressure at 750'C of SrSnxFe,03-6 (0 x
1). Conductivities at highly reducing and oxidizing conditions were measured, but not at intermediate
oxygen partial pressure where an ionic regime is expected. Reprinted from a reference [13]................ 19
Figure 7. A schematic illustration of the current collector pattern. UV light passes through the white
region (transparent, not coated with chrome), cross-linking the photoresist polymer. Photoresist on the
black region (coated with chrome) dissolves in the developer..............................................................
21
Figure 8. A Schematic illustration of the SSF bulk sample with four-point probe setup for simultaneous
conductivity/chem ical expansion measurem ents...................................................................................
23
Figure 9. Schematic illustrations of asymmetric cells with YSZ electrolyte, SSF working electrode with
(a) buried Pt current collector and (b) top Au current collector. Both cells had porous Ag counter
e le ctro d e s....................................................................................................................................................
24
Figure 10. XRD pattern of SrSn -,Fex03.6 (x
=
0, 5, 35) powders synthesized .....................................
Figure 11. HRXRD patterns of SrSnO3 thin films deposited on YSZ (100) single crystal substrates by
PLD at 850'C and 950'C heater tem peratures .......................................................................................
26
27
Figure 12. HRXRD patterns of SSF50 thin films deposited on YSZ (100) single crystal substrates by PLD
at 850'C heater temperature, and SSF35 thin film at 850'C and 9500 C.................................................
28
Figure 13. HRXRD patterns of STF35 thin films deposited on YSZ (100) single crystal substrates by PLD
at 600'C and 850'C heater tem peratures.................................................................................................
28
Figure 14. Oxygen nonstoichiometry - 6 - as a function of p02 and temperature for SSF35. .............. 30
Figure 15. Comparison of oxygen nonstoichiometry of STF35 [19] and SSF35 at 700'C and 1000'C.... 31
Figure 16. Bulk conductivity of SSF35 as a function of p02 and temperature. ....................................
6
32
Figure 17. Comparison of bulk conductivities of STF35 [20] and SSF35 at 8001C. ..............................
32
Figure 18. A schematic model of a material's expansion as a function of temperature, with the chemical
expansion contribution noted. Reprinted from a reference [22]............................................................ 33
Figure 19. SSF bulk sample length in different oxygen partial pressures at a constant temperature of
900'C m easured by a dilatom eter...............................................................................................................
34
Figure 20. Expansion of SSF35 as a function of 6 at 750'C. ..................................................................
34
Figure 21. Typical electrochemical impedance spectroscopy (EIS) spectrum of a SSF35 thin film
deposited onto a single crystal YSZ electrolyte with buried Au current collectors and porous Ag counter
35
electrode m easured in air at 468'C .............................................................................................................
Figure 22. Comparison of temperature dependence of area specific resistance (ASR) of SSF35 and STF35
36
th in film electro des.....................................................................................................................................
Figure 23. Schematic defect diagram using Brouwer approximation for the model system SSF and STF.41
Figure 24. Carrier concentrations from defect equilibria at 10000 C. ....................................................
42
Figure 25. Natural log of equilibrium constants K as a function of 1/kT...............................................
43
Figure 26. Temperature dependence of p-type (p 0 2= 0.316 atm) , n-type (pO2 = 10-18 atm) , and PO2insensitive ionic conductivities of SSF35..............................................................................................
45
Figure 27. Oxygen partial pressure dependence of (a) n-type and (b) p-type conductivities of SSF35
extracted by subtracting the respective ionic conductivities at each isotherm. .......................................
45
Figure 28. Temperature dependence of electron and hole mobilities of SSF35......................................
47
Figure 29. Expected conduction and valance band positions of STF and SSF as a function of Fe content,
x . .................................................................................................................................................................
7
52
LIST OF TABLES
Table 1. Summary of properties of different types of fuel cells. Reprinted from a reference [5]........... 13
Table 2. Ionic radii of Ti"* and Sn", lattice parameters and volumes of STF35 and SSF35 calculated from
X RD peak s. ................................................................................................................................................
29
Table 3. The predicted solutions to the defect solutions with Brouwer approximations.........................
40
Table 4. Thermodynamic parameters for STF35 [19] and SSF35 from fitting the nonstoichiometry 6 as a
function of oxygen partial pressure at different temperatures.................................................................
43
Table 5. Thermodynamic parameters for STF35 and SSF35 from fitting nonstoichiometry - 6 - data as
functions of oxygen partial pressure over a range of temperatures simultaneously...............................
44
Table 6. Electron and hole mobilities of SSF35 calculated from defect chemistry and conductivity data. 46
Table 7. Activation energies of n-type conductivity, electron mobility, electron concentration, and the
reductio n enthalpy . .....................................................................................................................................
48
Table 8. Activation energies of p-type conductivity, hole mobility, hole concentration, and activation
energies of hole concentration at intermediate and high pO2.................................................................
48
Table 9. Activation energies of ionic conductivity, oxygen vacancy concentration, and enthalpy of oxygen
vacan cy m igratio n . .....................................................................................................................................
49
8
ACKNOWLEDGEMENTS
I would like to express my sincerest gratitude to my advisor, Professor Harry Tuller, both
intellectually and personally. He is a genuine educator, an insightful researcher, a passionate lecturer, and
my personal mentor. His guidance was available whenever I needed, and was always above my
expectations.
Members of the Tuller group must be acknowledged for their helpful discussions, comments,
suggestions in our offices, group meetings, and labs. Dr. Jae Jin Kim has guided me since I first arrived at
the Institute. I cannot count how many questions I have asked and how many discussions we had. Drs.
Sean Bishop and Nicola Perry have generously shared their experience and insight. Dr. Stuart Cook has
helped with fitting the thermogravimetric data. I would also like to acknowledge Drs. Di Chen, Kunal
Muhkerjee, Nikolai Tsvetkov, and Michael Campion for their helpful discussions, and all other members,
visiting students and scholars in the group, as well as Elisabeth Anderson.
There are so many individuals and groups at and outside MIT I would like to thank: Kurt
Broderick and other members at the Microsystems Technology Laboratories (MTL), members of
KGMSE and MIT Sailing, just to name a few, have broadened my views.
Lastly, I would like to thank my parents, my only brother, and my wife Kyoung-Won for their
support and unconditional love.
Funding for this research was provided by Skoltech Center for Electrochemical Energy Storage
(CEES).
July 17, 2015
Chang Sub Kim
9
CHAPTER 1. INTRODUCTION
1.1 Motivation
The discovery of electricity has changed the way people live. Everyday life begins and ends with
electricity. The world's electrical energy consumption has been rapidly increasing over the years, and the
majority of energy production finds its source from fossil fuels [1]. However, there are limited fossil fuel
reserves, and the efficiency of converting fossil fuels into electricity is limited by the efficiency of the
Carnot cycle:
W
TC
QH
TH
W
where W is the work produced by the system,
QH
_
_1
(1)
is the heat going into the system, Tc is the temperature
of the cold reservoir, and TH is the temperature of the hot reservoir.
Demand for electricity continues to rise, and the world needs more sustainable, renewable, and
efficient way of generating electricity. Solar, wind, hydro, and geothermal technologies are sustainable
and have become the most prominent sources of renewable energy. All sustainable technologies,
however, have limitations: solar energy is limited by time (only available during the day), while wind,
hydro and geothermal by location. Therefore, efficient means of energy storage and transportation are
required to control the supply of sustainable energy. Electrochemical devices, such as batteries and fuel
cells, can solve the problem by redistributing the uneven supply of renewable energy. Batteries store
electrical energy in chemical form while charging, and it is converted back to electrical form during
discharge. Batteries are widely used in portable electronics such as cell phones and laptops, but low
energy density limits their scalability. Fuel cells are also electrochemical devices, but unlike batteries,
fuel cells do not store energy in devices, but generate electrical power from the electrochemical reaction
of fuels - usually hydrogen or hydrocarbons - with oxygen (Figure 1). Therefore, fuel cells are able to run
continuously as long as the fuel and oxidant are provided. The high energy density of chemical fuels is
particularly advantageous when it comes to the transportation sector and thus the use of fuel cells rather
than combustion engines to provide the needed energy offers higher efficiency as well as reduced
10
difficulties and
emissions. Despite these advantages, fuel cells are not widely used due to technological
high cost.
