The Effect of Barium Non-Stoichiometry on the Phase Structure,
Sintering and Electrical Conductivity of BaZr0.7Pr0.1Y0.2O3-ẟ
by
Kaamil Ur Rahman Mohamed Shibly
A thesis in partial fulfilment
of the requirements for the degree of
Master of Science
in
Material Science and Engineering
King Abdullah University of Science and Technology, Thuwal,
Kingdom of Saudi Arabia
05 May 2015
The thesis of Kaamil Ur Rahman Mohamed Shibly is approved by the examination
committee.
Committee Chairperson:
Professor Enrico Traversa
Committee Member:
Professor Enzo Di. Fabrizio
Committee Member:
Professor Alexander Rothenberger
ii
© 05 May 2015
Kaamil Ur Rahman Mohamed Shibly
All Rights Reserved
iii
ABSTRACT
The Effect of Barium Non-Stoichiometry on the Phase Structure, Sintering and Electrical
Conductivity of BaZr0.7Pr0.1Y0.2O3-ẟ
by
Kaamil Ur Rahman Mohamed Shibly
This thesis attempts to test the effects of barium non stoichiometry and varying
calcination temperatures on the microstructure and electrical conductivity of
BaxZr0.7Pr0.1Y0.2O3- ẟ (x = 0.9, 1.0, 1.1). BZPY powders were fabricated using a
combustion method, with the quantity of barium carefully controlled to create powders
with a 10% molar excess or deficiency of barium. Then, portions of the precursor were
calcined at 900 ºC, 1000 ºC, 1100 ºC, 1200 ºC and 1300 ºC for 5 h. The resulting
calcined powders were pressed into pellets and sintered at 1600 ºC for 10 h, in a powder
bath of the same chemical composition. In all, three chemically different powders were
synthesized, and each composition was subjected to five different calcination
temperatures, resulting in fifteen different samples to characterise.
The precursor from the combustion method was characterised by using an STA to
perform both TG and DSC simultaneously. The chemical composition of the precursor
and calcined samples was analysed using ICP-OES. XRD was used to characterise the
phases of both the powders and the sintered pellets. Lattice parameter indexing using
Topaz and Scherrer's equation were used to extract the lattice parameters and crystallite
sizes respectively. The microstructure of the pellets was examined using an SEM, the
grain size measured using a linear intercept method and pore size using ImageJ. Finally,
EIS was used to measure the conductivity of the pellets in dry and wet Argon
atmospheres, with silver electrodes.
Unfortunately, neither changes to barium stoichiometry nor partial calcination could
improve the performance of BZPY. Partially calcined samples did not give rise to dense
pellets, barium deficient samples showed inferior conductivity and barium excess
samples, while showing higher conductivity than the barium deficient pellets at high
temperature, were fragile and had to be handled carefully.
Ultimately, the attempt to improve the performance of BZPY did not succeed and
alternate methods of improving the grain growth need to be sought.
iv
ACKNOWLEDGEMENTS
This dissertation would not be possible without the help of several individuals who were
generous with both their time and their thoughts.
I would like to express my greatest appreciation to my advisor, Professor Enrico
Traversa, for his support and forbearance during my time under his supervision. It was a
challenge and eye-opening experience to engage in this research project under his
guidance.
My sincerest thanks to Dr. Bin Lin for his continuous tutelage and assistance during the
time of my thesis and for helping me understand many practical aspects in the laboratory
and for sharing with me many helpful and time saving techniques.
Wandi Wahyudi has my gratitude for his friendly and helpful demeanour and for always
being quick to help or encourage me during difficult times.
Lastly, I am grateful to my family members for their unending love and support during
my life and for encouraging me to pursue my graduate studies.
v
TABLE OF CONTENTS
1. Introduction ............................................................................................................... 1
1.1 Definition of a Fuel Cell ..................................................................................... 2
1.2 Historical Perspective of Fuel Cells .................................................................... 3
1.3 The SOFC Structure ............................................................................................ 5
1.3.1 Electrolyte ..................................................................................................... 5
1.3.2 Anode............................................................................................................ 7
1.3.3 Cathode ......................................................................................................... 8
1.4 Proton Incorporation ........................................................................................... 8
1.5 Mobility of Protonic Defects ............................................................................. 10
1.6 High Temperature Proton Conducting Oxides .................................................. 12
1.7 Proton Conducting Electrolyte Materials .......................................................... 12
1.8 Improved Electrolyte Performance ................................................................... 14
2. Experimental Methods ............................................................................................ 17
2.1 Synthesis............................................................................................................ 17
2.2 Simultaneous Thermal Analysis (STA) ............................................................ 17
2.3 Scanning Electron Microscopy (SEM) ............................................................. 18
2.4 Archimedes Method .......................................................................................... 18
2.5 X-Ray Diffraction ............................................................................................. 19
2.6 Electrochemical Impedance Spectroscopy ........................................................ 19
2.7 Induction Coupled Plasma - Optical Emission Spectroscopy (ICP-OES) ........ 20
2.7.1 Instrumentation ........................................................................................... 20
2.7.2 Reagents...................................................................................................... 20
2.7.3 Digestion of the samples and determination of the elements ..................... 21
4. Results and Discussion ........................................................................................... 22
4.1 Chemical Composition ...................................................................................... 22
4.2 TGA and DSC ................................................................................................... 23
4.3 XRD .................................................................................................................. 24
4.4 SEM................................................................................................................... 31
4.5 Proton Conductivity .......................................................................................... 35
5. Conclusion .............................................................................................................. 37
6. Bibliography ........................................................................................................... 39
vi
LIST OF ABBREVIATIONS
AFC -
Alkaline fuel cell
BCY -
Yttria doped barium cerate
BCZY -
Yttria and zirconia co-doped barium cerate
BZPY -
Yttria and praseodymium co-doped barium zirconate
BZY -
Yttria doped barium zirconate
DMFC -
Direct methanol fuel cell
DSC -
Differential Scanning Calorimetry
H-SOFC -
Proton conducting solid oxide fuel cell
ICP-OES -
Induction Coupled Plasma - Optical Emission Spectroscopy
MCFC -
Molten carbonate fuel cell
MIEC -
Mixed ionic and electronic conductor
O-SOFC -
Oxygen ion conducting solid oxide fuel cell
PAFC -
Phosphoric acid fuel cell
PEMFC -
Proton exchange membrane fuel cell
QM -
Quantum molecular dynamics
SOFC -
Solid oxide fuel cell
STA -
Simultaneous Thermal Analysis
TGA -
Thermogravimetric Analysis
TPB -
Triple phase boundary
XRD -
X-Ray Diffraction
YSZ -
Yttria Stabilised Zirconia
vii
LIST OF FIGURES
Figure 1: Schematic representation of a proton conducting fuel cell........................................3
Figure 2: Instantaneous proton transfer configurations obtained from quantum molecular
dynamics (QM) simulations. ...........................................................................................11
Figure 3: Calculated proton conductivities based on proton concentrations and mobilities for
varoious oxides. Conductivities of oxides with perovskite-type structure are shown by bold
lines. ................................................................................................................................14
Figure 4: a)The Thermogravimetric results of the different pre-calcined powders. b) The DSC
result ............................................................................................................................... 23
Figure 5: The XRD results for powders calcined at different temperatures for 5h for a)B0.9ZPY,
b)B1.0ZPY and c)B1.1ZPY ............................................................................................27
Figure 6: Crystallite size of fresh powders. .............................................................................28
Figure 7 Lattice Parameter vs Calcination Temp. for: left: Fresh Powder and Right: Sintered
Pellets ..............................................................................................................................29
Figure 8: The XRD of the sintered pellets, using different calcination temperatures. All pellets
were sintered at 1600 ºC for 10h. ....................................................................................30
Figure 9: SEM pictures of the as fractures cross-sections of sintered pellets. .........................32
Figure 10: a)Relative density of pellets relative to sintered BaCO3. b) Open porosity % as
determined by the Archimedes Method. .........................................................................33
Figure 11: a) Average Grain Size and b)Average Pore Size as determined from SEM images.34
Figure 12: a)All stoichiometries in wet argon, b) B1.0ZPY in wet argon and c)B1.1ZPY in wet
argon................................................................................................................................35
LIST OF TABLES
Table 1: Data for different types of fuel cells. ...........................................................................4
Table 2: ICP results showing the composition of the synthesised (fresh) powders. Ash refers to
the powder produced after combustion and before calcination. ......................................22
viii
CHAPTER 1: INTRODUCTION
The worldwide energy demand is expected to continue growing at an annual rate of 1.8%
through 2030 in response to population growth and worldwide industrialization.
