MECHANOCHEMICALLY SYNTHESIZED
NANOMATERIALS FOR INTERMEDIATE
TEMPERATURE SOLID OXIDE FUEL CELL
MEMBRANES
James Pieter Hos BE(Hons) BSc(Hons)
School of Mechanical Engineering
The University of Western Australia
Nedlands WA 6009
This thesis is presented for the degree of Doctor of Philosophy of the University of
Western Australia.
April, 2005
Abstract
In this dissertation an investigation into the utility of mechanochemically synthesized
nanopowders for intermediate temperature solid oxide fuel cell components is
reported. The results are presented in the following parts: the synthesis and
characterisation of precursors for ceramic and cermet components for the fuel cell; the
physical and electrical characterisation of the electrolyte and electrodes; and the
fabrication, operation and analysis of the resulting fuel cells.
Samarium-doped (20 mol%) ceria (SDC) nanopowder was fabricated by the solid-
state mechanochemical reaction between SmCl3 with NaOH and Ce(OH)4 in 85 vol%
dilution with NaCl. A milling time of 4 hours and heat treatment for 2 hours at 700°C
yielded a material with equivalent particle and crystallite sizes of 17 nm. The
existence of a complete solid solution was affirmed by electron energy loss
spectroscopy and x-ray diffraction analysis. Doped-ceria compacts were sintered for
4 hours at 1350°C forming ceramics of 88% theoretical density. The ionic
conductivity in flowing air was 0.009 S/cm, superior to commercially supplied
nanoscale SDC.
Anode precursor composite NiO-SDC nanopowder was synthesized by milling
Ni(OH)2 with the previously defined SDC formulation. Larger batches of materials
were obtained in an attrition mill and when reduced to a cermet at 50 vol%Ni-SDC
the electrical conductivity was 1068 S/cm at 600°C. These cermets were
impermeable to O2 and N2 at room temperature hence it was necessary to introduce
porosity by the inclusion of 17 vol% graphite into the composite. This reduced the
cermet’s density from 88% to 55% whilst maintaining a conductivity 468 S/cm and
ii
the permeability was significantly increased. Image analysis of electron micrographs
showed that the triple-phase boundary area between electrolyte, nickel catalyst and
pore space was 211 m/cm2. This was directly attributable to the ultra-fine
composition of 158 nm SDC and 196 nm Ni grains.
Anode-supported fuel cells were fabricated on a substrate of at least 500 μm
55wt%NiO-SDC with 17vol% graphite pore formers. Suspensions of SDC were
deposited by aerosol on the sintered bilayer at a thickness around 5 μm. A cathode of
10% SDC – (SmSr)0.5CoO3 was deposited onto the sintered electrolyte and after firing
had a thickness of around 25 μm.
Operation of fuel cells in single-chamber mixtures of CH4 and air diluted in argon
were successful and gave power outputs of 483 μW/cm2. Operation in undiluted 25
vol% CH4:air gave a power output of 5.5 mW/cm2. It was shown that a large
polarisation resistance of 4.1 Ω.cm2 existed and this was assigned to losses in the
anode, namely mass transport limitation associated with the catalytic combustion of
methane and insufficient porosity. The large surface area of Ni appeared to allow
more methane to combust and hence prevented its electrochemical reaction from
occurring, thus limiting the performance of the cell. The synthesis procedures,
ceramic processing and fabrication techniques and testing methods are discussed and
contribute significant understanding to the fields of ceramic science and fuel cell
technology.
iii
Acknowledgements
I would like to extend sincere appreciation to the following individuals for the
assistance they provided to me while I was completing this project.
Prof. Paul McCormick for allowing me to undertake this project, providing critical
feedback and for his sense of humour. I am also grateful to Dr Karl Foger for
allowing me to liaise with researchers at Ceramic Fuel Cells Ltd.
I am indebted to the technical and workshop staff from the Centre for Microscopy and
Microanalysis, the School of Mechanical Engineering and the School of Biomedical
and Chemical Sciences. I would like to acknowledge Mr Greg Cole for the
construction of the quartz test rig. Special thanks to Dr Martin Saunders for
assistance with high-resolution TEM and to Dr Frank Lincoln for proofreading this
thesis. Thanks to the research and development staff of APT/ANT who I worked with
over the years, in particular Geoff Trotter, Michihito Muroi, Takuya Tsuzuki, Aaron
Dodd and John Robinson.
I am grateful to have received financial support from the Minerals and Energy
Research Institute of WA.
Finally for my wonderful wife Linda, without whose support and encouragement I
would have been lost, thankyou for your patient understanding and endless love.
iv
Contents
Abstract ii
Acknowledgements iv
Contents v
Figure Captions ix
List of Tables xv
1 Introduction 1
1.1 Solid Oxide Fuel Cells 5
1.2 Synthesis of Nanoparticles and Nanostructured Powders 7
1.3 Study Objectives 9
2 Literature Survey 13
2.1 Principles of SOFC Operation 13
2.1.1 State-of-the Art SOFC Technology 17
2.1.1.1 The Influence of Anode Microstructure 19
2.2 Intermediate Temperature SOFC Technology 22
2.2.1 Ionic Conduction in Doped Ceria Phases 22
2.2.2 Intermediate Temperature SOFC based on doped ceria 24
2.2.3 Single-Chamber SOFC 28
2.3 Processing Methods for Improved SOFC Performance 31
2.3.1 Toward Nanostructured Interfaces in Solid Oxide Fuel Cells 32
2.3.1.1 Microstructural Refinement of Electrolytes 32
2.3.1.2 Microstructural Refinement of Electrodes 35
2.3.1.3 Mechanochemical Synthesis of Nanopowders 37
3 Experimental 40
v
3.1 Powder Handling 40
3.2 Powder Synthesis 40
3.2.1 Reagents 40
3.2.2 Mechanochemical Processing 41
3.3 Fabrication Techniques 42
3.4 Characterisation Techniques 43
3.4.1 X-ray Diffraction (XRD) 43
3.4.2 Specific Surface Area Measurement 44
3.4.3 Electron Microscopy 44
3.4.4 Thermal Analysis 44
3.4.5 Dilatometry 45
3.4.6 Gas Permeation Measurements 45
3.4.7 Conductivity and Fuel Cell Performance Apparatus 46
3.4.7.1 Total Conductivity Measurements 46
3.4.7.2 Design and Evaluation of Testing Rigs 47
3.4.7.3 Fuel Cell Performance Measurements 48
4 Synthesis and Characterisation of Nanopowder Precursors 50
4.1 Sm-Doped CeO2 50
4.1.1 Mechanical Alloying of Sm2O3 with CeO2 50
4.1.2 Mechanochemical Synthesis of Sm0.2Ce0.8O1.9-x 52
4.1.2.1 Mechanochemical Synthesis of SDC Nanopowders in an Attrition
Mill 57
4.1.2.1.1 Attrition Milling of Commercial SDC 57
4.1.2.1.2 Mechanochemical Synthesis of SDC in an Attrition Mill 59
4.2 Mechanochemical Synthesis of NiO-SDC Composite Powders 62
vi
4.2.1 Mechanochemical Synthesis of NiO-SDC in an Attrition Mill 65
4.3 Synthesis of (SmSr)0.5CoO3 68
4.3.1 Characterisation of Commercial 10%SDC-SCC 70
5 Processing of Mechanochemically Synthesized Nanopowders 73
5.1 Slip Characterisation and Tape Casting of CeO2 Nanopowder 73
5.1.1 Slip Formulation 73
5.1.2 Thermal Analyses and Sintering Behaviour 78
5.2 Compaction of Nanopowders 82
5.2.1 SDC Electrolyte 82
5.2.2 NiO-SDC Anode Precursor 86
5.2.3 SDC-SCC Cathode 89
6 Electrical and Microstructural Characterisation of Fuel Cell Components 91
6.1 SDC Electrolyte 91
6.2 Ni-SDC Anode 92
6.2.1 Cermet Precursor 93
6.2.2 Cermet Formation 96
6.2.2.1 Pore Forming Additives 100
7 Fabrication and Performance of a Single Chamber SOFC 111
7.1 Electrolyte Fabrication by Spin Coating 112
7.1.1 In-Situ Reduction of Cermet Precursors 115
7.1.2 Cell construction using Pt mesh current collector 116
7.1.3 Cell construction using Ag wire current collector 119
7.1.4 In-Situ Reduction of Cathode 122
7.2 Electrolyte Fabrication by Airbrush 125
7.2.1 Cell construction using Ag wire current collector 126
vii
7.2.2 Cell construction using Ag wire current collector and Ag composite
paste on Cathode 130
7.3 Single Chamber Fuel Cell Operation in Undiluted Gas Mixtures 131
8 Conclusions 139
8.1 Summary 139
8.2 Suggestions for Further Work 140
9 References 143
10 Publications and Presentations during PhD Canditature 150
viii
Figure Captions