2e-2
2e
2e
2e-
0
H2gas
H 2 0 + 2e (- H 2 +
2e
/1202 +2e-
2
02-
,gas
--
t2e2
4-
-4
--
H
22
-
2e
>
--0
2.cras
H 2 0gas
H 2 0 + 2e -
H2 + 02
2
/202 + 2e
(bottom) modes.
Figure 1. A reaction diagram of a SOFC running in fuel cell (top) and electrolysis
11
1.2 Solid oxide fuel cells (SOFCs)
1.2.1
Principlesof operationand characteristics
Fuel cells can be categorized according to the types of electrolyte membranes (Table 1). Different
electrolytes conduct different ions, and thus exhibit different reactions at the electrodes, operate with
different fuels and in different temperature regimes. SOFCs are characterized by metal oxide electrolytes
which conduct oxygen ions. Solid electrolytes have several advantages over liquid or polymer
electrolytes, including thermal, mechanical and chemical stability against corrosion and contamination.
Operation at high temperatures (between 700'C and I 0000 C) gives SOFCs fuel flexibility: certain
hydrocarbons, such as methane, can be used directly as a fuel without external reforming, because C-C
bonds in hydrocarbons can be easily broken at the temperature range. Furthermore, in contrast to polymer
electrolyte membrane (PEM) fuel cells, they do not require noble metal catalysts such as Pt and are less
susceptible to poisoning.
SOFCs can also be operated reversely as electrolysis cells by applying an electric potential, as
shown in Figure 1. A number of studies were conducted using conventional SOFCs running in reverse
mode, or symmetric cells with electrodes more chemically stable in both oxidizing and reducing
conditions [2-4]. Excess energy from solar, wind, or hydro can be stored as hydrogen or hydrocarbon
fuels by running reversible SOFCs in electrolysis mode to decompose water or C0 2, and the fuels can
then be converted back to electrical power by running the SOFCs in the fuel cell mode as demand
increases. This enables a carbon-free or carbon-neutral energy supply by combining renewable energy
with SOFCs (Figure 2).
12
I
Table 1. Summary of properties of different types of fuel cells. Reprinted from a reference [5].
- Backup power
* Portable
power
<120C
<1 kW
100 kW
- Distributed
generation
-
Fransportation
- Specialty
vehicles
- Solid electrolyte
reduces corrosion
electrolyte
management
problems
- Sensitive to fuel
<100 C
-
100kW
*Space
- Wider range of
stable materials
a lloxw s lower cost
- Sensitive to C02
in fuel and air
components
* Low temperature
- Backup power
- Quick start-up
* Transportation
Phosphoric acid
soaked in a
porous matrix or
imbibed in a
polymer
membrane
100 kW
150 -200C
module
.liqum
d
- Distributed
generation
- Suitable for
combined heat and
power (CHP)
sxstem
Flectrolvte
management
(aqueous)
- Electrolyte
conductivity
(polymer)
5 Expensive
catalysts
Long start-up time
PAFC);
(polymer
membane)to
membrane)
Molten lithium,
sodium, and/or
potassium
carbonates.
soaked in a
porous matrix
impurities
- Flectric utility
600
6
-
700 C
fuel
.o
.mpri-e
.ue
impurities
- Hybrid/gas turbine
cycle
300 kW - 3
M W, 300
kW module
- Increased tolerance
- Distributed
generation
Fe lxblt
Fuel flexibility
b for CP
,
. .s
Sulfur sensitivity
.
Aqueous
potassium
hydroxide
soaked in a
porous matrix, or
alkaline polymer
membrane
- Military
- Expensive
catalysts
- Quick start-up and
load following
&
acid
-
Perifluorosulfon ic
Low temperature
High temperature
corrosion and
breakdown of cell
components
- Long start-up time
- Low power
density
- High efficiency
Yttria stabilized
zirconia
500 - I 000"C
I kW
-
- Auxiliary
power
. .
- Hligh efticiency,
- Limited number
of shutdowns
- Electric utility
- Fuel flexibility
- Long start-up time
- Solid electrolyte
IMW
- Suitable for CIP
- High temperature
corrosion and
breakdown of cell
components
2
"Distributed
generation
.
- Hybrid/gas turbine
Cycle
13
A
160SolarBatteries
IMM-Wind
Electricit
HydroStrchmclerg
Exrach
c ess:
g
Electrolysis
SOFC
Figure 2. Renewable energy cycle diagram. SOFCs running in electrolysis mode to generate fuel
with excess energy and run in fuel cell mode to convert back to electrical power.
1.2.2
Technical challenges of SOFCs
While high operating temperature offers fuel flexibility and high efficiency, it is also responsible
for the high cost of interconnect, support and sealant materials. Less expensive materials can be used in
intermediate temperature SOFCs (IT-SOFCs); however, lower temperatures exponentially decrease
electrochemical reaction rates and ionic transport. To overcome these drawbacks, two approaches can be
followed: thinner electrolyte to reduce loss in ionic transport - less ohmic resistance - and enhance
surface reaction kinetics - less polarization resistance. Slow surface reaction kinetics at the SOFC cathode
is believed to be the major contributor to losses (Figure 3). The processes controlling these losses are also
not well understood.
14
,.
E
C
2.0
-Cell resistance
-- Cathode
1
Electrolyte
C
-
-*-Anode
Diffusion
1.0
0.5
0.0
600
700
900
800
1000
Temperature [*C]
Figure 3. Area specific resistance of a SOFC, its components - cathode, electrolyte, anode - and
charge carrier diffusion. Reprinted from a reference [6].
1.2.3
SOFC Cathodes
The overall reaction at the SOFC cathode known as the oxygen reduction reaction (ORR), can be
described in Kriger-Vink notation as:
102(gas) + 2e' + V0 -
L
(2)
where VI represents a doubly positive charged oxygen vacancy, e' an electron, and Ox an oxygen ion in a
normal oxygen lattice site.. Both electrons and oxygen vacancies must be transported to the site of
reaction. Three reaction paths are possible as illustrated in Figure 4: electrode surface path, bulk path, and
electrolyte surface path [7]. In the electrode surface or electrolyte surface paths, oxygen gas molecules
adsorb on the surface of the cathode or electrolyte and diffuse to a triple-phase boundary (TPB) where
electrolyte, electrode, and gas phase co-exist. ORR occurs along the TPB for these paths, and the number
of active sites is thus directly proportional to the TPB length. When the cathode is able to conduct both
electrons and ions, ORR can take the bulk path, and the entire surface area of the cathode becomes active.
15
electrode surfacepath
bulk path
electrolyoe surfacepath
02
Figure 4. Sketches of the three reaction paths of the oxygen reduction and incorporation reaction
and some possible rate-determining steps. Modifications of the paths (e.g., adsorption of a
molecular rather than an atomic species or diffusion along the cathode/electrolyte interface) and a
combination of electrode and electrolyte surface paths (adsorption on cathode and surface diffusion
onto the electrolyte surface) are also possible. Reprinted from a reference [7].
The current state-of-the-art cathode material, lanthanum strontium manganite (LSM), exhibits
high electronic conductivity, but poor ionic conductivity, and therefore ORR is limited to the TPB length.
There is great interest in new cathode materials with mixed ionic and electronic conductivities to
maximize the ORR rate.
1.3 Mixed ionic and electronic conducting (MIEC) oxides
Mixed ionic and electronic conducting (MIEC) oxides have high electronic carrier densities,
typically associated with multivalent transition metal ions, and high deficiency or excess of oxygen,
which facilitates migration of oxygen ions in the electrode lattice via oxygen vacancies or interstitials.