According to the U.S. Energy Information Administration, Fossil fuels such as
petroleum, coal and natural gas supply 83% of the of the worlds energy[1].
Unfortunately, fossil fuels reserves are concentrated in a few regions of the world while
the demand for fossil fuels is worldwide. To make matters worse, fossil fuels are finite
and are being depleted at an alarming rate. According to the Energy Watch Group
(EWG) 2013 report, fossil fuel production is expected to peak at 2020 and thereafter
begin to decline[2]. The supply of crude oil is expected to last for only 40 - 50 years
more. To minimise the danger of disruption and price volatility, alternate fuel sources are
necessary to maintain energy security and prosperity.
One way to avoid this danger is to pursue research in technologies which can produce
electricity from hydrogen, since it can be produced locally by most nations. Depending
on how the hydrogen is generated, this method can help mitigate green house gas
emissions - a critical advantage as global warming becomes an ever growing issue. Solid
oxide fuel cell systems are promising candidates for next generation energy conversion
and offer the additional benefit of being scalable, with SOFCs ranging from small
portable power generation to large centralized power stations.
1
1.1 Definition of a Fuel Cell
A fuel cell is an electrochemical device which converts chemical energy from a
fuel into electricity through an electrochemical reaction with oxygen or another
oxidizing agent [3]. SOFCs can achieve an efficiency as high as 60%, much higher than
that of internal combustion engines which usually only reach 25% efficiency.
Additionally, SOFCs can also produce thermal energy, which, when used with cogeneration can further increase the system efficiency to 75%. Since the SOFC efficiency
is independent of its size, it can be made highly modular, making it easy to scale up and
to be placed in different locations. The most common fuel used is Hydrogen, but
other hydrocarbons such as methanol may also be used.
The rudimentary structure of a proton conducting SOFC is shown in Figure 1. The
typical fuel cell consists of a porous cathode, porous anode and a dense
electrolyte. The electrolyte is an electronic insulating ion conducting ceramic which
blocks the passage of electrons but conducts oxygen ions (Oxygen Conducting
Membrane) or hydrogen ions (Proton Conducting Membrane). The cathode and
anode are Mixed Ionic and Electronic Conductors (MIEC).
Because the electrolyte is an electronic insulator, electrons from the chemical
reaction need to go from anode to cathode by an external circuit, which provides the
power to run electronic devices. As shown in Figure 1, in the production of electricity,
fuel and oxidative gases flow along the surface of the anode and cathode, respectively,
and react electrochemically at the triple phase boundaries (TPB) where the electrode,
2
electrolyte and pore meet.
Therefore, the consistent generation of electricity can be
ensured when fuel and oxidative gases are constantly fed to the porous electrodes.
Normally, a single cell produces a voltage around 1.0 V and a series connection
between cells is needed for obtaining higher voltage for actual use.
Figure 1: Schematic representation of a proton conducting fuel cell.
Based on this concept, many different types of fuel cells have been developed. Fuel cells
can be classified based on different parameters such as operating temperature or type of
fuel used, but the literature predominantly classifies fuel cells based on the electrolyte
material used. Table 1 lists information on different types of fuel cells.
1.2 Historical Perspective of Fuel Cells
Sir William Grove first operated a successful hydrogen-oxygen cell in 1839. He built a
cell in which hydrogen gas served as a fuel, oxygen as the oxidative gas and dilute acid
solution served as the electrolyte. The fuel cell produced both electricity and water.
German scientist and Nobel laureate, Hermann Walther Nernst, invented the first solid
electrolyte in 1897. He demonstrated that zicronia-based solid oxides can have very
3
high electrical resistance and, when electric current is applied, can be used for
lighting purposes [5].
Table 1: Data for different types of fuel cells [4].
Baur and Preis demonstrated the first working SOFC in 1937. Coke and magnetite were
used as fuel and oxidant with zirconia ceramics used as the electrolyte. The fuel cell had
an operating temperature of 1273 ˚K. Professor Davtyan, a Russian scientist, attempted
to improve the performance of zirconia ceramics by mixing them with several
carbonates, but still experienced low power output. This limited the practical feasibility
of the technology at the time [6]. Fortunately, research in solid oxide technology began
to accelerate in the late 1950's and began spreading across the world to various
laboratories.
SOFC research intensified in the 1960s, driven by the exacting demands of space
programs and military applications. Most of the research at the time focused on
4
improving the conductivity of the electrolyte. A 'second wave' of fuel cell research began
in the mid 1980's and continues to the present day, focusing on developing new materials
and their technological applications in devices. In recent years, more and more
companies, universities and national laboratories are engaging in SOFC research due to
the growing concerns of energy security and rising energy prices [7].
1.3 Structure of the SOFC
1.3.1 Electrolyte
Solid oxide fuel cells are usually operated at temperatures between 600 and 1000 ˚C,
depending on the design of the cell configurations. A high operating temperature is
required for the SOFC to increase the ionic conductivity of the electrolyte to generate
usable power output. The high operating temperature also gives the additional benefit
of not needing precious metals, such as platinum, for electrocatalysis, and the
possibility of internal hydrocarbon reforming at the anode, which means
hydrocarbon fuels such as methane (natural gas), ammonia and propane can be used
as fuel directly.
Many SOFCs can be classified into oxygen ion conducting SOFC (O-SOFC) and proton
conducting SOFC (H-SOFC), depending on the nature of the electrolyte.