Figure 1.1. Schematic of a single solid oxide fuel cell ..................................................3
Figure 1.2 Graphical representation of a fuel cell anode/electrolyte interface ..............6
Figure 1.3 Schematic representation of the mechanochemical process........................9
Figure 2.1 Typical voltage losses for an operational fuel cell after Larminie [16]......14
Figure 2.2 Planar and Tubular SOFC schematics ........................................................18
Figure 2.3 Theoretical percolation according to the General Effective Media theory
(left) [25] and Ni percolation in YSZ at 1000°C experimental(right) [26].........21
Figure 2.4 Single Chamber Fuel Cell schematic [49] .................................................28
Figure 3.1 Stylised SPEX mill operating ....................................................................41
Figure 3.2 Schematic of permeability test setup .........................................................45
Figure 3.3 Digital photograph of a typical sample prepared for 4-pt conductivity
measurements. Electrical contact achieved with Ag wire and Pt paste ..............46
Figure 3.4 Flow cell used for conductivity and fuel cell performance measurements.
The thermocouple is located at A, leads to specimen are crimped onto Pt wires at
B ...........................................................................................................................48
Figure 3.5 Fuel cell clamped between porous alumina with attached leads ...............48
Figure 3.6 Circuit diagram for measuring fuel cell output. Battery element is the fuel
cell test specimen .................................................................................................49
Figure 4.1. TEM Micrograph of Mechanically Alloyed (Sm2O3)0.2(CeO2)0.8 .............50
Figure 4.2 Lattice parameter vs Sm content in SDC...................................................52
Figure 4.3. Development of mechanochemically milled SDC lattice parameter and
particle/crystallite sizes with annealing temperature and time ............................54
Figure 4.4. TEM Lattice image of mechanochemically synthesized SDC ..................55
ix
Figure 4.5. EFTEM Elemental maps and reference image of SDC Nanopowder. From
bottom left: Ce-M edge, Sm-M edge, O-K edge. All micrographs have the same
field of view. ........................................................................................................56
Figure 4.6. AMR SDC as received and after (a) 4 hrs SPEX milling, (b) 80 mins
attrition milling; both received heat treatment for 30 mins at 500°C ..................58
Figure 4.7 SPEX Milled AMR SDC heat treated 2 hrs at 700°C ................................59
Figure 4.8 X-ray diffractogram of 60 min attrition milling of SDC precursors,
anomalous peaks suggesting unreacted SmCl3 indicated with arrows.................60
Figure 4.9. TEM Micrograph of mechanochemically synthesized SDC, milled twice
(80mins + 80mins) in a horizontal attrition mill, heat treated at 700°C for 2hrs.61
Figure 4.10 XRD Pattern of 50 wt%NiO-SDC nanocomposite, NiO peaks indicated
with arrows...........................................................................................................63
Figure 4.11 Energy filtered transmission electron micrographs of 50%NiO-SDC
nanopowder: (Top-left) Unfilitered, (Top-right) Oxygen map, (Bottom-left)
Cerium map, (Bottom-right) Nickel map.............................................................65
Figure 4.12 35 wt%NiO-SDC milled for 60 mins in SPEX showing wide particle size
distribution and hard agglomerates ......................................................................67
Figure 4.13 Plot of lattice parameter and crystallite size against attrition milling time
for the 55wt%NiO-SDC reaction mixture ...........................................................68
Figure 4.14 X-ray diffractogram of cubic (SmSr)0.5CoO3 ..........................................69
Figure 4.15 Transmission electron micrographs of SCC and 10%SDC-SCC supplied
by MicroCoating Ltd............................................................................................71
Figure 4.16 MCT 10%SDC-SCC as received (bottom) and after calcination at 350°C
for 2hrs (top). Squares denote SCC reflections, circles SDC, question marks
denote un-indexable reflections ...........................................................................72
x
Figure 5.1 Rheometry titration of 40wt% 25nm CeO2/ethanol with DisperByk®-111
..............................................................................................................................74
Figure 5.2 Shear stress and viscosity vs shear rate for 58.5wt% 25nm CeO2/ethanol
suspension dispersed with 3.5% DisperByk®-111, .............................................75
Figure 5.3 Electron micrograph of green 25nm CeO2 tape..........................................77
Figure 5.4 Transmission electron micrograph of a diluted slip sample .......................78
Figure 5.5 TGA of green 25 nm CeO2 tape. Squares = weight loss, Circles =
derivative weight loss with respect to temperature ..............................................79
Figure 5.6 Dilatometry curves for green 25 nm CeO2 pellets. Squares = temperature
program, diamonds = shrinkage vs. time, circles = shrinkage vs. temperature ...80
Figure 5.7 XRD Crystallite size vs. sintering time at temperature for tape cast CeO2
..............................................................................................................................81
Figure 5.8 Electron micrograph of the sintered grain structure on the material’s
surface (left); sintered tape cross-section fracture surface(right).........................81
Figure 5.9 Dilatometry curve for MCP-SDC...............................................................83
Figure 5.10 SEM fracture surface of sintered MCP-SDC, 91% dense ........................84
Figure 5.11 Dilatometry curve for AMR-SDC, heat-treated for 2 hrs at 700°C..........85
Figure 5.12 Dilatometry curves for compacted NiO/SDC composites.......................87
Figure 5.13 Sintered 35%NiO composite pellet micrographs: (Left) Polished surface,
(Right) Fracture surface .......................................................................................88
Figure 5.14 TEM reference image and EFTEM Elemental Map 75% dense
35wt%NiO-SDC: NiO = Blue, SDC = Red .........................................................89
Figure 5.15 Dilatometry curves for 10%SDC-SCC....................................................90
Figure 6.1 Temperature dependence of the conductivity of doped ceria specimens ...91
xi
Figure 6.2 Cermet precursor conductivity as a function of NiO-SDC composition
showing percolation of acceptor-doped NiO. ......................................................93
Figure 6.3 Conductivity vs reciprocal temperature for NiO-SDC cermet precursors..95
Figure 6.4 Cermet formation via reduction of NiO-SDC composites in 10%H2/Ar at
600°C ...................................................................................................................96
Figure 6.5 Electron micrograph of 65wt%NiO-SDC ceramic showing the interface
between Ni and NiO(left), backscattered electron image of the same region(right)
..............................................................................................................................97
Figure 6.6 Conductivity of Ni-SDC cermets at 600°C as a function of Ni content.....98
Figure 6.7 Temperature dependence of conductivity for 55 wt%NiO-SDC and 50
vol%Ni-SDC ........................................................................................................99
Figure 6.8 Scanning electron micrograph of 50vol%Ni-SDC cermets.....................100
Figure 6.9 Electron micrograph from a fractured pellet of 55 wt%NiO-SDC with 20
vol% graphite .....................................................................................................101
Figure 6.10 Dilatometry curve for 55 wt%NiO-SDC with 10 vol% graphite included.