Certain fluorite oxides, such as CeO2_6 and its derivatives, and perovskite-structured oxides, such as La
xSrxCoO3 (LSC)-based materials, meet these criteria and show mixed conductivity [8-9]. During
processing and/or operation at high temperatures, LSC-based cathodes react with YSZ, and therefore
require use of a different electrolyte material such as (La,Sr)(Ga,Mg)03 (LSGM), or an additional
interlayer - "buffer" layer - such as Gd-doped ceria [10-11].
16
1.3.1
Model system STF
SrTilxFex03-6 (STF) is a solid solution of the wide band gap semiconductor strontium titanate
(SrTiO3) with the MIEC strontium ferrite (SrFeO 2.5). STF is a MIEC with p-type behavior in high
pO2
(i.e. cathode conditions), p0 2-independent ionic behavior at intermediate pO2, and n-type behavior at low
pO2 (i.e. anode conditions). More description of the relevant defect model is included in the discussion
chapter.
By varying the ratio of Fe to Ti, one can systematically control both the ionic and electronic
conductivities of STF over many orders of magnitude. This aspect was viewed as useful in attempting to
better understand the role of these defect species in influencing the surface kinetics of the model SOFC
cathode. The position of the conduction band (Ec), relative to Fermi energy (EF), defines to minority
charge carrier (electron) density in conduction band (n) and is given by the following expression
Ec - EFF
=
-kn
( --n )
(3)
where k is the Boltzmann constant, T is the absolute temperature, and N( is the effective density of
conduction band states.
A strong correlation was found between the magnitude of Ec - EF and the activation energy (Ea)
associated with the surface exchange coefficient (k) from electrochemical impedance spectroscopy- see
Figure 5 [12]. This suggested that the minority carrier density n could be the rate determining species in
the ORR in STF.
17
3
30
310
0 30
6000C in air
D
E -EF
Ea of ASR (k)
2.5
2.55 m
0
2.0
uJ' 2 .0
I
U
<
1.5 -1.5
1.0
1.0
0
20
60
40
80
100
Fe mol%
Figure 5. Ec-EF (left axis) and activation energy of RSTF (or k) measured by EIS (right axis) as a
function of Fe composition at 600 0 C in air. Reprinted from a reference [121.
Model system SSF
1.3.2
The analogous model system SrSnIl-Fex036 (SSF) is proposed in this work to serve as an extension
to the previous study. In the solid solutions of the wide band gap semiconductor strontium stannate
(SrSnO3) with MIEC strontium ferrite (SrFeO 2 ,5), the high electron mobility Sn 5s orbital in SSF replaces
the Ti 3d orbital in STF as the conduction band. This is expected to lead to a considerable increase in
electron mobility for SSF for electrons in the conduction band. There have been some studies on crystal
structure and electrical properties of SSF [13-16]. However, they were limited to very specific conditions
and no comprehensive studies, such as defect chemical modeling, were carried out.
18
0
hA61.......
-1
..
.
.....
10 % Fe in
.. ...+ -4
.A
S..
v
C.
-2 -
-
tP,
miss+ng
i*n.
-3-mCls
-
50
60
o' 70
A
-4
90
v SrFeO 3
o SrSn 3
75 *C
....
-5...
-25
20
30
40
x
..........
o
SrSn03
15
-20
-15
-10
-5
0
log, 0P0 (atm)
Figure 6. Variation of the total conductivity vs. oxygen partial pressure at 750*C of SrSn..Fe,03-6
(05x<1). Conductivities at highly reducing and oxidizing conditions were measured, but not at
intermediate oxygen partial pressure where an ionic regime is expected. Reprinted from a reference
[131.
1.4 Objectives of the research
Previous results from the STF model system presented an interesting aspect of surface reaction
kinetics. In order to extend the study, an analogous perovskite-structured oxide SSF is selected as a new
candidate material. Key objectives of this research lie with assessing electro-chemo-mechanical properties
of SSF as a potential SOFC cathode material:
1.
Establish a defect chemical model and derive the relevant thermodynamic parameters.
2.
Develop a transport model based on the above defect equilibria.
3.
Investigate surface reaction kinetics by electrochemical impedance measurements.
4.
Investigate dilation (chemical expansion) of SSF as function of pO2 at varying temperatures to
obtain the coefficient of chemical expansion (CCE) and evaluate chemo-mechanical stability.
5.
Compare the defect chemistry, conductivity, surface reaction, and chemical expansion of SSF
with STF to assess the origin of electro-chemo-mechanical properties.
6.
Determine the suitability of SSF as a superior SOFC cathode material
19
CHAPTER 2. EXPERIMENTS
2.1 Sample Preparation
2.1.1
Powder synthesis
SSF solid solutions were prepared by the ball milling method. Strontium carbonate (Alfa Aesar,
99.99%), tin (IV) oxide (Alfa Aesar, 99.9%), and iron (III) oxide (Alfa Aesar, 99.945%) powders were
mixed with the desired Sr/Sn/Fe ratio, ball milled with deionized water and ethanol for six hours, dried,
and calcined in air at 1300'C for six hours. The cubic perovskite phase was confirmed by powder X-ray
diffraction (XRD).
2.1.2
Bulk samples
Two different shapes of bulk samples were prepared. A cylindrical sample was prepared to serve
as a pulsed laser deposition (PLD) target, and a rectangular bar sample was prepared for bulk conductivity
and chemical expansion studies. The prepared powders were placed in either a cylindrical or rectangular
bar shaped stainless steel die, uniaxially pressed at 20 MPa for two minutes, and then sintered in air at
1500'C for six hours.
2.1.3
Deposition of STF and SSF thinfilms by PLD
STF and SSF thin films were prepared by means of PLD (Neocera Inc., Beltsville, MD). After
loading (001) oriented YSZ single crystal substrates (10xlOxO.5
mm3 ) from MTI Corporation
(Richmond, CA) and a bulk target of desired composition, the PLD chamber was pumped down to a base
pressure of 9x 106 Torr, and then the substrate was heated to 850'C or 950'C. The substrate temperature
was calibrated to be about 150'C lower than the heater temperature. A Coherent COMPex Pro 205 KrF
eximer laser (Santa Clara, CA) of 248nm wavelength with 400 mJ/pulse at 10Hz was used to ablate STF
and SSF targets. The surfaces of the targets were pre-ablated with 3,000 pulses before every deposition.
The chamber was maintained at 10 mTorr oxygen pressure during deposition and 10 Torr after deposition,
before cooling down to room temperature.
20
2.1.4
Deposition ofAu or Pt currentcollectors
to
Because the mobilities of the majority electron holes are low in both STF and SSF, this leads
either the top
high resistance within the film plane. Metal current collectors were therefore deposited on
of the
of the thin films, or between the thin films and the YSZ substrates to facilitate lateral transport
were
electronic carriers for through plane measurements to evaluate ORR kinetics. Micro-patterns
fabricated using wideband ultraviolet photo-lithography. Patterns of different width/spacing (20, 40, 80,
with chrome
160 and 320 Rm) were drawn with AutoCAD (Figure 7) and a soda-lime glass base mask
spincoating was custom-ordered at Advance Reproductions Corporation (North Andover, MA). After
coating five drops of photoresist NR71-3000P (Futurex, Inc., Franklin, NJ) at 2,500 rpm for 30 seconds,
0
light for 40 seconds
samples were baked for two minutes at 160 C and exposed to broadband ultraviolet
Germany). The
through the chrome coated mask using a MA4 mask aligner (SUSS MicroTec, Garching,
Inc.).
samples were post-baked for two minutes at 105'C and developed for two minutes (RD6, Futurex,
to remove water
All samples were cleaned with oxygen plasma at 200W in 0.5 Torr for three minutes
molecules from the sample surface.
the
Figure 7. A schematic illustration of the current collector pattern. UV light passes through
polymer.
white region (transparent, not coated with chrome), cross-linking the photoresist
Photoresist on the black region (coated with chrome) dissolves in the developer.