5
YSZ is the most commonly used electrolyte material for O-SOFC. The advantages of
using YSZ include [8], [9]:
1. High strength and toughness, which provide good fracture resistance when
serving as the electrolyte in a SOFC.
2. High oxygen ion conductivity at temperatures higher than 700˚C, 0.052 S cm-1 at
800˚C and 0.18 S cm-1 at 1000˚C.
3. High chemical stability and pure oxygen ion conductivity.
4. Good sinterability results in high relative density, which is required in order to
serve as an electrolyte material in an SOFC.
Unfortunately, YSZ requires a high operating temperature in the range of
800 - 1000 oC. This high operating temperature is a major drawback for SOFC practical
applications because of the need for expensive materials, long startup times, large energy
inputs to warm up the cell and the occurrence of thermal stresses that reduce the cell
lifetime.[10]
Reducing the operating temperature would help overcome some of these
limitations. However, this leads to a significant drop in performance of the SOFC
due to the increased resistance in the electrolyte and electrodes. [11]
It was only in the late 1960s that proton conducting ceramics were discovered. The first
systematic investigation was performed in 1980. As proton conducting ceramics are a
newer discovery, there is far less history and research than what there is for YSZ.
6
Theoretical calculations indicate that using proton conducting ceramics as the electrolyte,
instead of oxygen ion conducting ceramics is advantageous.
These advantages include [12]–[14]:
1. H-SOFC produces steam at the cathode, thus preventing fuel dilution at the
anode. This makes the complete utilisation of fuel highly plausible.
2. The activation energy of proton conductivity is lower than that of oxygen ions.
This is to be expected, given the fact that protons are significantly smaller than
oxygen ions.
3. Higher conductivity, which means H-SOFC can be operated at a lower
temperature such as 500 - 700˚C.
1.3.2 Anode
Most SOFC anodes consist of a mixture of the electrolyte material and a metal oxide,
with the metal oxide being subsequently reduced to metal under reducing atmosphere
prior to operation. Not only does the mixing of the two materials help inhibit the
sintering of the metal particles, it also reduces the difference in thermal expansion
coefficient between the anode and the electrolyte. Unlike the dense electrolyte, the anode
structure is deliberately fabricated with a high porosity (30% - 50%) to facilitate gas
transportation within the anode. Nickel is the most common anode material given its
catalytic activity under operating temperatures and its chemical stability and low cost
[15].
7
1.3.3 Cathode
The cathode material needs to be both an oxygen ion and electronic conductor under
oxidative atmosphere, while having a high catalytic activity for oxygen molecule
dissociation or reduction. La1-xSrxMnO3(LSM), is one of the most commonly used
cathode materials and is a p-type semiconductor with a perovskite crystal structure. The
ionic conductivity is seen to increase if the transition metal Mn is substituted with Co
[20]. However, the doping of Co also leads to higher mismatch of thermal
expansion coefficient with electrolyte materials. Some other perovskite materials such
as lanthanum strontium ferrite, samarium strontium cobaltite and n-type semiconductors
also show good electrocatalysis.
1.4 Proton Incorporation
Since the material under investigation in this thesis is a proton conducting electrolyte, it
would be useful to briefly examine how proton incorporation occurs in these electrolyte
materials.
The most important reaction for perovskite ceramics leading to proton uptake is
the dissociative absorption of water, which requires the presence of oxygen ion
vacancies. Perovskite structure ceramics such as BaCe0.8Y0.2O2.9 have been shown
to be good protonic conductors. [15], [16] The Y3+ ion substituting on the perovskite B
sites for Zr4+ require an O2- ion vacancy for every two Y3+ ions to maintain charge
neutrality, if there are no protons present. The reaction can be written by Kroger-Vink
notation as:
8
(1)
From a valence point of view, when Y3+ ions replace some Zr4+ ions, an
oxygen with a Y3+ neighbour should welcome one proton to keep the neutrality of
the structure. If protons are added by exposure of the ceramic to hydrogen or steam,
charge neutrality requires that one O2- ion be added for every two protons added
and can be presented as:
(2)
(3)
Since oxygen vacancies are lost when protons are added into the structure, the proton
concentration is limited to the same value as the dopant concentration. In
BaZr0.8Y0.2O2.9, this is 0.2 protons per formula unit, if it is assumed that oxygen ions can
only be added until all oxygen ion vacancies are filled. This limitation of proton
concentration not usually exceeding Y+3 concentration has been validated
experimentally. Schober and Bohn found a maximum proton concentration of 0.083 for
BaZr0.9Y0.1O2.95 which is close to the assumed limit of 0.1[17]. It has been proposed that
the remaining reduction of the saturation limit from the dopant concentration may be the
result from a slight incorporation of the dopant to the A-site where it acts as a donor.
9
1.5 Mobility of Protonic Defects
High diffusivities of protonic defects in the perovskite oxides are a general
phenomenon at high temperature. The principle of the transport mechanism
involves rotational diffusion of protonic defects and proton transfer between adjacent
oxygen ions rather than hydroxyl ion migration. This means that only protons show
long-range diffusion while oxygen ions reside on the crystallographic positions for
the proton transport mechanism. Kreuer et al. proved this mechanism experimentally
using BaCeO3-based oxides [18].
It is not clear what the rate-limiting step of the proton transfer mechanism is. Quantum
molecular dynamics (QM) simulations and experimental study indicate that the proton
transfer reaction in the perovskites is the rate-limiting step because the rotational
diffusion has been shown to be fast with low activation barriers [19], [20]. However, the
infrared spectra indicate that the strong hydrogen bond interactions favour faster proton
transfer reactions than rotation processes because the rotation process requires the
breaking of H-bonds [21].
The distance between oxygen ions in the perovskite oxides is usually larger than 290 pm
and strong hydrogen bonds in H-bonded crystals have 250 to 280 pm. Therefore,
the gain in free energy of the system due to hydrogen bond formation is
competing with the free energy required for the lattice distortion that is required
for hydrogen bonding. The QM simulations of cubic perovskites find protons can locally
“soften” the lattice.[20]
10
Thus, short oxygen ion separations, which favour proton transfer, and long oxygen ion
separations, which allow hydrogen bond breaking, can be achieved. Figure 2 shows a
diagram of these protons hopping between neighbouring oxygen ions. For the time
average, the presence of the protonic defect only leads to a slight reduction of the O-O
separation. However, for an instant one of the eight O-O separations is reduced to
about 280 pm to permit the hydrogen bonding and makes the proton defect behave
almost like a free hydroxyl.
Figure 2: Instantaneous proton transfer configurations obtained from quantum molecular dynamics (QM)
simulations.[20]
QM simulations also reveal the proton is not found between the two oxygen ions on the
edge of the octahedron but outside the BO6 octahedron as part of a strongly bent
hydrogen bond. The reason for this is probably the repulsive interaction between
the proton and the B-site cation. The repulsion between proton and B-site cation
is also proved by the finding that the activation enthalpies of proton mobility in
cubic perovskites with pentavalent B-site cations are significantly higher than for
tetravalent B-site cation perovskites.