............................................................................................................................102
Figure 6.11 Electron micrographs of cermet formed from 55 wt%NiO-SDC with 15
vol% graphite; Secondary image(right), backscattered image(left)...................103
Figure 6.12 Scanning electron micrographs of reduced and polished 50vol%Ni-SDC
from 9.1vol%(left) and 16.8vol%(right) included graphite ...............................103
Figure 6.13 Permeability vs pressure difference for reduced cermets with added
graphite...............................................................................................................104
Figure 6.14 Backscattered electron micrograph of 50vol%Ni-SDC with 13.2vol%
graphite added. White regions=SDC, Grey=Ni, black=pores............................106
xii
Figure 6.15 Backscattered electron micrograph of 50vol%Ni-SDC showing the pore
under consideration ............................................................................................108
Figure 6.16 Microstructure model used to determine TPB length............................109
Figure 7.1 55wt%NiO-SDC//SDC anode substrate spin-coated, sintered at 1300°C
(left) vs 1400°C (right). Arrows denote the anode//electrolyte interface .........113
Figure 7.2 Micrograph of an oven-dried spin coated layer of AMR-SDC showing
agglomerates and surface cracking ....................................................................114
Figure 7.3 Fuel cell microstructure showing spin-coated SDC electrolyte and SDC-
SCC cathode.......................................................................................................115
Figure 7.4 Conductivity and temperature vs annealing time for the reduction of
55wt%NiO-SDC ceramic under 200mL/min FCG............................................116
Figure 7.5 OCV as a function of furnace temperature for the fabricated fuel cell ...117
Figure 7.6 Power vs current drawn for fuel cell 7.1.2 at selected temperatures between
550 and 750°C in 260 mL/min FCG..................................................................118
Figure 7.7 Power vs current drawn for fuel cell 7.1.3 over selected temperatures
between 550 and 750°C in 260 mL/min FCG....................................................119
Figure 7.8 OCV and Maximum power output for fuel cell 7.1.3 as a function of
temperature.........................................................................................................120
Figure 7.9 Polarisation Loss as a function of current drawn from the fuel cell ........121
Figure 7.10 Conductivity vs temperature for sintered 10 wt%SDC-SCC in flowing air
............................................................................................................................123
Figure 7.11 Temperature and conductivity vs. annealing time for 10%SDC-SCC
ceramic. Gas composition is FCG(6%CH4, O2) and is changed as marked. ....124
xiii
Figure 7.12 Electron micrographs of airbrushed SDC on NiO-SDC substrates, sintered
at 1350°C for 4hrs. The cross section shows the slightly porous SDC layer on
the base(left); a typical SDC surface region(right) ............................................126
Figure 7.13 Power vs current drawn for the fuel cell at selected temperatures in 30
mL/min FCG ......................................................................................................127
Figure 7.14 OCV and Maximum power output for fuel cell 7.2.2 at a function of
temperature.........................................................................................................128
Figure 7.15 Temperature dependence of ohmic (Rint) and interfacial (Rpol)
resistances for fuel cell from Section 7.2.2 ........................................................129
Figure 7.16 Cross section of fuel cell after testing with the anode support on the left
side .....................................................................................................................130
Figure 7.17 Dried 160°C (left) and fired 800°C (right) Ag-based cathode current
collector paste.....................................................................................................131
Figure 7.18 Cross section of cell failed during exposure to CH4:Air atmosphere.
Extensive crack propogation in the lower anode region has caused the electrolyte
and cathode (upper region) to also fracture........................................................133
Figure 7.19 Cell power generation from 50mL/min CH4, 250mL/min air ...............135
Figure 7.20 Variation of open-cell voltage and power output for SOFC in 50 mL/min
CH4, 250 mL/min air..........................................................................................136
Figure 7.21 Cell losses as a function of temperature as determined by GCI .............137
xiv
List of Tables
Table 2.1 Comparison of Commercial SOFC Designs ................................................18
Table 2.2 Summary of recent intermediate temperature SOFC developments............25
Table 2.3 Comparison of nanopowder and particle characteristics according to
synthesis method ..................................................................................................39
Table 4.1. Comparison of Literature and Industry Values for the Lattice Parameters of
the Sm-CeO2 system ............................................................................................51
Table 4.2. Summary of BET and XRD characterisation data for NiO/SDC
nanopowders ........................................................................................................63
Table 5.1 CeO2 nanopowder slip properties compared to literature source [105] ......76
Table 5.2 Typical slip formulation...............................................................................76
Table 5.3 Densification data for AMR-SDC and MCP-SDC .....................................86
Table 5.4 Sintering data for NiO/SDC compacts........................................................87
Table 6.1. Comparison of Conductivity Parameters for the Sm-CeO2 System ...........92
Table 6.2 Contiguity data for Ni-SDC(E)-Pore cermet ............................................107
Table 7.1 Summary of cell component resistances based on GCI measurements ....122
Table 7.2 ICP Elemental assay for 10%Sm0.2Ce0.8O1.9-x-(SmSr)0.5CoO3 .................125
Table 7.3 Anode supported NiO-SDC//SDC//SCC-SDC solid oxide fuel cell
fabrication regime ..............................................................................................132
xv
1 Introduction
Since the industrial revolution the quality of life has been largely dictated by access to
energy for heating, power and transportation. The combustion of fossil fuels is the
conventional means to this end. Centrally located coal fired power plants distribute
electricity generated at 40% efficiency over a grid to end users. Automobiles are
powered by liquid fuels of high specific energy density: diesel and petrol distilled
from crude oil and combusted internally at around 30% efficiency. As the standard of
living in the developing world increases, the challenges to supply the automotive and
domestic energy markets will be significant. To supply this energy in a sustainable
manner to maintain air quality is a challenge for which conventional power generation
technologies are not ideally suited.