21
Dense Pt or Au films were deposited on the photoresist patterned samples by DC sputtering (Kurt
J. Lesker, Clairton, PA). After reaching a base pressure of < 5x10-6 Torr using a cryogenic pump
(Cryotorr 8, CTI Cryogenics, Chelmsford, MA), DC power of 50 W was used to sputter Pt (10 mTorr of
Ar) or Au (10 mTorr with Ar:02 = 90:10) at room temperature. The lift-off process was completed by
soaking the samples in a solvent (MICROSTRIP@ 2001, FUJIFILM Electronics Materials U.S.A. Inc.,
North Kingstown, RI) at I 00 0C for ten minutes to completely dissolve the photoresist.
2.2 Physical Characterization
2.2.1
X-ray diffraction
X-ray diffraction (XRD) measurements for the powders and bulk samples were performed using a
PANalytical
X'pert
Pro
diffractometer
(Westborough,
MA)
in
Bragg-Brentano
geometry.
A
monochromated Cu Ka () = 1.541 A) X-ray source, Open Eularian Cradle (OEC) sample stage, and
PANalytical HighScore software package were used.
A Bruker D8 High-Resolution XRD (Madison, WI) with four-bounce Ge (022) incident beam
monochromator was used for thin film analysis. Alignment to the symmetric substrate (single crystal YSZ
with (001) orientation) peak was performed prior to each scan.
2.2.2
Density measurement
The density of the bulk samples was obtained by applying Archimedes' principle. The weight of
the samples in air and in water were measured, and the relative density of the bulk samples was calculated
to be approximately 97%.
2.3 Thermogravimetry
SSF bulk samples were broken into small pieces of at most few millimeters in each dimension,
and placed in an alumina crucible hanging on one side of a Cahn 2000 microbalance beam with a Pt wire.
22
ranges were used.
Depending on the amount of change in sample mass, 10 jg, 100 jig, and I mg recorder
Oxygen partial pressures were controlled by 0 2-N 2 (high pO2), H2-H 20-N 2 (low pO2), and CO-CO 2
controllers
(intermediate pO2) gas mixtures using MKS Instruments 1179A and M100B mass flow
sensor.
(MFCs) (Andover, MA). The oxygen pressure was monitored using an in situ Nernst based YSZ
2.4 Simultaneous conductivity and chemical expansion measurements
2.4.1
Bulk sample configuration
.
3
Rectangular bar-shaped SSF bulk samples were prepared with dimensions of 21 x2.7x2.47 mm
ensure
Four sets of Pt wires were wrapped around the bars as illustrated in Figure 8. Pt paste was used to
2.93 mm
good adhesion of the wires. The Pt pasted areas were approximately 3.17 mm wide and spaced
apart.
Pt Wires
'
Pt paste
Figure 8. A Schematic illustration of the SSF bulk sample with four-point probe setup for
simultaneous conductivity/chemical expansion measurements.
2.4.2
Measurement setup
The SSF bulk sample was loaded into Linseis L75 PT Horizontal Dilatometer (Robbinsville, NJ)
on
with quartz sample holder and push-rod. Electrical leads were introduced using electrical feedthrough
23
a vacuum port. The four Pt electrical leads were connected to HP 4192A impedance analyzer for AC
electrical characterization.
2.4.3
Measurement conditions
0
Conductivity and chemical expansion measurements were performed at 700-1000 C over a wide
range of oxygen partial pressure (10-20 ~ I atm) controlled by 0 2-N 2 (high pO2), H 2-H 20-N 2 (low pO2),
and CO-CO2 (intermediate pO2) gas mixtures with the aid of MKS Instruments MIOOB (Andover, MA)
MFCs. The oxygen pressure was monitored using an in situ Nernst based YSZ sensor.
2.5 Thin film EIS measurements
2.5.1
Cell configuration
Asymmetric cells were prepared by (a) depositing SSF thin films by PLD over the patterned Pt
current collectors on (001) oriented YSZ electrolytes, and by (b) depositing Au current collectors over
pulsed laser deposited SSF thin film on (001) oriented YSZ electrolytes. Both types of cells had porous
Ag applied to the other side of the YSZ to serve as a counter electrode as shown in Figure 9.
(a)
-N-
SSF
-
-
-
Buried Pt current collector
YSZ
(b)
SSF
Top Au current collector
YSZ
Figure 9. Schematic illustrations of asymmetric cells with YSZ electrolyte, SSF working electrode
with (a) buried Pt current collector and (b) top Au current collector. Both cells had porous Ag
counter electrodes.
24
2.5.2
Measurement conditions
Electrochemical impedance measurements were performed with a Modulab-MTS Test System
(Solartron Analytical, Hampshire, UK). Samples were held by metal clips that were wrapped with 99.99%
pure Pt wire (Alfa Aesar) and provided electrical contact. An AC amplitude of 10-20 mV was used. A
YSZ Nernst-type oxygen sensor, maintained at a constant temperature in a separate furnace connected to
the sample furnace, was used to monitor the oxygen partial pressure.
25
CHAPTER 3. RESULTS
3.1 Physical/chemical characterization of SSF powder and thin films
3.1.1
SSF powders
The XRD spectra of SrSnIxFe,03.6 (x = 0, 5, 35) powders are shown in Figure 10. All three
-
compositions are consistent with a single cubic perovskite phase. All peaks shift toward the right
smaller lattice constants - as Fe content increases.
SSO
SSF5
SSF35
SSF50
>1
(10)
30
(310)
(220)
(0)(222)
(211)
(200)
(100)
(2 10)
45
60
75
90
2 Theta (*)
Figure 10. XRD pattern of SrSni..xFe,03-S (x = 0, 5, 35) powders synthesized
3.1.2
SSF thin films
The XRD spectra of SrSnO3 (SSO), STF35 and SSF35 thin films grown by PLD on single crystal
YSZ (100) substrates are shown in the figures below. Films with low Fe content and/or low deposition
temperatures show very weak peaks or are XRD amorphous - i.e. exhibited no XRD peaks. In these
26
cases, peaks start to appear above the background noise when the deposition temperature is raised. XRDamorphous films at low deposition temperatures is a commonly observed phenomenon [17]. As the lattice
mismatch between the YSZ substrate and the thin films increases, the thin films become less crystalline,
i.e. weaker peaks. SSO thin films show almost no peak when pulsed laser deposited at 850'C heater
temperature (substrate temperature ~ 700'C) as shown in Figure 11, while SSF50 (Figure 12) and STF35
(Figure 13) thin films deposited at 850'C heater temperature showed strong peaks.
SSO_8500C
SSO 9500C
to
C
-1
-
-
--
Aj
A.
-A
0
YSZ (400)
YSZ (200)
SSO (110)1
SSO (220)
L11
-.- -.-.- L1
20
I
I
60
40
80
2 Theta (0)
Figure 11. HRXRD patterns of SrSnO3 thin films deposited on YSZ (100) single crystal substrates
by PLD at 8501C and 9500 C heater temperatures.
27
SSF50_8500C
SSF35_8500C
SSF35_9500C
(A
a,
4.'
0)
0
I
I
SSF (110)
I
YSZ (200)
-
SSF (220)
--
'-----
40
20
YSZ (400)
80
60
2 Theta (*)
Figure 12. HRXRD patterns of SSF50 thin films deposited on YSZ (100) single crystal substrates by
PLD at 850 0 C heater temperature, and SSF35 thin film at 850 0 C and 9500 C.
STF35_600 0 C
STF35_850*C
Mkh6ALAj.- I.L.,-u J-
_________1,____A,
____________I_____
,4 - , ,
II
I
I I
n,
I
,
I
STF (110) YSZ (200)
0)
0
STF (220)
20
40
60
YSZ (400)
80
2 Theta (*)
Figure 13. HRXRD patterns of STF35 thin films deposited on YSZ (100) single crystal substrates by
PLD at 6000 C and 8500 C heater temperatures.