11
1.6 High Temperature Proton Conducting Oxides
Perovskite type oxides including rare-earth doped BaCeO3, BaZrO3, SrCeO3 and SrZrO3
were first systematically investigated by Takahashi and Iwahara in the 1980s and the
proton conductivity of the materials was revealed when exposed to H2 or water
containing atmospheres.[17] These compounds had many applications and could be used
for hydrogen and hydrogen containing compound sensors, solid oxide fuel cells,
hydrogen separation membranes and electrochemical reactors. On exposing the material
to hydrogen gas at high temperature, usually higher than 400˚C, they become almost
pure protonic conductors. The protonic conduction in these oxides was verified by
electrochemical hydrogen transport under hydrogen or water vapour containing
atmosphere at high temperatures [17].
Usually, the conductivity of the proton conductive ceramics is in the
order of 10-2 to 10-4 S cm-1 at 600 ~ 1000 ˚C. Figure 3 shows the proton conductivities of
various oxides, which are calculated from available data on proton concentrations and
motilities, for a wide range of temperatures.
1.7 Proton Conducting Electrolyte Materials
Among proton conductors, Y-doped barium cerate (BCY) shows high proton
conductivity but it tends to decompose into protonic insulating barium carbonate
(BaCO3) or barium hydroxide (Ba(OH)2) and cerium oxide (CeO2) when exposed to CO2
or high H2O containing atmosphere [22]–[26]. In contrast to BCY, yttria doped barium
12
zirconate (BZY) shows high chemical stability but has relatively low proton conductivity
due to its high proton resistance at grain boundaries [27]–[30].
Solid solutions of Zr into barium cerate (BCZY) have been proposed to compromise
between the properties of both materials. For instance BCZY shows high protonic
conductivity and suitable sintering properties for the fabrication of anode-supported
cells[17]. However, reports indicate the formation of barium carbonate and CeO2
secondary phases after exposing BCZY powders to 3% CO2 at 600˚C for 3h, indicating
chemical reactivity towards CO2 even under mild conditions [31].
A recent study has indicated that the addition of Pr as a second dopant to BZY has
beneficial effects. Fabbri et al. showed that introducing 10% Pr into BZY dramatically
improved its sinterability while retaining the excellent electrical performance [33].
BaZr0.7Pr0.1Y0.2O3-ẟ (BZPY) powders were manufactured by a combustion method. After
sintering at 1600˚C for 8 hours, the pellets achieved high density (98.6% of theoretical
density relative to BaZrO3) and showed large grain sizes of 1.7µm. BZPY displayed only
proton conductivity under humid conditions and at had a conductivity of about 10-2 S cm1
at 600˚C. Furthermore, the pellets showed excellent chemical stability under both CO2
and H2O environments. As this new material displayed both excellent sinterability
(leading to good protonic conductivity) and chemical stability, it warrants further
consideration.
13
Figure 3: Calculated proton conductivities based on proton concentrations and mobilities for various
oxides. Conductivities of oxides with perovskite-type structure are shown by bold lines. [32]
1.8 Improved Electrolyte Performance
The operating temperature of the SOFC is usually determined by the nature of the
electrolyte. Since proton conduction has a lower activation, it makes the operation of
SOFCs at intermediate temperatures far more feasible. Operation at intermediate
temperatures of 500 - 700˚C, while maintaining high proton conductivity is desirable
since it brings several benefits such as[11]:
1. Faster start-up times.
14
2. Lower heating costs.
3. Feasibility of using stainless steel metallic interconnects and components instead
of expensive metallic alloys.
4. Enhanced lifespan of the fuel cell module.
The overall benefit is expected to result in reduced costs for SOFCs, thereby making
them a more financially attractive option.
With this goal in mind, it is important to improve the electrolyte performance. One of the
greatest obstacles to the performance of doped barium zirconate is the large grain
boundary density and the resulting grain boundary resistance. Two notable techniques
for improving the sinterability of the ceramic powders are to deliberately use partially
calcined powders and to prepare powders with barium non-stoichiometry.
Haile et al have shown increased grain size and proton conductivity of BaZr0.8Y0.2O3-ẟ by
sintering partially calcined powders.[34] The BZY powders were synthesized by a
combustion method and then calcined at various temperatures. The partially calcined
samples displayed varying amounts of barium carbonate which had an influence on the
sintering process. When compared to the fully calcined samples, these displayed greater
relative density (pellets calcined at 950˚C had higher relative densities compared to
pellets made of powders calcined at 1150˚C), grain size and electrical conductivity. It
must be noted that in previous publications Sin et al reported that the presence of barium
carbonate in partially calcined powders actually inhibits grain growth in undoped barium
15
zirconate and recommended full calcination to ensure the formation of dense pellets [35].
The difference in effects can be attributed to the presence of a dopant. Since BZPY is
more similar to BZY, it was expected that the partially calcined samples may have a
positive impact on its sinterability.
A-site cation non stoichiometry has varying effects in the reported literature. Most
publications insist that barium deficiency in perovskite structures degrades the
performance of the electrolyte [27], [36]–[39] but Zhang et al. reported that barium
deficiency helped increase the sinterability of BCZY[39]. It has been proposed that this
occurs due to A-site vacancies increasing the rate of diffusion of Barium [40].
Barium excess leads to a more encouraging response. Barison et al reported an increase
in conductivity for 5% and 10% molar barium excess, stating it aided densification and
even improved chemical stability for BaCe0.65Z0.2Y0.15O3- ẟ. [41] Shima et al. concluded
that a 4% molar excess of barium resulted in increased conductivity for undoped barium
cerate[37]. Zhang et al. shows increased conductivity at 5% molar excess barium in
BaCe0.5Zr0.4Y0.1O3- ẟ [39].
This work varies both the calcination temperature and the barium stoichiometry and
investigates the effect on the properties of BZPY.
16
CHAPTER 2: EXPERIMENTAL METHODS
2.1 Synthesis
Powders of BaxZr0.7Pr0.1Y0.2O3-d (BZPY; with x = 1.1, 1.0, 0.9) were prepared using a
combustion synthesis process. The appropriate quantities of Ba(NO3)2, (Aldrich 99.9%),
Zr(NO3)2·2H2O (Wako 97%), Pr (NO3)3·6H2O (Aldrich 99.8%) and Y(NO3)3·6H2O
(Aldrich 99.8%) were dissolved in deionised water. Citric acid was used as a chelating
agent with the ratio of citric acid to total metal cations being 2:1. After obtaining a
transparent solution, NH4OH was added to adjust the pH to between 6 and 8. The
solution was heated at 80 ºC until a transparent gel was obtained, which was pre-heated
at 350 ºC for 6 h until it formed a solid pre-cursor. Portions of this solid were further
calcined at 900 ºC, 1000 ºC, 1100 ºC, 1200 ºC and 1300 ºC for 5 hours.