The availability of coal and the magnitude of power that can be generated result in
coal power supplying almost 50% of total OECD demand [1]. However the
environmental impact of coal combustion products, both gaseous chemicals and
particulate soots, are more readily visualised than they are quantified. It is debateable
whether current reserves of oil are sufficiently large for foreseeable demand, yet there
are unavoidable issues associated with the resource’s geographical location and its
subsequent extraction and distribution. It is also recognised that the reserves of coal
and natural gas far outstrip those of oil, but both resources suffer in terms of a poor
energy density for gas and undesirable emissions from coal. Natural gas is the
cleanest of the fossil fuels so in order to sustainably increase power generation
capacity, it would appear sensible to increase the utilisation of gas in conventional or
alternative power plants.
1
Natural gas is commonly encountered as a heating fuel, yet centrally located gas-fired
power plants are often employed to supplant base-load generation during periods of
excessive demand. In transmitting this energy to the customers over a grid the losses
are estimated at 6.8% of total OECD demand [2], which is equivalent to the annual
energy demands of the global transport sector. Since distributed gas networks exist
within the infrastructure of most developed centres, the ability to generate power at
smaller stations closer to customers could eliminate these losses. Nations such as The
Netherlands, Finland and Denmark all have at least 40% of their electricity demand
provided by distributed sources [3] which yield a 20-30% reduction in greenhouse gas
emissions. The power generation devices are predominantly diesel or gas
reciprocating engines but the drive for higher efficiency and lower emissions means
that micro-turbines are gaining acceptance and alternative technologies are becoming
more viable.
The losses associated with turbine-electric generators are fundamentally bound by the
Carnot cycle, before the effects of friction reduce the efficiency to 40%, at best. By
reacting the fuel electrochemically all of the free energy can theoretically be converted
into electricity. Furthermore, because there is no combustion central to the operation
of a fuel cell there are no emissions such as partial combustion products, nitric or
sulphur oxides. This is the concept of a fuel cell, first demonstrated by William Grove
in 1839, electrochemically reacting hydrogen and oxygen gases over platinum
electrodes in a sulphuric acid electrolyte to produce a potential difference and flowing
current. A fuel cell is analogous to a battery except that charge is continually
generated rather than stored, provided there is a continual supply of fuel. At that time
2
there were no means to convert electricity into mechanical energy, or a practical
means to mass produce hydrogen, so for some time fuel cells remained an academic
pursuit. This research lead to the emergence of distinct types of fuel cells that can be
differentiated by the material employed as an electrolyte. The “solid oxide” fuel cell,
shown in Figure 1.1 for example, has an ionically conductive ceramic as an electrolyte.
Grove’s original cell, the “Gas Voltaic Battery” relied on hydrogen ion conduction
through sulphuric acid, but the fundamentals components remain an anode, electrolyte,
and cathode. The technical challenges arise when connecting the cells together in
series with seals and manifolds for gas supply.
Figure 1.1. Schematic of a single solid oxide fuel cell
The nature of the electrolyte and its corresponding operating temperature and suitable
fuel has lead to the development of the following types of fuel cells:
• Polymer Electrolyte Membrane
• Phosphoric Acid
3
• Molten Carbonate
• Solid Oxide
Polymer Electrolyte Membrane Fuel Cells (PEMFC) use platinum catalysts finely
dispersed in an organic matrix and are ideal for applications around 1kW. Whilst
there are variants that can operate using methanol as a fuel, PEMFC require high
purity hydrogen to prevent catalyst poisoning. Because hydrogen is considered a
chemical, not a commodity or fuel, the commercialisation of hydrogen powered fuel
cells is linked to the economic production of hydrogen.
It is therefore necessary for fuel cells to operate from fuels other than hydrogen.
Methane and methanol can be reformed or partially oxidised into mixtures of carbon
monoxide and hydrogen, or the cells can be designed to operate from these fuels
directly. Molten Carbonate Fuel Cells (MCFC) have a readily scaleable molten bath
for an electrolyte and can produce hundreds to thousands of kilowatts. Solid Oxide
Fuel Cells (SOFC) employ an oxide-ion conducting ceramics as an electrolyte. SOFC
technologies have been shown in various implementations to operate from 750°C to
1000°C from hydrogen, or reformed methane. Solid oxide fuel cells are tolerant to
carbon monoxide and can utilise it as a fuel, making them ideal when coupled to a
methane steam reformer for operation from natural gas. The absence of corrosive
liquids, high pressures or excessive temperatures means they are ideal for local
combined heat and power generation. Demonstration units up to 2 MW have been
exhibited in recent years by companies including Siemens-Westinghouse(US-Europe),
Sulzer(Europe), Rolls Royce(UK), Ceramic Fuel Cells Ltd (Australia) plus a host of
contributions from research institutes, laboratories and universities. Despite this
4
considerable effort there remain significant engineering challenges to bring SOFC
technology to the marketplace, that has been valued at US$123 million and forecast to
reach $335 million by 2008 [4]. In order to access this market, the solid oxide fuel
cells must be ready for mass production at a cost of US$400/kW.
1.1 Solid Oxide Fuel Cells
The SOFC requires elevated temperatures to achieve sufficiently high conductivity in
its electrolyte, which in state-of-the-art examples is a yttria-stabilized zirconia (YSZ).
On the fuel side an anodic material is required, typically a composite between the
electrolyte and a catalytically active metal with appreciable electronic conductivity,
typically nickel. Hydrogen gas reacts with the oxygen ions at the triple-phase
boundary inside a pore where fuel gas, nickel and electrolyte interconnect. The
oxygen ions have diffused through the electrolyte from oxygen reduced at the cathode.
The nickel phase must allow the electrons released by the oxidation of hydrogen gas
to percolate out of the anode, so that their energy can be collected upon returning to
the cathode. Furthermore the anode must have open porosity to allow diffusion of
products away from the reaction zone. A simple schematic of the triple phase
boundary is shown in Figure 1.2 for an SDC (samarium doped ceria) electrolyte.