28
The lattice parameters of the STF35 and SSF35 thin films were calculated from the position of
their peaks (Table 2). SSF35 has a larger lattice parameter than STF35 as the smaller cation, Ti", is
.
4
replaced by the larger cation, Sn
Table 2. Ionic radii of Ti4 and Sn"', lattice parameters and volumes of STF35 and SSF35 calculated
from XRD peaks.
Ionic radius
0.605 (Ti VI coord) [18]
0.69 (Sn4+ VI coord) [18]
Lattice parameter
3.906 A
4.002 A
3
59.592 A
Volume
64.114 A3
3.2 Thermogravimetric analysis
Oxygen nonstoichiometry 6 in SSF35 as a function of pO2 over the temperature range of
800-1000*C is plotted in Figure 14. Since the preferred valence state of Fe at intermediate pO2 is 3+, for
4
every two Fe 3 cations substituted onto the Sn B-site, one oxygen vacancy is created. Therefore, the
nonstoichiometry of oxygen can be written as 3 -
2
+ 6, where x is the Fe dopant concentration. This
suggests that SSF has a built-in oxygen deficiency given by x.
2
In oxidizing conditions (high oxygen partial pressure regime), SSF35 takes oxygen from the gas
phase increasing 6; as pO2 decreases, the oxygen vacancy concentration increases and 6 first decreases
and then changes sign and becomes increasing more negative as the material becomes net oxygen
deficient relative to 3 -
2
. The Temperature dependence
shows higher reducibility at higher
temperatures, noticeably under reducing conditions.
Compared to STF35, SSF35 is more readily reducible in oxidizing conditions, though the
difference decreases with increasing temperature (Figure 15). In reducing conditions, STF35 and SSF35
show almost identical oxygen nonstoichiometry
at 700*C, but as temperature increases, SSF35
nonstoichiometry shifts to higher pO2, i.e., lower 6. In oxidizing conditions, SSF35 shows much lower
29
0
oxygen excess than STF35 at 700 C, but the difference becomes smaller at 1000'C, as SSF35
nonstoichiometry shifts much slower to the right than the STF35 counterpart. In other words, the width of
the plateau for SSF shrinks much faster than STF as temperature increases.
SSF5 showed no measureable change in nonstoichiometry under both intermediate and high pO2.
Under reducing conditions, an irreversible phase transition occurred, confirmed by XRD, and the original
mass could not be recovered after reoxidation. Also, during the phase transition, the equilibrium took
significantly more time (days) compared to reversible nonstoichiometric changes (hours).
0.03
0.00
-0.03
0
A
700*C
800"C
*
9000C
1000*C
-0.06
-25
-20
-10
-15
log p
0 2
-5
0
(atm)
Figure 14. Oxygen nonstoichiometry - 6 - as a function of pO2 and temperature for SSF35.
30
m
0.06-
0
A
SSF35
STF35
SSF35
STF35
700*C
700 0C [191
1000*C
10000 C [19]
0.03-
0.00
-
CO
-0.03-
-
Increasing V0
-0.06-25
-20
-10
-15
log p 0
2
-5
0
(atm)
0
Figure 15. Comparison of oxygen nonstoichiometry of STF35 [19] and SSF35 at 7004C and 1000 C.
3.3 Bulk conductivity
The log of the bulk conductivity of SSF35 is shown plotted vs log pO2 in Figure 16 for a series of
0
isotherms ranging from 700 - 1000 C in one hundred degree intervals. The SSF35 exhibits p-type
conductivity with a slope of 0.254 0.017 under oxidizing conditions, a pO2 insensitive ionic region at
intermediate pO2, and n-type conductivity with a -0.180 0.007 slope in reducing conditions. This trend is
similar to that exhibited by STF35, as shown in Figure 17.
31
0
SSF35 conductivity
V
-1
V
~
AV,
_
-
-2
0
a
A
Man
-3-
v
-20
7000C
8000C
U
-15
-10
9000C
1000*C
0
-5
log p02 (atm)
Figure 16. Bulk conductivity of SSF35 as a function of PO2 and temperature.
0-
a
*
a
STF35 [201
SSF35
U
a
U
-1
a a 0
a a
-15
-10
1
0
-2-
-3-20
-5
0
log p02 (atm)
0
Figure 17. Comparison of bulk conductivities of STF35 [201 and SSF35 at 800 C.
32
3.4 Bulk chemical expansion
When an oxide material changes its oxygen nonstoichiometry and localized electrons are added or
subtracted, the radii of the cations change such that a noticeable contraction or expansion of the crystal
lattice occurs. The change in cation radii is much greater than the strain induced by oxygen vacancy
formation or annihilation, so the lattice expands as oxygen vacancies are formed and electrons are
to
localized to the cations, and vice versa [21-22]. This is known as chemical expansion, EC - analogous
thermal expansion, ET - and can be expressed as
EC =
(4)
acAS
where ac is the coefficient of chemical expansion (CCE), and AS is the change in oxygen nonstoichiometry. Since changes in oxygen nonstoichiometry become significant at higher temperatures
1.6%
-
(>400'C), chemical expansion is observed at those elevated temperatures (Figure 18).
1.2%
"AA-
0
S0.8%
CU
.
x
p
0.4%
chemical
0.0%
'--
0
200
600
400
800
1000
Temperature, *C
the
Figure 18. A schematic model of a material's expansion as a function of temperature, with
[221.
chemical expansion contribution noted. Reprinted from a reference
The SSF bulk sample was observed to expand systematically with decreasing oxygen partial
pressure, as illustrated in Figure 19. Figure 20 shows the chemical expansion of SSF35 as a function of
linear
oxygen nonstoichiometry (6) at 750'C as derived from the thermogravimetric analysis data. It has a
33
dependence on 6 with a slope (ac) of -0.0392 0.0006. It agrees well with the CCE of STF35, which
ranges from -0.0398 (700 0 C) to -0.049 (1000 C) [22].
224SSF35 900*C
.%
2
1% 02
-
222
2% 02
4%0O
E
-
220
20%4%
-J
410%02
218-
-
216
100
0
02
400
300
200
100
Time (minutes)
Figure 19. SSF bulk sample length in different oxygen partial pressures at a constant temperature
of 900 0 C measured by a dilatometer.
Chemical Expansion
-2-
*
750 0C
-3-
0
slope (CCE)
-0.0392
-4-
-5
-6-7 -I
0.008
0.016
0.012
6
0
Figure 20. Expansion of SSF35 as a function of 6 at 750 C.
34
0.020
3.5 EIS results of SSF thin film electrodes
Figure 21 shows a typical impedance spectroscopy spectrum of the asymmetric cell. There is an
offset resistance, independent of oxygen partial pressure, representing the ohmic resistance of the YSZ
electrolyte, and a highly distorted semicircle in the high frequency domain showing very similar
characteristics to porous Ag counter electrode, as studied by Di Chen in our group [23]. The large, nearly
ideal semicircle in the low frequency domain can therefore be attributed to the STF35 or SSF35 thin film
surface exchange kinetics. The offset resistance and the diameters of the distorted and nearly ideal
semicircles
show Arrhenius
behaviors. The diameter of the semicircle in the low frequency,
corresponding to the SSF35 or STF35 thin films, is normalized by the film area to give area specific
resistance (ASR) and plotted against temperature in Figure 22. The activation energies of STF35 and
SSF35 thin films are 1.70eV and 1.65eV, respectively.
-900-
-600-
w
-300-
0.
400
1200
800
1600
Z' (0)
Figure 21. Typical electrochemical impedance spectroscopy (EIS) spectrum of a SSF35 thin film
deposited onto a single crystal YSZ electrolyte with buried Au current collectors and porous Ag
counter electrode measured in air at 4681C.
35
3.5
ASR of SSF35 vs STF35 in air
3.0
E (STF35) = 1.70 eV
E 2.5-
C0
2.0-
Ea (SSF35) = 1.65 eV
0
*
0
1.50
U
I
*
I
.