The calcined samples were made into pellets of about 13 mm in diameter and about 1
mm in thickness by uniaxially pressing 0.40g of powder at 220 MPa for 1 min. All of the
green pellets were sintered at 1600 ºC for 10 hours. To limit barium evaporation during
sintering, the pellets were immersed in a powder bath consisting of calcined powders of
the same composition.
2.2 Simultaneous Thermal Analysis (STA)
The precursor was also characterised by simultaneous Thermogravimetric Analysis
(TGA) and Differential Scanning Calorimetry (DSC) using a Netzsch STA 449 F1
17
Jupiter. The samples were heated up to 1300 ºC at a heating rate of 3 ºC min-1 and were
kept there until no further mass change was seen.
2.3 Scanning Electron Microscopy (SEM)
The as fractured cross sections of the sintered pellets were studied using an FEI Quanta
200 field emission scanning electron microscope (5.00kV excitation voltage) after being
sputter coated with a 2nm layer of iridium. The grain size of the sample was measured
using a mean linear intercept method[42]. An image processing software, ImageJ, was
used to measure the pore size distribution of the sample using the SEM images. ImageJ
was not used to measure the grain size, since the grain boundaries did not show a
sufficiently uniform contrast, whereas the pores did.
2.4 Archimedes Method
The relative density of the sintered pellets was measured using the Archimedes principle
and was compared to the theoretical density of BaZrO3 (Dtheory = 6.21 g cm-3). The
diameter (d, mm) and thickness (t, mm) of the sintered pellets were measured using a
vernier caliper. An analytical balance was used to measure the dry mass (m1, g) of each
pellet before immersing the pellet in deionised water heated at 100 ºC for 2h to ensure
that all the open pores in the pellet were filled with water. The pellet was then quickly
wiped to remove surface absorbed water and the wet mass (m2, g) measured.
The open porosity, bulk density (Dbulk) and relative density (Drel) were calculated using
equations (4), (5) and (6) respectively.
18
(4)
(5)
(6)
2.5 X-Ray Diffraction
Phase identification of the calcined powders and sintered pellets was performed by X-ray
powder diffraction using a Bruker D8 Advanced diffractometer (Cu Kα λ=1.5406Å,
40kV, 40mA) with a Bragg-Brentano configuration in the range of 20˚≤ 2θ ≤ 80˚ with a
step width of 0.02˚ and a scan speed of 1˚ min-1. The software Topaz was used to
perform indexing and extract the lattice parameters and the Scherrer equation was used
to calculate the crystallite size.
2.6 Electrochemical Impedance Spectroscopy
The sintered pellets were prepared for electrochemical impedance spectroscopy (EIS) by
coating both sides with silver paste. The pellets were then dried at 700 ºC for 2 h. The
conductivity was measured under dry and humidified Argon atmospheres (3 vol.% H2O)
using a multichannel potentiostat (Bio Logic SAS) in the 0.1Hz to 1MHz frequency
range with an applied ac voltage amplitude of 500mV. The temperature was varied from
300 ºC to 700 ºC with 50 ºC steps. At each temperature, the conductivity was tested
every 30 mins until the equivalent resistance no longer changed.
The resulting impedance plots were analysed in terms of an equivalent circuit model in
which (depending on the temperature regime) distinct RQ subcircuits were used to
represent the grain interior and grain boundary regions. Here, R is an ideal resistor with
impedance ZR = R and Q is a constant phase element with impedance ZQ = (Y(j ω)n)-1
19
where j = √-1, ω is frequency, and Y and n are constants with 0 < n < 1. Data analysis
was performed using the software package EC-Lab (Bio Logic, Claix, France).
2.7 Induction Coupled Plasma - Optical Emission Spectroscopy (ICP-OES)
2.7.1 Instrumentation
The multi-element determination of Ba, Zr, Pr and Y was performed using the following
equipment: an inductively coupled plasma optical emission spectrometer (ICP-OES)
model Varian 720 E-S with axial viewing and charge-coupled detector. The instrumental
parameters used for the multi-element determination were as follows: RF generator of 40
MHz, power of 1.20 kW, plasma gas flow rate of 16.5 L min-1. The elements and the
analytical spectral lines (nm) used were Ba II 230.424, Zr II 267.869, Pr II 390.843 and
Y II 324.228 where 'II' is the ionic emission line.
A digester Milestone Ethos 1 was used for the acid digestion of the samples.
2.7.2 Reagents
Ultrapure water produced from a Milli-Q purification system from Millipore (Bedford,
MA, USA) with resistivity of 18 MΩ cm-1 was used throughout the experiment. The
nitric acid was obtained from BDH and was ultrapure. A working standard solution was
prepared freshly by serial dilution from stock solutions containing 1000 mg L-1
(Inorganic Ventures) of the elements Ba, Zr, Pr and Y.
20
2.7.3 Digestion of the Samples and Determination of the Elements
Triplicates of each sample were run for the determination of the total contents of the
elements. Approximately 5 mg of each sample was placed in a digester tube and 4 mL of
concentrated (70%) nitric acid was added. The mixture was heated on a digester block
with the settings as follows: heating to 200 ºC over 20 mins followed by a holding time
of 20 mins and then cooldown. The digested solution was then quantitatively transferred
to centrifuge tubes and diluted with ultrapure water up to a final volume of 50 ml. A
blank digest was carried out in the same way as the sample of interest.
21
CHAPTER 4: RESULTS AND DISCUSSION
4.1 Chemical Composition
The results of the ICP-OES in Table 2 indicate that the desired stoichiometries of BZPY
powders have been achieved. Of particular note is the B1.1ZPY sample which, even after
calcination at 1300 ºC has retained almost all of its Barium. The stoichiometries are
normalised relative to Pr and discrepancies between the values can readily be accounted
for by uncertainties in the raw ICP results [43].
Table 2: ICP results showing the composition of the synthesised (fresh) powders. Ash refers to the powder
produced after combustion and before calcination.
Sample
B0.9ZPY Ash
B0.9ZPY 900 ºC
B0.9ZPY 1000 ºC
B0.9ZPY 1100 ºC
B0.9ZPY 1200 ºC
B0.9ZPY 1300 ºC
Molecular Ratio (Relative to Pr = 0.10)
Ba
Zr
Pr
Y
0.916
0.757
0.100
0.214
0.916
0.710
0.100
0.200
0.914
0.706
0.100
0.211
0.916
0.704
0.100
0.208
0.913
0.703
0.100
0.214
0.914
0.712
0.100
0.203
B1.0ZPY Ash
B1.0ZPY 900 ºC
B1.0ZPY 1000 ºC
B1.0ZPY 1100 ºC
B1.0ZPY 1200 ºC
B1.0ZPY 1300 ºC
1.006
1.021
1.036
1.009
1.017
1.031
0.696
0.715
0.719
0.723
0.718
0.717
0.100
0.100
0.100
0.100
0.100
0.100
0.204
0.203
0.203
0.203
0.202
0.203
B1.1ZPY Ash
B1.1ZPY 900 ºC
B1.1ZPY 1000 ºC
B1.1ZPY 1100 ºC
B1.1ZPY 1200 ºC
B1.1ZPY 1300 ºC
1.123
1.093
1.133
1.137
1.166
1.096
0.698
0.730
0.713
0.705
0.686
0.704
0.100
0.100
0.100
0.100
0.100
0.100
0.208
0.205
0.198
0.202
0.204
0.203
22
4.2 TGA and DSC
Figure 4: a)The Thermogravimetric results of the different pre-calcined powders. b) The DSC results
The TGA results in Figure 4a show what would generally be expected. The mass% of the
pre-calcined powders falls with increasing temperature, reaching a plateau at about
23
1100˚C. The transformation begins at about 400˚C, is almost completely finished by
900˚C and displays very little mass change afterwards.