5
Figure 1.2 Graphical representation of a fuel cell anode/electrolyte interface
The cathode layer exists on the air side and the material must allow the facile
dissociation of oxygen and diffusion of the atoms. Doped perovskites of the type
(A,B)XO3 such as strontium-doped lanthanum manganese oxide, La0.8Sr0.2MnO3,are
selected to have chemical and mechanical compatibility with the electrolyte. As for
the anode, the triple phase boundary between electrolyte and cathode is the location
where incoming electrons reduce oxygen atoms dissociated on the perovskite surface,
so the cathode must also have a degree of porosity.
The power generating characteristic of a fuel cell is distinct from that of a battery, in
that the voltage extracted decreases non-linearly with the current drawn over the load.
The two factors that contribute to these internal losses are the ohmic resistances of the
anode, cathode and electrolyte and the non-ohmic anodic and cathodic polarisation
resistances, associated with the electrochemical processes occurring at each interface.
To achieve the desired benchmarks for SOFC commercialisation, the performance
6
must be increased as economically as possible. The selection of more exotic, highly
conductive materials will improve the performance; however the costs make this
approach prohibitive. More significant gains could be achieved by engineering ever
finer, more open and function microstructures. This implies that the microstructure of
the ceramic will be sub-micron and entering the nano-scale regime, which is only
feasible if the starting materials are of similarly small dimensions.
As the particle size decreases from micron, to sub-micron, ultra-fine and finally
nanosized, a number of phenomena can be observed in the ceramics formed after
sintering. The sintering temperature and duration required to achieve full density are
lowered and while this is normally accompanied by rapid grain growth, in some cases
it is possible to preserve nanosized grains [5]. Grains of such small dimensions can
theoretically increase the ionic conductivity by virtue of an increased grain boundary
volume and its associated grain boundary diffusion coefficient, which is normally four
times larger than the bulk conductivity [6]. Furthermore, on the anode or cathode side,
nanoscale microstructures have a higher internal surface area and hence a longer
triple-phase boundary length. When using nanoparticles there are considerable
challenges to obtain and maintain stable suspensions and nanoscale features, but the
novelty and opportunity for improved properties are so great that it was a strategy that
was actively pursued in this thesis.
1.2 Synthesis of Nanoparticles and Nanostructured Powders
Nanopowder synthesis methods can be classified as top-down or bottom-up. Most fall
into the bottom-up method, where particles are precipitated from a gas or liquid and
their growth is restrained by various chemical (surfactant), or physical (vapour
7
combustion) means. The measure employed to minimise agglomeration of the
product phase will be discussed in more detail in the literature review. Top-down
methods are far less common and generally less effective, involving the mechanical
attrition of large grains into smaller ones. The energy efficiency of such a process is
low and the probability of contamination from mill lining and media is high.
Mechanochemical synthesis is the solid-state mechanical activation of precursor
chemicals in a ball mill. Nanopowders are obtained when this process occurs in an
excess volume of an unreactive and water soluble salt phase [7, 8], shown
schematically in Figure 1.3A for the reaction of ZrCl4 with CaO in excess CaCl2. The
typically oxide nanopowder product phase reaches a steady-state as a nanocomposite
within the salt matrix as a result of repeated welding and fracturing from the milling
action shown in Figure 1.3A to B. The reaction may occur during milling, but final
dehydration and calcination is effected by a mild heat treatment, below the melting
temperature of the salt phase in Figure 1.3B to C. Because the majority salt phase
maintains the separation of nanoparticles no agglomeration occurs during this heat
treatment. Washing the soluble phase away yields nanoparticles that may be as small
as 4 nm, but typically range from 10-70 nm in Figure 1.3C to D.
8
Figure 1.3 Schematic representation of the mechanochemical process
Agglomeration during washing is not significant and provided the slurries are not
completely dried, the particles adopt a re-dispersible flocculated state. Bottom-up
methods generally suffer from excessive agglomeration as batch sizes in either
solution or vapour phases are increased to an economic volume. The
mechanochemical method suffers no such drawbacks and the equipment is more
readily scaled for mass production.
1.3 Study Objectives
There exists a considerable research and development effort largely focussed on the
commercialisation of SOFC’s based on YSZ electrolyte technology, which could be
considered mature. Research into the use of alternative electrolyte materials to enable
9
intermediate temperature of solid oxide fuel cells will allow the development of the
second generation of SOFC’s. Electrolytes such as doped cerium oxides and doped
lanthanum gallium oxides receive the most attention for their high ionic conductivities
from 400-700°C and favourable stability and compatibility properties. Mixed-
conducting materials such as ceria, in combination with more robust metallic catalysts
can give enhanced anode performance and there are a wide range of doped perovskite
phases suitable for cathodes. As in all ceramic processing science, the physical and
electrical properties are dictated by the microstructure which is a product of several
processing steps, starting with powder synthesis and ceramic preparation.
Mechanochemical processing techniques, to be discussed more thoroughly in section
2.3.1.3, provide nanoparticles with unique characteristics but have not been applied
directly to the fabrication of solid oxide fuel cell components.
Samarium-doped cerium oxide (SDC) has been chosen for this study because of its
high oxide-ion conductivity, which at 500°C is fractionally higher than lanthanum-
doped gallium-oxide and equivalent to YSZ’s conductivity at 700°C [9]. The
mechanochemical synthesis of SDC nanopowder electrolyte powders would also
allow the fabrication of a Ni-SDC cermet from NiO-SDC precursors. The optimum
cermet composition will be determined by measuring the conductivity as a function of
Ni content; the transition from ionic to electronic conduction corresponding to the
formation of a percolated Ni network. Nanoparticulate SDC can then be deposited as
a thin film supported on the cermet precursor and the microstructure investigated.
It has been shown that the cathode composition depends on the environment and
temperature and, recently, a single-chamber fuel cell, where the cell is exposed to a
10
uniform mixture of fuel and oxygen, has received attention [10, 11]. In these single-
chamber solid-oxide fuel cells a strontium-doped samarium cobalt oxide
((SmSr)0.5CoO3, SSC) perovskite was employed as a cathode [12]. The cobaltite
cathode materials such as SCC, or more generically lanthanum-strontium cobaltite
(LSC) exhibit higher performance than the manganese oxide based perovskite
cathodes (LSM) that are commonly employed with YSZ electrolytes [13]. This is
because LSM has far less propensity to react with YSZ to form resistive zirconates
phases at high temperatures and has prompted some researchers to employ an SDC
interlayer between cathode and electrolyte [14], but it would appear to be more
sensible to use the high conductivity electrolyte exclusively with high performance
cathode material.
The single chamber fuel cell sacrifices the oxygen gradient achieved when the fuel
and air are separate, but its strength is believed to lie in the resulting simplifications to
gas delivery. The electrochemical driving force is generated by differing rates of
catalytic activity: the anode catalyses the oxidation of fuel faster than the cathode, and
oxygen is reduced faster at the cathode. By choosing the single-chamber fuel cell
design for testing, considerable savings can be made which are significant for an
investigation at this level. These experimental details will be explained in Chapter 3,
following a more thorough survey of the literature of ceramic processing as it relates
to solid oxide fuel cells and mechanochemical synthesis of nanopowders, in Chapter 2.