1 .U
SSF35
STF35
1.15
1.20
1.25
1 .30
1.35
1.40
1000/T (K')
Figure 22. Comparison of temperature dependence of area specific resistance (ASR) of SSF35 and
STF35 thin film electrodes.
36
CHAPTER 4. DISCUSSION
4.1 Defect chemical model
4.1.1
Defect chemistry of STF and SSF
Oxygen nonstoichiometry was derived from the measurement of the relative mass change of SSF
as a function of oxygen partial pressure at different temperatures. As discussed in section 3.2, the
nonstoichiometry of oxygen in SSF or STF can be written as 3 -
2
+ 6, where x is the Fe dopant
concentration, and a built-in oxygen deficiency of x/2 due to the difference in valence states between Fe
and Sn or Ti. As a consequence, once they become occupied by insertion of extra oxygen, these normally
unoccupied sites are denoted as 'interstitials'. Anion Frenkel disorder occurs as a result, and the reaction
can be written in Kr6ger-Vink notation as:
ON + Vix <
V" + 0'
(5)
where Ox is an oxygen on an oxygen site with no charge with respect to the lattice, Vix is a vacancy on an
interstitial site (built-in oxygen deficiency) with no charge with respect to the lattice, V5 is a vacancy on
an oxygen site with two positive charges relative to the lattice, and 0' is an oxygen ion on an interstitial
site with two negative charges relative to the lattice.
The equilibrium constant can be written as
K',
af
[V6][O7']
[OK][yM]
(6)
Provided the change in [V] and [0'] remains small relative to [0x] and [Vix], the latter can be
assumed to remain approximately constant, and it becomes convenient to introduce a new constant K.1
given by,
if
0
Haf(7
Kaf = K'f[0x][Vix] = [V5][0'] = Kaf ekT
(7)
Intrinsic electron-hole pair generation can be written as:
nil +-* e' + h'
37
(8)
and the equilibrium constant Ki as
Ki = np = NcNve-EglkT
(9)
where n and p are the concentrations of electrons and holes in the conduction band and valence band,
respectively, and Nc and Nv are effective density of conduction band and valence band states,
respectively.
Similarly, the oxygen reduction reaction can be described by the following reaction and
equilibrium constant, Ked.:
-> V +2e' +O2
(10)
1
K'e
2
- [V]n (PO2)f
(11)
where n is the concentration of electrons in the conduction band.
As for the anion Frenkel equation above, provided the change in [V] is much less than [Ox], the latter
can be treated as a constant, and the defect equations simplified as done below.
Kred = Kred[Oox]
(12)
Equation (11) can be written as
[Vj]n 2 (pO 2 )f = Kred = [VS]on1 2 (P02 0)f
(13)
assuming [Ox] is much greater than [V ]O and [O'], where [V]O and ni are the concentrations at the
stoichiometric oxygen partial pressure P02 0 . Therefore, one can write
ni =
R
(14)
from equation (9).
Lastly, charge neutrality gives:
2[V ] + p = 2[0j'] + n
38
(15)
[V
Writing n as a function of other variables from equation (15) and substituting [0k"] and p as [v~lI and
from equations (7) and (9), respectively,
- [0']) + p = 2
=2([V-]
K4
Kaf'\
]+--
(16)
n
[Vd]/
Rewriting equation (13), one can write oxygen vacancy concentration as:
1
1
Ka2 K-(p02)21
1
[VI] =
2
n p0g
(17)
By substituting equation (17) into equation (16) and then solving the quadratic equation for n,
[V ]
(18)
K
Kaf +y
[V+]
[Vd]
Substituting equation (17) into equation (18) and solving for [Vs] and n using Mathematica 10.1,
analytical solutions for the defect concentrations as a function of thermodynamic parameters were found:
n=Root[4KafKiP0202
-
4
0
Ki P0 2 PO2 #1
2
+ (8KafKi2P0
2 PO2
0
4
+ 2Ki 3 p0PO2"0)#1
-0Root[4KaKi4p0
2
2"
-
4
0
2
Ki p02 pO2 #16 + 4Kafp0 2 2#1,]
0
0
Ki P02p02 #12 + (8KfKi2P0 2P0 2 + 2Ki 3 p02 P02 0 )#14
(19)
(20)
0
P02
KajK
IVi]J
-
2
0
- K P02P"2 #16
+ 4Kap02 2#1B]
where Root [ ] is the root of the polynomial of variable #1 inside the brackets. There are eight analytical
roots to the above function, and two of the roots have physical solution. The third root corresponds to the
solution for P02 0 > P02, and the fourth root for P02 0 < P0 2The oxygen nonstoichiometry
Kaf
(21)
S
-
-rs
-
[SSF]
can now be obtained analytically as a function of pO2, and is fitted to the thermogravimetry data using
fmincon function in MATLAB, as in Figure 14.
39
4.1.2
Model defect diagram
In different PO2 regimes, certain defect species dominate and approximations can be made. For
example, at low PO2, n and [V6] will be much larger than p and [O'], such that n ~ 2[V6] from the
charge neutrality equation. For the model systems SSF and STF, three different pO2 regimes and
corresponding approximations can be obtained, as shown in Figure 23. Along with the three key reactions
- anion Frenkel disorder, intrinsic electron-hole pair generation, oxygen reduction - and electroneutrality
equation, one can obtain analytical solutions for each defect species and their pO2 dependence, as
calculated in Table 3. A schematic defect diagram for SSF or STF model systems from these solutions are
shown in Figure 23.
Table 3. The predicted solutions to the defect solutions with Brouwer approximations.
I
1
( 2 Kred)P
II
1
P02
S1e1
Kred2
_
\
P0
1
2
2 aJ
1
Ka
P6
Kredz
2KafK 2 K 1P
Kred
)
( 2 Kred)3
K6K
red
K
1
Kaf 4
1
SPO26
K-
III
0
af P 2 6
[V1
KI
f
1
3
1Kre
QKred)
Kaf
0
6
Kaf
P 2_
Ka
40
2
K
K
2
1
(1Kred)3
4K
2
4Kred)
1KafKi1
PO26
v =-!o']
2[V0 ] =n
p
III
II
nI
2[O1']
1/6
1/6
M
U-4
[01
p
-00
~-1/6
1/6
Pi
10g P02
SSF and
Figure 23. Schematic defect diagram using Brouwer approximation for the model system
STF.
4.1.3
Defect diagrams and thermodynamic parameters of STF35 and SSF35
The carrier concentrations can be obtained as a function of pO2 (Figure 24) and the equilibrium
25 are the
constants as a function of reciprocal temperature (Figure 25). Slopes derived from Figure
reactions (Table 4).
thermal band gap and the enthalpies of the anion Frenkel and oxygen reduction
41
22.
E
0
20
STF35 [19]
-
SSF35
Vo
...------- Vo
0
a..
18 18-n
o
16-
-25
......
p
.-^--
- n-
-20
-15
-10
-5
0
5
log p02 (atm)
0
Figure 24. Carrier concentrations from defect equilibria at 1000 C.
SSF35's defect equilibria is shifted to the right (higher pO2) upon replacing Ti with the larger Sn
leading to a larger lattice parameter. This implies that SSF reduces more readily than STF. This is similar
to BaSrjTij-yFey03-y/21 (BSTF), where adding the larger Ba atoms for Sr shifts the defect equilibria to
the "right".
As expected, we find a larger band gap for SrSnO3 than for SrTiO3. However, the enthalpies of
anion Frenkel and reduction are sensitive to the conditions of the fitting, and the combined values of the
two enthalpies should be considered. SSF35 has a slightly larger enthalpy sum: this agrees with the
thermogravimetric analysis of STF35 and SSF35 where the difference in reducibility is greater at high
temperatures than lower temperatures.