All of the DSC curves show a pattern of an exothermic peak, followed almost
immediately by an endothermic peak. The smallest of such fluctuations is seen in the
B0.9ZPY powders. The first exothermic peak at about 400˚C corresponds to the start of
the mass loss in the powders. According to Sin et al[35], the perovskite structure for
undoped BaZrO3 begins to form at about 800˚C, which corresponds with the
endothermic peaks seen in Figure 4b. The shift in temperature at which the peak appears
can be attributed to the presence of Y and Pr dopants and differences in barium
stoichiometry.
4.3 XRD
The XRD results of the calcined powders are shown in Figure 5. All of the calcined
powders show some amount of BaCO3 at 900˚C and 1000˚C. The calcined powders of all
three compositions reach phase purity at different calcination temperatures. B1.0ZPY
can obtain a pure phase at 1100 ºC, which is in agreement with previous data reported in
literature [33]. B0.9ZPY and B1.1ZPY form pure perovskite phase at 1200˚C.
For the barium deficient powder, the higher temperature may be necessary, as it requires
more energy to form the perovskite phase. Figure 4b shows that the perovskite phase for
the B0.9ZPY powder does not begin forming until 900˚C, whereas for the stoichiometric
compound, it began at 750˚C. That a pure perovskite phase can be obtained even with
24
10% cation deficiency implies that some of the dopant ions may reside on the Ba, rather
than Zr sites, as shown in[18]:
(7)
(8)
The barium deficiency drives the incorporation of the dopants into the A-site, thereby
reducing the number of oxygen vacancies.
The impact of barium deficiency on the properties of doped perovskites has not been
examined as extensively as the impact of barium excess, however, several important
observations have been made. Ma et al. found that the conductivity of Y-doped barium
cerate peaked at a composition of 5 mole% Ba deficiency[22]. Kreuer et al., in contrast,
observed a dramatic decrease in the conductivity of nominally stoichiometric La-doped
barium cerate after exposure to elevated temperatures for a prolonged period (25 h at
1700 ◦C), which presumably led to BaO loss [18]. The conductivity of 10% Ba deficient
barium cerate, doped with lanthanum, was comparable.
B1.1ZPY shows pure perovskite phase by 1200˚C. This is to be expected, given the high
barium content as it is possible that not all of the barium carbonate was incorporated into
the perovskite phase.
As suggested by Shima et. al., there are three distinct defect reactions one can consider
for the incorporation of excess BaO into the structure of doped barium zirconate [37].
These (written in Kroger-Vink notation) are:
25
(9)
(10)
(11)
The first method is highly unlikely since the large radius of the Ba3+ ion makes it
difficult to accommodate it as an interstitial within the perovskite structure. The second
is highly unlikely as it would result in the creation of a highly charged zirconium
vacancy. The third is far more likely, with the Ba dopant excess being incorporated into
the B site. It would also explain the increase in lattice parameter with increasing Ba
concentration.
Alternatively, one can consider the possibility of at least some of the excess barium
forming amorphous barium carbonate and occupying grain boundaries. Its amorphous
nature would make it difficult to detect via XRD.
26
a)
b)
c)
Figure 5: The XRD results for powders calcined at different temperatures for 5h for a)B0.9ZPY,
b)B1.0ZPY and c)B1.1ZPY
27
Upon examining the crystallite sizes of the calcined powders in Figure 6, it can be seen
that there is a general increase in size with increasing the temperature. This is to be
expected as higher temperature enhances diffusion and crystallite growth. However,
there is a noticeable drop in the crystallite size of all compositions at around 1100 ºC.
For B1.0ZPY and B1.1ZPY, this is accompanied by a drop in lattice parameter for the
calcined powders as shown in Figure 7. Moreover, the lattice parameter does not increase
linearly but instead decreases to a minima before rising again. This is similar to the
pattern witnessed by Yamazaki et. al. [34] who saw that the minima for calcined powder
was obtained at 950 ºC for BaZr0.8Y0.2O3- ẟ.
Figure 6: Crystallite size of fresh powders.
The lattice parameter of the calcined powders first decreases with increasing
temperature, before reaching a minima and increasing again. For B1.0ZPY and B1.1ZPY
this occurs at the 1100 ºC mark. For B0.9ZPY, the 1200 ºC mark. After sintering, both
B1.0ZPY and B1.1ZPY have the same lattice parameter (about 4.228 Å ). This indicates
28
that that the Barium has either been incorporated into the bulk or has been removed by
evaporation. It is also different from the results of Shima et. al. [37] with Gadolinium
doped BaZrO3 where increasing Ba stoichiometry increased the unit cell volume. Of note
is B0.9ZPY which consistently has a lower lattice parameter, reaching a maximum of
about 4.0128 Å by 1100 ºC and staying there.
Figure 7 Lattice Parameter vs Calcination Temp. for: left: Fresh Powder and Right: Sintered Pellets
XRD results show that regardless of the state of the calcined powders, the sintered pellets
all have achieved pure phase (Figure 8) . It must be noted that the B1.1ZPY pellets were
exceptionally fragile after sintering. This can be attributed to the incorporation of excess
barium as BaO at the grain boundaries. Upon exposure to air, the BaO forms BaCO3,
leading to mechanical weakness of the pellet.
Lattice Parameters obtained from the XRD results are shown in Figure 7b, showing that
once sintered, the lattice parameters of B1.0ZPY and B1.1PZPY become almost the
same (~4.228 Å) regardless of initial calcination temperature.
29
a)
b)
c)
Figure 8: The XRD of the sintered pellets, using different calcination temperatures. All pellets were
sintered at 1600 ºC for 10h.
30
4.4 SEM
Figure 9 shows the results of the SEM images from the as fractured cross sections of the
sintered pellets. While it is difficult to discern any pattern by eye, it can be seen that the
stoichiometric composition shows increased densification with calcination temperature.
Figure 11a) shows that the largest average grain sizes are obtained when powders
calcined at 900 ºC were sintered. This seems to lend credence to Haile et al's proposal
that BaCO3 acts as a sintering aid and that the reactive sintering can promote grain
growth. [34] Unfortunately, the open porosity for powders calcined at 900 ºC is also high
(~15% on average, among all the compositions for that temperature).