The results and discussion of the three phases of the project are dealt with in Chapters
4-7, namely the Synthesis and Characterisation of Nanopowder Precursors; a study of
the Processing of Mechanochemically Synthesized Nanopowders; and finally the
Electrical and Microstructural Characterisation of and the Fabrication and
11
Performance of a Single Chamber SOFC. The mechanochemical process was
successfully applied to the synthesis of a nanostructured cermet anode supported thin
electrolyte, this process and the performance of the fuel cell thus fabricated will be
presented.
12
2 Literature Survey
Many comprehensive reviews of SOFC technology exist [15]. The purpose of this
chapter is to introduce the relevant concepts, issues and existing solutions. Examples
where nanostructured interfaces have been engineered successfully into SOFC test
subjects will be discussed and the utility of mechanochemically synthesized
nanomaterials for this task will be revealed.
2.1 Principles of SOFC Operation
For the electrochemical reaction between hydrogen and oxygen in Equation 1 the
electromotive force is calculated using the Gibbs free energy change at 700oC via the
Nernst Equation in Equation 2, where z = 4 (number of electrons transferred) and F is
Faraday’s constant.
2H2(g) + O2(g) 2H2O(g) , ΔG°700 = -388 kJ/mol Equation 1
E°700 = -ΔG°700 / zF = 1.005 V Equation 2
This potential is the open cell voltage (OCV) and when a current, I, is drawn the
terminal voltage, V, drops due to the following irreversibilities:
• Activation losses, caused by reaction kinetics at the surface of either electrode
• Fuel crossover or internal currents, caused by electron conduction through
electrolyte or its permeability to fuel
• Ohmic losses associated with the electrical resistance of anode, cathode and
electrolyte
13
• Mass transport/concentration losses, caused by restrictions to the
supply/removal of reactants/products by diffusion as the reaction rate increases
These factors are related in Equation 3 [16] where RA, RE, RC are the ohmic
resistances of Anode, Electrolyte and Cathode; ηC and ηA are the cathodic and anodic
activation overpotentials and ηMT is the mass transport overpotential. Note that the
three overpotentials are themselves a function of the current density.
V(i) = E° – (RA + RE + RC)i – [ηC(i) + ηA(i)] - ηMT(i) Equation 3
More rigorously, the electrode activation polarisations are proportional to ln(i) and the
mass transport overpotential can be approximated according an exponential relation.
The resultant I-V characteristic is shown in Figure 2.1 and it applies in general to fuel
cells of any type.
(RA + RE + RC)I
Circuit current I [mA/cm2]
Cell voltage [V]
OCV
ηC
ηA
Figure 2.1 Typical voltage losses for an operational fuel cell after Larminie [16]
14
The performance of a SOFC is bounded by the open-cell voltage and the maximum
current extractable as determined by the activation and concentration polarisations.
Since the OCV is determined by the fuel and operating temperature, performance
improvements can be achieved by reductions in the activation, mass transport and
ohmic losses. Activation polarisation is strongly affected by the electrode’s
preparation and resultant morphology and is equally important to both electrodes.
Ohmic losses are decreased by having as thin and as highly conductive cell
components as possible, whilst mass-transport polarisation is minimised by having as
porous and open electrodes as possible. Obviously these demands cannot be met
simultaneously, especially when the structural integrity of a cell must be ensured.
The operation of a SOFC from pure hydrogen fuel are reasonably well understood and
the reactions are believed to take place in the so-called “electrochemical reaction
zone”[17] that is postulated to exist in the first 10-20mm of triple phase boundary
from the electrolyte/anode interface. Commercial operation from pure hydrogen fuel
is impractical, and given the abundance and relative ease of processing of methane, it
would appear to be the logical fuel cell fuel of choice and it is from methane that
hydrogen is produced. Despite its flammability, the chemical reactivity of methane is
comparatively low due to the strength and symmetry of its C-H bonds. It has the
thermodynamic properties approaching that of an ideal gas, evidenced most obviously
by a boiling temperature of -161°C. Hydrogen can be obtained from methane by two
rather extreme measures. Steam reformation, as shown in Equation 4, occurs in a
steam rich environment over alumina-supported nickel catalysts at pressures of 3 MPa
and temperatures around 900°C.
15
CH4 + H2O CO + 3H2 Equation 4
Alternatively there is partial oxidation which can be simply expressed in Equation 5.
The actual process involves a number of sub-reactions and requires temperatures in
excess of 1200°C and the exclusion of air to prevent ammonia formation.
CH4 + 0.5 O2 CO + 2H2 Equation 5
Partial oxidation is hampered by the propensity for methane decomposition to carbon,
in the form of carbon black or soot. Steam reformation is the favoured method to
generate fuel usable for SOFC’s, but partial oxidation is useful as there is no steam
management necessary. Domestic natural gas typically contains 70-90% CH4 with
smaller amounts of higher-hydrocarbons and trace amounts of sulphur containing
gases. The sulphur compounds will poison the anode catalysts at high temperatures
and therefore the gas must be treated in a desulphuriser before use. The fuel
processing sections of an SOFC system can represent a considerable fraction of the
total cost so technologies that are more robust will have an advantage.
It is feasible for methane to be reformed internally on the anode. Steam reformation
will occur if the temperature is above 700°C and there is sufficient steam to inhibit the
decomposition of methane via cracking or CO disproportionation. The carbon thus
deposited blocks the pores and channels in the anode microstructure and rapidly
degrades the fuel cell performance. Effective operation from methane is essential for
the progress of SOFC technology. The majority of commercially developed SOFCs
16
employ an external reforming and as such operate from hydrogen, with some notable
exceptions. Regardless of the fuel, the overall cell efficiency, η, is strongly dependent
on the fraction of fuel utilisation, α, shown in Equation 6 [9].
η = ηgηvα Equation 6
Where ηg = nFE°/ΔH is the Gibbs efficiency and ηv = (E° - IR)/E° is the voltaic
efficiency (R is the total area specific cell resistance, RA + RE + RC). Numerical
modelling of an internally-reformed SOFC stack showed an electrical efficiency of
62.2%, with heat recovery this increases to 83.8% [18].
2.1.1 State-of-the Art SOFC Technology
The cell shown previously in Figure 1.1 is the conceptual core of a planar SOFC. To
complete the system there must be a means to supply fuel and air to the anode and
cathode, remove exhaust gases and collect the current. The drive for more power
entails increasing both current drawn and voltage extracted: to increase the voltage
cells are combined in series and increasing the surface area of each cell generates
larger currents. The geometry of the cell, planar or tubular, as shown in Figure 2.2,
will determine the method by which cells are connected together.