42
140
m In(Kaf)
*
In(Ki)
In (Kred)
120-
100-
80-
9.0
10.0
9.5
10.5
11.0
I/kT (1/eV)
Figure 25. Natural log of equilibrium constants K as a function of 1/kT.
Table 4. Thermodynamic parameters for STF35 [191 and SSF35 from fitting the nonstoichiometry
6 as a function of oxygen partial pressure at different temperatures.
STF35[191
SSF35
AHaf (eV)
0.518
0.489
Eg (eV)
1.379
1.772
AHred (eV)
3.893
4.304
Fittings of nonstoichiometry 8 were very insensitive to the equilibrium constant of anion Frenkel
disorder. Therefore, further analyses on both STF35 and SSF35 were conducted by fixing He, from that of
STF35 in reference [19]. This time, data points from all temperatures were fitted simultaneously by
setting enthalpies and pre-exponential terms as variables, instead of fitting equilibrium constants at each
temperature. The pre-exponential term for the anion Frenkel disorder from the STF35 fitting was fixed as
a constant for fitting SSF35 data. The resulting enthalpies and pre-exponentials are shown in
. Note that Hred for both fittings were almost identical and Eg was within 3% of each other for SSF35. This
suggests that the two different fittings are consistent.
43
-
Table 5. Thermodynamic parameters for STF35 and SSF35 from fitting nonstoichiometry - 8
data as functions of oxygen partial pressure over a range of temperatures simultaneously.
STF35 (data points from [19])
SSF35
tHaf (eV)
0.518
0.518
K f
99.16
99.16
Eg (eV)
1.287
1.718
NcNv
AHred (eV)
99.63
3.758
103.28
4.302
Ke
159.95
166.74
4.2 Conductivity model
Four-point impedance measurements of SSF bulk samples showed a single semicircle in the
frequency range of the instrument (HP 4192A, 5 Hz ~ 13 MHz). This is equivalent to a parallel RC
circuit, where the resistance value can be used to calculate the electrical conductivity taking into account
the geometry of the device. The PO2 independent ionic conductivity can be subtracted from the total
conductivity in Figure 16 and n- and p-type conductivities can be extracted, as shown in Figure 26 and
Figure 27. With subtraction of the background ionic conductivity, now the PO2 dependence of both the ptype and n-type conductivities have slopes of 0.25 0.01 (negative for n-type) in log-log plots, as shown in
Figure 27.
44
0
-1-
Hole
Electron
Ion
A
*
*
-20
-3I
I
I
0.75
0.80
0.85
0.90
I
-
-
I
1.00
0.95
1.05
1000/T (K-')
Figure 26. Temperature dependence of p-type (pO2= 0.316 atm), n-type (pO2
pO2-insensitive ionic conductivities of SSF35.
-0.
10-18 atm) , and
(b)0
5
,
(a)
=
-10-0.6
A__
-2.0-
L
-"S
o -2.5-
01
m800*C
-1.0
700'C
e
-
*
0
-
-3
-0.8
.
-3.5
-22
-20
a
a
-18
i
16
A
V
900*C
1000*C
,
i
*
A
V
-1.2
-2.5
-14
-2
0
.
i i
-1.5
4
i
-1.0
05
-0.5
700*C
800C
900*C
1000*C
0.0
0.0
0,5
log p02 (atm)
log p02 (atm)
Figure 27. Oxygen partial pressure dependence of (a) n-type and (b) p-type conductivities of SSF35
extracted by subtracting the respective ionic conductivities at each isotherm.
The conductivity of a material is defined by
f=
45
niqiqi
(22)
where i refers to a charge carrier (electron, hole, or oxygen ion), n, is the concentration, q, is the charge
and p, is the mobility of the carrier. Since the charge carrier concentrations were obtained from the defect
equilibrium model, the conductivity dependence on temperature and pO2 can be predicted by assuming
carrier mobilities. Conversely, carrier mobilities can be calculated from the combination of the
conductivity data and the defect model, as shown in Table 6 and plotted in Figure 28.
The electron mobility of SSF35 (~0.0002 cm 2 /V-s) is less than a quarter that of STF35 (-0.0009
cm 2 /V-s [19]). This suggests that the conduction band of SSF35 is dominated by the Fe 3d derived band,
rather than the higher mobility Sn 5s derived band. STF35 seems to have Fe 3d orbital hybridized with Ti
3d, resulting in higher mobility than the sole Fe 3d derived band in SSF35. However, SSF35 (Table 6)
and STF35 (0.005 cm2 /V-s [19]) show similar hole mobilities; this can be explained by similar valence
band structure - Fe 3d and 0 2p hydbridization - in both SSF53 and STF35.
Table 6. Electron and hole mobilities of SSF35 calculated from defect chemistry and conductivity
data.
T (*C)
P. (cm 2/V-s)
pp (cm 2 /V-s)
700
0.000252
0.00238
800
900
1000
0.000152
0.000215
0.000204
0.00291
0.00429
0.00523
46
-2.0
p-p
P-fl
-3.0
-
c
E
0)
0
-
-3.5-
U
-4,0
0.75
0.95
0.90
0.85
0.80
1.00
1.05
1OOO/T (K)
Figure 28. Temperature dependence of electron and hole mobilities of SSF35.
The activation energies of each type of conductivity can be extracted from the slopes in Figure
26. These activation energies are the sum of the charge carrier concentration and mobility (migration)
activation energies. Therefore we can separate the contributions to the activation energy contributed by
the mobility and carrier concentration (Table 7).
Let us look at the electronic part first. From equations (11) and (12),
(
1
1
Kredf(PO2 )
n=
KI(23)
At intermediate pO2, [Vs] varies with temperature as -f-, therefore the activation energy related
to n generation will be approximately
Hred
2
-
2
. At low pO2, n ~ 2[V ], so:
1
n = ( 2 Kred)jPO2
1
and the activation energy related to n generation will be approximately
(24)
3
. Reducing conditions at
which n-type conductivity measurements were done were near the left edge of the ionic conduction
regime, but well within the regime in which the vacancy concentration dominates. This is consistent with
47
1
the P0 2 - dependence of the n-type conductivity. Therefore the activation energy for electron generation,
taking into account the effective activation energy for electron mobility of -0.040 eV, is 1.872 eV which
should be given by 4.262 eV as shown in Table 7. This
Thus Hred
2 .a.
-
should be equivalent to
value is within 0.040 eV of the value given for
Hred
in Table 5.
Table 7. Activation energies of n-type conductivity, electron mobility, electron concentration, and
the reduction enthalpy.
Eaayn
Eapn
Ea,n
Hred
1.833 eV
-0.040 eV
1.872 eV
4.262 eV
For holes, we have
K i Ki [V"](p0 2 )f
- 1
(25)
Kred2
At intermediate pO2, the activation energy of p will be given by Eg +
p
2
-af
Hred
2
At high pO2,
2[0'], so:
p= (Kr
iP02
(26)
KrafEgHe
and the activation energy of p will be approximately
the p-type conductivity, Ea, is given by Eg +
Again, given the P 0 2$ dependence of
Ha3+Eg-Hred
2 . Taking into account the derived value for
H
Hred= 4.262 eV, Haf = 0.518 eV, and Eg = 1.718 eV, one calculates a value for Ea,p = -0.154 eV.
This value is 0.091 eV smaller than that shown in Table 8 derived from the electrical data and assuming
an activated hole mobility with activation energy Ea,jp = 0.292 eV.
Table 8. Activation energies of p-type conductivity, hole mobility, hole concentration, and
activation energies of hole concentration at intermediate and high pO2.
Eagr
Eap,
Ea,p
0.047 eV
0.292 eV
-0.245 eV
48
An activation energy for a nonstoichiometric oxide semiconductor of close to zero, i.e. Ea,ap =
0.047 eV, is highly unusual. However, a similar observation was already made previously for the related
STF35 system. There it was shown that with the Fermi energy close to the valence band, Ea,p took on a
small positive value, while presumed phonon scattering resulted in a small effective negative value for
Ea,Up which served to nearly compensate Ea,Mp to result in a near zero value for Ea,,p[2 0]. This is likely
the case here as well. In subsequent work, I will re-examine the fitting routines to see where small errors
may have introduced themselves to result in a reversal of signs of the two terms contributing to EaapFor the ionic component, the activation energy for oxygen vacancy formation was calculated
from the defect model of SSF, i.e.