Figure 10b shows the open porosity measurements obtained via the Archimedes method.
It can be seen that for a fixed composition, the lowest porosity and highest relative
density occur at the same temperature. However, it must be noted that the relative
density is low compared to that obtained in literature (76% for B1.0PZY-1200 ºC
compared to about 95% in other fabrications reported in literature).
Figure 10b and Figure 11a show that it is difficult to find a correlation between the
calcination temperature and the porosity of the pellets. Nor does the open porosity
correlate well with the pore size obtained from calculating from the pictures taken.
31
Composition
B1.0ZPY
B1.1ZPY
900
1000
Calcination
Temperature
/ ˚C
1100
1200
1300
Figure 9: SEM pictures of the as fractures cross-sections of sintered pellets.
32
B0.9ZPY
a)
b)
Figure 10: a)Relative density of pellets relative to sintered BaCO3. b) Open porosity % as determined by
the Archimedes Method.
33
a)
b)
Figure 11: a) Average Grain Size and b)Average Pore Size as determined from SEM images.
34
4.5 Proton Conductivity
a)
b)
c)
Figure 12: a)All stoichiometries in wet argon, b) B1.0ZPY in wet argon and c)B1.1ZPY in wet argon.
35
EIS measurements were performed on pellets made of powders calcined at 1300 ºC.
Figure 12a shows that the stoichiometric composition possessed the highest conductivity
in wet conditions at all temperatures between 300 to 700 ºC. Interestingly, while
B1.1ZPY had higher conductivity at higher temperatures, B0.9ZPY had higher
conductivity at lower temperatures. This could be explained by the fact that excess
barium may form barium oxide at the grain boundaries, increasing the overall grain
boundary resistance. Hence, the excess barium pellets had a lower conductivity at lower
temperatures where grain boundary effects dominate. B0.9ZPY on the other hand,
probably has a lower oxygen vacancy concentration since some of the dopants such as Y
or Pr may occupy A-sites in order to maintain the perovskite structure. This would
explain its inferior conductivity when compared to B1.0ZPY pellets.
36
CHAPTER 5: CONCLUSION
An attempt was made to improve the sintering of BaZr0.7Pr0.1Y0.2O3- ẟ to enhance grain
growth and hence, performance. Two strategies were employed in attempt to achieve
this. Firstly, a reactive sintering method involving partially calcined powders containing
some amount of BaZrO3, as proposed by Yamazaki et al. [34]. The second method was
to prepare 10% molar barium excess and deficient powders as some in the literature had
reported superior conductivity with barium excess[41] and slightly enhanced sintering
with a small amount of barium deficiency.[39]
Contrary to expectations, the presence of BaCO3 was detrimental to the sintering process
of BZPY, with fully calcined samples displaying higher densities to their partially
calcined counterparts. In this respect, these findings are more in line with Sin et al. who
also reported detrimental effects on sintering if BaZrO3 powders were not fully calcined
beforehand[35].
Barium deficiency was detrimental to the sintering process and is a result in line with the
majority of publications. Unfortunately, barium excess also did not improve the
sinterability of BZPY and in terms of electrical performance, was inferior to the barium
deficient samples at lower temperatures. Both barium deficient and barium excess
samples showed poorer conductivity to the pellets made out of the fully calcined
stoichiometric BZPY.
37
In light of these findings, other methods to enhance the grain growth of BZPY need to be
sought.
38
BIBLIOGRAPHY
[1]
U.S. Energy Information Administration, “World Energy Consumption Report
January 2015.” U.S. Energy Information Administration, 2015.
[2]
W. Zittel, J. Zerhusen, and M. Zerta, “Fossil and Nuclear Fuels – the Supply Outlook,” no. March, pp. 1–41, 2013.
[3]
R. . Khurmi, Materials Science. 2008.
[4]
J. Larminie and A. Dicks, Fuel Cell Systems Explained, 2nd ed. Wiley, 2003.
[5]
E. Subhash, C. Singhal, and K. Kendall, High Temperature Solid Oxide Fuel
Cells: Fundamentals, Design and Applications. Elsevier, 2003.
[6]
H.-H. Mobius, “On the history of solid electrolyte fuel cells,” J. Solid State
Electrochem., vol. 1, pp. 2–16, 1997.
[7]
B. Viswanathan and M. A. Scibioh, Fuel Cells: Principles and Applications.
Universities Press, 2007.
[8]
S. P. S. Badwal and K. Foger, “Materials for Solid Oxide Fuel Cells,” Mater.
Forum, vol. 21, p. 183, 1997.
[9]
S. P. S. Badwal, F. T. Ciacchi, and V. Zelizko, “The effect of alumina addition on
the conductivity, microstructure and mechanical strength of zirconia — yttria
electrolytes,” Ionics (Kiel)., vol. 4, no. 1–2, pp. 25–32, 1998.
[10] S. P. S. Badwal and K. Foger, “Solid oxide electrolyte fuel cell review,” Ceram.
Int., vol. 22, no. 3, pp. 257–265, 1996.
[11] D. J. L. Brett, A. Atkinson, N. P. Brandon, and S. J. Skinner, “Intermediate
temperature solid oxide fuel cells.,” Chem. Soc. Rev., vol. 37, no. 8, pp. 1568–
1578, 2008.
[12] F. a. Coutelieris, S. Douvartzides, and P. Tsiakaras, “The importance of the fuel
choice on the efficiency of a solid oxide fuel cell system,” J. Power Sources, vol.
123, no. 2, pp. 200–205, 2003.
[13] S. Assabumrungrat, W. Sangtongkitcharoen, N. Laosiripojana, a.
Arpornwichanop, S. Charojrochkul, and P. Praserthdam, “Effects of electrolyte
type and flow pattern on performance of methanol-fuelled solid oxide fuel cells,”
J. Power Sources, vol. 148, no. 1–2, pp. 18–23, 2005.
39
[14] W. Jamsak, S. Assabumrungrat, P. L. Douglas, N. Laosiripojana, and S.
Charojrochkul, “Theoretical performance analysis of ethanol-fuelled solid oxide
fuel cells with different electrolytes,” Chem. Eng. J., vol. 119, no. 1, pp. 11–18,
2006.
[15] N. Bonanos, “Perovskite solid electrolytes: Structure, transport properties and fuel
cell applications,” Solid State Ionics, vol. 79, pp. 161–170, 1995.
[16] T. Yajima, H. Suzuki, and S. Crystal, “SOLID STATE Protonic conduction in
SrZrO3-based oxides,” vol. 51, pp. 101–107, 1992.
[17] H. Iwahara, T. Yajima, T. Hibino, K. Ozaki, and H. Suzuki, “Protonic conduction
in calcium, strontium and barium zirconates,” Solid State Ionics, vol. 61, no. 1–3,
pp. 65–69, 1993.
[18] K. Kreuer, E. Schonherr, and J. Maier, “Proton and oxygen diffusion in BaCeO3
based compounds: A combined thermal gravimetric analysis and conductivity
study,” Solid State Ionics, vol. 70–71, pp. 278–284, 1994.