17
Figure 2.2 Planar and Tubular SOFC schematics
Solid oxide fuel cell stacks are manufactured by interconnecting individual cells with
bipolar current-collecting plates, which must also serve to feed fuel and air to anode
and cathode respectively forming the stack. Alternatively, tubular fuel cells increase
their power by increasing the number of tubes, current is collected at the end of each
bundle where gas is fed in. These two designs illustrated schematically in Figure 2.2
have sufficiently distinct issues and solutions and the planar type is adopted by the
majority of researchers and developers. For this discussion only the planar system
will be considered, anode-support is compared to electrolyte support in Table 2.1.
Table 2.1 Comparison of Commercial SOFC Designs
Category Electrolyte Supported [19, 20]
Anode Supported [21]
Methane reforming method
Internal Externally
Electrolyte 10YSZ(+ 15wt% Al2O3) 150 μm
8YSZ 5-10 μm
Cathode LSM 50 μm LSM 50 μm Anode Ni/YSZ 50 μm Ni/YSZ 1 μm Interconnect Ni/Ag Ferritic stainless steel Stack Size 150W 800W Temperature 875°C 750°C Specific Power Output(Fuel)
0.1 W/cm2 (CH4/H2O) 1.4 W/cm2, (H2/H2O)
18
The anode supported fuel cell logically has a thinner electrolyte than the electrolyte-
supported design which can account for some of its enhanced power output.
Operation with methane via internal reforming, for the electrolyte supported case,
gives a less complex system but produces significantly a less power. To achieve
acceptable performance at temperatures below 750°C in a fuel cell that operates
directly from methane alternative electrolyte materials are necessary. Ceria has a
variable oxidation state and an oxygen storage capacity, desirable properties for an
anode to operate from methane. These concepts will be discussed in 2.2, taking
examples from the academic and patent literature. Before this is done, however, it is
worthwhile exploring further the importance of the anode microstructure, as it is
central to SOFC operation regardless of the temperature or fuel.
2.1.1.1 The Influence of Anode Microstructure
The anode of an SOFC must exhibit the following properties to achieve minimal
polarisation losses and robust function:
• Catalytic activity towards the fuel
• Open porous structure for diffusion and flow of gases
• Thermal expansion compatibility with the electrolyte
• Percolation of electronically conducting phase
• Contiguity with electrolyte to form a large triple phase boundary
In the case of an anode supported fuel cell, the anode must also be mechanically self-
supporting, so it is difficult for any design to satisfy all demands completely. The
19
high temperatures that may be encountered during processing and operation will
fundamentally determine the anode’s initial structure and its evolution.
Conventionally this is achieved by sintering a NiO-YSZ composite to give a good
bond to the electrolyte. Upon exposure to fuel cell conditions, the volumetric
shrinkage occurring as NiO is reduced to Ni introduces a 34% increase in pore/void
volume which is generally sufficient to form open porosity, but pore-forming
additives are often added. The most popular of these are carbon blacks or graphitic
phases, which tend to induce less shrinkage during firing than organics such as
cellulose or starch. Corbin and Apte [22] report that 150 μm graphite introduces
porosity without increasing the shrinkage of the specimen.
Alternatively, Kim and co-workers contend that when compacting an anode, its
microstructure can be better engineered by the addition of a deformable fugitive phase
such as thermoset polymers [23]. The nature of the starting materials and the
processing ultimately determine the anode’s performance, and there are a number
reviews that discuss the NiO-YSZ anode system [17, 24]. The relative particle sizes
of the NiO and YSZ starting powders are of particular importance. For operation at
1000°C the coarsening of Ni is rapid, so the use of sub-micron or finer NiO is
wasteful. To achieve electrical percolation, large spherical metallic particles will
yield larger thresholds so large NiO powders are not a disadvantage. The percolation
of biphasic, randomly packed spheres of identical diameters [25] has a characteristic
as shown in Figure 2.3(left), and a percolating network is formed at a volume fraction
of 0.162 in the conductive phase. This is compared to the percolation of Ni in YSZ,
but in a triphasic system pore percolation occurs at 30vol%Ni of total solids [26],
shown in Figure 2.3(right).
20
Figure 2.3 Theoretical percolation according to the General Effective Media theory (left) [25] and Ni
percolation in YSZ at 1000°C experimental(right) [26]
When the conductive particles are anisotropic then percolation can be achieved at
lower volume fractions, but the cubic symmetry of both Ni and NiO makes the
synthesis of such particles difficult. Fine YSZ particles in the anode are beneficial,
giving a larger triple phase boundary length and short diffusion distances. As the
microstructure becomes finer there is the risk of excessive tortuosity that may hinder
gas permeation.
The fabrication of a SOFC anode is significantly more complicated than other state-
of-the art technologies such as the anodes of photovoltaic solar cells and tantalum
capacitors. The two-dimensional electrode in a solar cell should have as large an area
as possible to collect the current with minimal losses, yet because it can never be
optically transparent, it can only exist at the expense of the optical flux. This is
comparable to the need for a SOFC anode to have a high electrical conductivity, but if
this were so continuous as to be two dimensional then gas flow and oxygen ion
21
percolation would be compromised. High C-V tantalum capacitors have anodes
fabricated from high surface area Ta powders [27]. When compacted and sintered,
Ta2O5 dielectric is electrolytically grown on the free surfaces, however the
densification must be controlled to give a good interparticle contact without excessive
coarsening. The SOFC anode must possess a contiguity of three continuous phases –
metal, ceramic and air so the sintering process must be even more carefully controlled
to retain gas permeability.
2.2 Intermediate Temperature SOFC Technology
Ceria based solid electrolytes are a popular choice for intermediate temperature SOFC
research. Doped ceria is more structurally and crystallographically stable, unlike
doped bismuth phases, and are far cheaper than gallate based phases.
2.2.1 Ionic Conduction in Doped Ceria Phases
Pure, dense CeO2 is a weak n-type (electronic) conductor and at high temperatures
and/or reducing atmospheres it loses oxygen and becomes non-stoichiometric.
Oxygen vacancies, VO••, and Ce3+ defects are thus formed via Equation 7 and the ceria
is often denoted as CeO2-x (Kröger-Vink notation, Ce`Ce is a Ce3+ defect).
OO + 2CeCe VO•• + 2Ce`Ce + 0.5 O2(g) Equation 7
Doping the fluorite structure of CeO2 with lower valency cations, M = Ca, Gd, Sm,
has a similar effect. The degree of doping is associated with changes in the cubic
lattice parameter, a=5.411Å for pure ceria which can form extensive solid solutions
22
with many elements. The substitution for Ce4+ and concomitant oxygen vacancy
generation are shown in Equation 8. It should be noted that this will affect the
equilibrium of Equation 7 by altering the concentration of vacancies and CeCe [28].
2MO1.5 2M``Ce + OO + VO•• Equation 8
When M=Sm, the ionic conductivity of 20 mol% samarium-doped ceria,
Sm0.2Ce0.8O1.9 (SDC) is an order of magnitude greater than pure CeO2 and increases
from 0.001 S/cm at 400°C to 0.1 S/cm at 800°C. Ionic conduction occurs via the
diffusion of oxygen ions through these vacancies according to σi=Cqμ, where C is the
number of anion vacancies per unit volume, q is the charge per carrier and μ the
carrier mobility. The Nernst-Einstein relation for mobility and diffusivity gives the
temperature dependence of conductivity according to the Arrhenius-type Equation 9
[29], where [VO••] is the vacancy concentration and pre-exponential A are constant for
a given composition.