= Ea,[ve, while the migration enthalpy Hmig was obtained by
subtracting that activation energy from the total ionic conductivity activation energy, as shown in Table 9.
Table 9. Activation energies of ionic conductivity, oxygen vacancy concentration, and enthalpy of
oxygen vacancy migration.
Eaaion
Ea,[v, I
0.947 eV
0.245 eV
Hmig
0.702 eV
The calculated migration enthalpy of SSF35 is in the same range as other MIEC perovskite
oxides (0.5~0.9 eV) [24].
4.3 Chemical expansion
When a Fe cation changes its states between Fe2+, Fe'+, and Fe 4 +, its ionic radius changes
significantly. As SSF is reduced, extra electrons are created and localized on the Fe ion, decreasing its
oxidation state (i.e. from Fe4 to Fe' to Fe2+). SSF35 has very similat CCE to that of STF35. This is
expected as electrons are localized on the'Fe cations, and the mechanism is essentially identical for both
STF35 and SSF35. Both STF35 and SSF35 have a CCE nearly an order of magnitude lower than that of
49
the fluorite oxides, and therefore will induce much less stress on the electrolyte under reducing conditions
[21].
4.4 Investigation of surface exchange kinetics of STF/SSF thin films
Although some bulk properties differ from that of the thin film due to strain from lattice
mismatch against a substrate, bulk trends are good indicators in predicting the corresponding properties in
thin films.
The surface exchange coefficient k is inversely proportional to the area specific resistance (ASR)
of STF/SSF thin films as [12,25]:
k =kBT
4e 2 Rsco
(27)
where kB is the Boltzmann constant, Rs is the area specific resistance, and co is the total concentration of
lattice oxygen. A strong correlation between k and electron carrier density, n, was found in a previous
study (reviewed in section 1.3.1). SSF35 has slightly larger n derived from defect equilibria, but a lower
electron mobility. Interestingly, SSF35 showed slightly better surface exchange kinetics, and nearly the
identical activation energy of the ASR (or k) (Figure 22). This is not inconsistent with the hypothesis that
the electrons in the conduction band are rate limiting in the oxygen surface exchange reaction.
4.5 Problems with SSF5 and SSF50
4.5.1
Phase decomposition of SSF5
Thermogravimetric measurement of SSF5 in reducing conditions (mixture of wet hydrogen and
nitrogen) resulted in phase decomposition of SSF5. Therefore the instability of a lower valent Sn seems to
be the source of this instability. The addition of Fe appears to stabilize the stannate to lower PO2.
50
4.5.2
Degradationof SSF50 thinfilm electrode
Degradation of the SSF50 thin films occurred noticeably faster than that of SSF35. Degradation
of thin films during EIS measurements made it impossible to obtain meaningful impedance spectra
(significantly distorted).
51
..........
CHAPTER 5. CONCLUSION
5.1 Summary
The electro-chemo-mechanical properties of a new SOFC model cathode material SrSnl.xFe03.
xI2+n
have been explored in this study. Defect chemistry and conductivity models of SSF35 were
established with key thermodynamic and kinetic parameters derived, namely the enthalpy of reduction
(4.302 eV), electronic band gap (1.718 eV), migration enthalpy of oxygen vacancies (0.702 eV), and
2
electron (0.0002 0.00005 cm 2 /V.s) and hole (0.0037 0.0015 cm /N-s) mobilities.
The SSF defect equilibria was distinctly different from that of STF largely under oxidizing
conditions at the lower temperatures of this study for which the regime of oxygen excess was shifted to
higher pO2. This could be attributed to its higher band gap energy. The n-type electrical conductivity of
SSF35 followed the same trend as its STF counterpart, but with lower magnitude, suggesting that the Fe
3d band is largely responsible for electron conduction in the conduction band rather than the Sn 5s band
(Figure 29). STF35 showed higher electron mobility, possibly due to the Ti 3d derived band hybridizing
with the Fe 3d derived band, while only Fe 3d derived band is conducting electrons for SSF35. The hole
mobility of SSF35 was almost identical to that of STF35, as expected, since Fe 3d and 0 2p orbitals
constitute the valence bands of both SSF35 and STF35.
E--E-
Conduction band
EC
EF
EC
EE
F
--- -- --
EV
Ev
Valence band
x in
Tij Fex)03-x/2+
x in S (SnFe )O 3-x/,
6
Figure 29. Expected conduction and valance band positions of STF and SSF as a function of Fe
content, x.
52
5% Fe doped SrSnO3 (SSF5) showed low chemical stability under reducing conditions. In order
to be used as an electrode for a reversible SOFC, stability in both oxidizing and reducing conditions must
be fulfilled, and thus a sufficient Fe concentration is required for SSF to serve as a potential electrode.
However, too much Fe content in SSF50 resulted in fast degradation, possibly from Sr segregation to the
surface. Sr, due to its large size, is believed to segregate to the surface as a result of compressive stress
[26]. Because Fe is smaller than Sn, higher Fe content decreases the lattice constant, increasing the
compressive stress on Sr and the degree of its surface segregation. Therefore, the SSF35 composition
seems to be the most suitable in Fe-doped SrSnO3 perovskite-structured oxide as a SOFC electrode. The
chemical expansion of SSF35 was found to be similar to STF35 and other perovskite oxides. The surface
exchange rate and its activation energy for SSF35 were similar to those of STF35. SSF35's higher
electron density in the conduction band seems to be compensated by its lower electron mobility in terms
of the magnitude of the n-type conductivity.
5.2 Recommendations for Future Work
The electro-chemo-mechanical properties of the perovskite-structured SSF have been explored;
however, this study has also raised a number of questions. Instead of the expected large increase in
electron mobility associated with the Sn conduction band, an electron mobility even lower than in STF35
was obtained. A more detailed examination of the energy band structure of SSF35 could lead to an
improved insight into this question.
The activation energy and magnitude of the surface exchange coefficient of SSF were found to be
very similar to those for STF. This is likely due to the fact that Sr segregates to the surface of both while
both also depend on Fe for charge transfer. It would thus be interesting to study surface segregation in
SSF and see how it compares with that in STF [26]. Surface exchange kinetics were found to be strongly
dependent on surface termination in metal oxides, such as Sn02 [27]. Surface segregation has been
53
considered as one of the major issues for SOFC cathodes at high temperatures and in long term operations
[26,28-30]. Systematic variations in surface by surface decoration or etching and surface analysis such as
X-ray photoelectron spectroscopy can be employed to further study these phenomena.
Surface exchange kinetics can be further studied with the aid of optical transmission spectra. A
case study of Pr-doped ceria was demonstrated by characterizing changes in optical absorption of
different valence states of Pr [31]. Fe cations also have different optical absorption spectra for different
valence states, and so a similar technique can be employed. Unlike AC impedance spectroscopy with their
need for metallic current collectors, optical measurements do not affect or interfere with the surface
exchange kinetics, and therefore a direct, operando measurement is possible.
Relaxation in both bulk conductivity and chemical expansion were examined following changes
in PO2. However, analysis of the data was complicated by the porosity of the bulk samples. This resulted
in higher than expected diffusion coefficients, given much reduced diffusion lengths than the outside
dimensions of the specimens would suggest. The use of thin films is recommended for in-plane
conductivity relaxation using interdigitated electrodes, as well as chemical expansion relaxation using
high-temperature XRD (HTXRD), which can be done simultaneously to systematically compare the two
methods - electrical and mechanical - on extracting surface exchange and diffusion coefficients. In
addition, by direct measurement of the kinetics, optical relaxations studies can help complete the broader
picture for this model perovskite oxide.
54
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