[19] M. S. Islam, P. R. Slater, J. R. Tolchard, and T. Dinges, “Doping and defect
association in AZrO(3) (A = Ca, Ba) and LaMO(3) (M = Sc, Ga) perovskite-type
ionic conductors.,” Dalton Trans., vol. 3, no. 19, pp. 3061–3066, 2004.
[20] R. a. Davies, M. S. Islam, a. V. Chadwick, and G. E. Rush, “Cation dopant sites in
the CaZrO3 proton conductor: A combined EXAFS and computer simulation
study,” Solid State Ionics, vol. 130, no. 1, pp. 115–122, 2000.
[21] J. Müller, K. . Kreuer, J. Maier, S. Matsuo, and M. Ishigame, “A conductivity and
thermal gravimetric analysis of a Y-doped SrZrO3 single crystal,” Solid State
Ionics, vol. 97, no. 1–4, pp. 421–427, 1997.
[22] G. Ma, “Ionic conduction and nonstoichiometry in BaxCe0.90Y0.10O3−α,” Solid
State Ionics, vol. 110, no. 1–2, pp. 103–110, 1998.
[23] W. Suksamai and I. S. Metcalfe, “Measurement of proton and oxide ion fluxes in
a working Y-doped BaCeO3 SOFC,” Solid State Ionics, vol. 178, no. 7–10, pp.
627–634, 2007.
[24] B. R. Sneha and V. Thangadurai, “Synthesis of nano-sized crystalline oxide ion
conducting fluorite-type Y2O3-doped CeO2 using perovskite-like BaCe0. 9Y0.
1O2. 95 (BCY) and study of CO2 capture properties of BCY,” J. Solid State
Chem., vol. 180, no. 10, pp. 2661–2669, 2007.
40
[25] S. V Bhide and A. V Virkar, “Stability of AB 1 / 2 ′ B ″ 1 / 2 O 3 ‐ Type Mixed
Perovskite Proton Conductors,” J. Electrochem. Soc. , vol. 146 , no. 12 , pp.
4386–4392, Dec. 1999.
[26] F. Chen, O. T. Sørensen, G. Menga, and D. Penga, “Chemical stability study of
BaCe Nd O ceramic,” J. Mater. Chem., vol. 7, no. 3, pp. 481–485, 1997.
[27] S. M. Haile, G. Staneff, and K. H. Ryu, “Non-stoichiometry, grain boundary
transport and chemical stability of proton conducting perovskites,” J. Mater. Sci.,
vol. 36, no. 5, pp. 1149–1160, 2001.
[28] J. Tong, D. Clark, M. Hoban, and R. O’Hayre, “Cost-effective solid-state reactive
sintering method for high conductivity proton conducting yttrium-doped barium
zirconium ceramics,” Solid State Ionics, vol. 181, no. 11–12, pp. 496–503, 2010.
[29] S. Ricote, N. Bonanos, H. J. Wang, and B. a. Boukamp, “Conductivity study of
dense BaZr0.9Y0.1O(3−δ) obtained by spark plasma sintering,” Solid State Ionics,
vol. 213, pp. 36–41, 2012.
[30] W. Sun, L. Yan, Z. Shi, Z. Zhu, and W. Liu, “Fabrication and performance of a
proton-conducting solid oxide fuel cell based on a thin BaZr0.8Y0.2O3-??
electrolyte membrane,” J. Power Sources, vol. 195, no. 15, pp. 4727–4730, 2010.
[31] M. Scholten, J. Schoonman, J. Vanmiltenburg, and H. Oonk, “Synthesis of
strontium and barium cerate and their reaction with carbon dioxide,” Solid State
Ionics, vol. 61, no. 1–3, pp. 83–91, 1993.
[32] K. D. Kreuer, “Proton Conducting Oxides,” Annu. Rev. Mater. Res., vol. 33, no. 1,
pp. 333–359, 2003.
[33] E. Fabbri, L. Bi, H. Tanaka, D. Pergolesi, and E. Traversa, “Chemically stable Pr
and y Co-doped barium zirconate electrolytes with high proton conductivity for
intermediate-temperature solid oxide fuel cells,” Adv. Funct. Mater., vol. 21, no.
1, pp. 158–166, 2011.
[34] Y. Yamazaki, R. Hernandez-Sanchez, and S. M. Haile, “High total proton
conductivity in large-grained yttrium-doped barium zirconate,” Chem. Mater., vol.
21, no. 13, pp. 2755–2762, 2009.
[35] A. Sin, B. El Montaser, P. Odier, and F. Weiss, “Synthesis and Sintering of Large
Batches of Barium Zirconate Nanopowders,” J. Am. Ceram. Soc., vol. 85, no. 8,
pp. 1928–1932, 2002.
[36] Y. Yamazaki, R. Hernandez-Sanchez, and S. M. Haile, “Cation non-stoichiometry
in yttrium-doped barium zirconate: phase behavior, microstructure, and proton
conductivity,” J. Mater. Chem., vol. 20, no. 37, p. 8158, 2010.
41
[37] D. Shima and S. M. Haile, “The influence of cation non-stoichiometry on the
properties of undoped and gadolinia-doped barium cerate,” Solid State Ionics, vol.
97, no. 1–4, pp. 443–455, 1997.
[38] J. S. Park, Y. B. Kim, J. An, J. H. Shim, T. M. Gür, and F. B. Prinz, “Effect of
cation non-stoichiometry and crystallinity on the ionic conductivity of atomic
layer deposited Y:BaZrO3 films,” Thin Solid Films, vol. 539, pp. 166–169, 2013.
[39] C. Zhang and H. Zhao, “Structural characteristics, sinterability and electrical
conductivity of Ba-site non-stoichiometric BaxCe0.50Zr 0.40Y0.10O3-??,”
Materials Research Bulletin, vol. 45, no. 11. pp. 1659–1663, 2010.
[40] L. Ge, W. Zhou, R. Ran, S. Liu, Z. Shao, W. Jin, and N. Xu, “Properties and
performance of A-site deficient (Ba0.5Sr0.5)1-xCo0.8Fe0.2O3-δ for oxygen
permeating membrane,” J. Memb. Sci., vol. 306, no. 1–2, pp. 318–328, 2007.
[41] S. Barison, M. Battagliarin, T. Cavallin, S. Daolio, L. Doubova, M. Fabrizio, C.
Mortalò, S. Boldrini, and R. Gerbasi, “Barium Non-Stoichiometry Role on the
Properties of Ba 1+ x Ce 0.65 Zr 0.20 Y 0.15 O 3-δ Proton Conductors for ITSOFCs,” Fuel Cells, vol. 8, no. 5. pp. 360–368, 2008.
[42] A. M. Gokhale, “ASM Handbook,” in ASM Handbook, Ohio: ASM International,
2004, pp. 428–443.
[43] P. R. (Inorganic V. Gaines, “10,000 ppm Bismuth for ICP.” p. 2012, 2012.
42