σiT = A[VO••]exp(-Ea/kT) Equation 9
Samarium-doped ceria is characterised by an activation energy Ea of 0.78 eV [30].
The dependence of ionic conductivity on dopant radii and concentration implied
defect ordering via complexation so Ea is more properly expressed as the sum of the
association enthalpy, ΔHA and the enthalpy for migration, ΔHm for which values of
0.5-0.61 eV are generally accepted [29]. Most instances in the literature index their
conductivity measurements to the combined activation energy, Ea.
23
The maximum conductivity in the ceria system occurs when the host and dopant ionic
radii are similar. This is when the association enthalpy between dopant and vacancy
is minimised. Since the ionic radii of Ce4+ and Ce3+ are 0.97 and 1.14 Å respectively,
there is more than enough room for the 1.1 Å Sm3+ and Ca2+ dopants that have
traditionally high conductivities. By correlating an effective crystallographic index,
Mori et al. were able to show an order of magnitude improvement by further doping
SDC with both Cs and Li [31]. The extra doping was also claimed to limit the
reduction of ceria in reducing atmospheres.
2.2.2 Intermediate Temperature SOFC based on doped ceria
Whilst 20 mol% samarium-doped ceria would appear to be the most obvious choice
for an intermediate temperature electrolyte based on its well documented conductivity,
there is some controversy in the literature regarding the superiority of gadolinium
doping. It is suggested that Gd0.1Ce0.9O1.95 (GDC) is a better electrolyte based on
predicted low vacancy-solute complex binding energies [32], but this is hardly evident
on a practical scale. Nevertheless, there is much research activity into fuel cells based
on Gd or Sm doped CeO2. Table 2.2 summarises recently published examples of such
research, the thickness of the specified component indicates that it is the support on
which the cell is fabricated.
24
Table 2.2 Summary of recent intermediate temperature SOFC developments.
LSCF = (LaxSr1-x)(CoyFe1-y)O3, LSM = (LaxSr1-x)MnO3, SCC = (SmxSr1-x)CoO3
Author Electrolyte Anode Cathode Atmos-phere
Performance
Metcalfe [33] GDC Ni-8YSZ 1mm LSCF CH3OH
reformate
60 mW/cm2
@ 600°C
Ormerod [34] GDC 280 μm Ni-GDC LSCF 10%H2-N2 90 mW/cm2
@ 600°C
Mori [31] (Cs, Li)SDC
500 μm
Ni-YSZ LSM 3%H2O-
H2
220 mW/cm2
@ 700°C
Gorte [35] YSZ Ni(Cu, CeO2)-
YSZ 230μm
LSM CH4 130 mW/cm2
@ 800°C
Liu [36] SDC Ni-SDC 300μm SCC 3%H2O-
CH4
304 mW/cm2
@ 600°C
Pham [37] SDC Ni-SDC 200μm LSCF CH4 320 mW/cm2
@ 550°C
Barnett [38] (YDC)YSZ Ni-YSZ LSM 1 mm CH4 370 mW/cm2
@ 650°C
Hibino [39] GDC Ni(Ru)-GDC 1
mm
SCC CH4 750 mW/cm2
@ 600°C
The first two examples highlight the futility of using conventional ceria based
electrolytes with hydrogen atmospheres. The poor power output of Ormerod’s cell is
directly associated with the relatively thick Gd0.1Ce0.9O1.95 electrolyte’s partial
reduction, evidenced by an OCV below 1V. Only by increasing the level of dopant in
the electrolyte, as Mori shows, can one begin to obtain power densities in excess of
200 mW/cm2 from hydrogen fuel. Before examining methane, the efforts of Metcalfe
deserve a mention for extracting 60mW/cm2 from partially reformed methanol,
despite the strange choice of YSZ in the anode supported cell when it would appear
more sensible to use GDC since it is also the electrolyte.
The five remaining cells exemplify the difference between direct methane utilisation
and internal reforming, as steam must be present with the fuel stream to allow steam
25
reformation without carbon deposition over Ni-YSZ. The YSZ based direct methane
fuel cell shown by Gorte cannot be classified as intermediate temperature, but its
performance is reported [40-42]. The precipitation of both CuO and CeO2 within the
porous anode just prior to operation facilitated the direct electrocatalytic oxidation of
methane without carbon deposition. Barnett and co-workers utilise ceria in a different
manner, the (YDC)YSZ refers to a porous interlayer of (Y2O3)0.15(CeO2)0.85 (YDC)
deposited on either side of the electrolyte by D.C. magnetron sputtering. In real terms,
the YDC interlayer reduces the interfacial resistance by a factor of 6 giving the 370
mW/cm2 of power in either moist or dry methane.
When SDC was used as the electrolyte by Liu et al., a comparable power density was
obtained at a lower operating temperature. This work introduced the notion of
engineering ever finer structures and thinner electrolytes to improve performance and
employs a glycine-nitrate technique to synthesize nanopowders. Pham and Glass [37]
divulged the process for depositing an electrolyte via ultrasonically assisted colloidal
spraying. The reported energy density from methane is impressive given the
operating temperature, but there are no microstructures presented nor have there been
any supporting publications.
Hibino et al. were able to produce more than twice the power of any cells listed in
Table 2.2 from a cell modified only by the inclusion of 3wt% Ru into the GDC based
anode. This had a profound affect and the authors claim it was primarily responsible
for catalysing the complete anodic oxidation of methane via Equation 10. This would
give a theoretical OCV of 1.04 V at 600°C, so the reported values in the 890 mV
26
range suggest that this conversion is incomplete. Furthermore, the use of costly
metals such as Ru, Pt and Pd limits and commercial application.
CH4 + 4O2- 2H2O + CO2 + 8e- Equation 10
Higher performances were reported with methane over hydrogen on account of lower
OCVs in the partially reduced GDC electrolyte. The increased OCV loss when 2SDC
was used was rationalised by the assumption that Ce4+ is more readily reduced as the
degree of doping increases. This is at odds with the findings of Mori in the previous
section.
For some years Hibino and co-workers [43] have developed the idea of operating a
fuel cell in a single chamber of gas, where the fuel and air are mixed. The concept
originated when researchers in the nuclear fission industry considered the means to
electrochemically recombine dilute hydrogen and oxygen mixtures generated by
radiolytically split water [44]. The initial fuel cells were alkaline-solution based, but
the concept is equally applicable to solid electrolytes as shown by Dyer [45]. In 1990
Dyer reported 1-5 mW/cm2 associated with a potential difference between two Pt
electrodes separated by a semi-permeable membrane, either pseudoboehmite
(AlO(OH)) or Nafion™ (sulfonated fluoroethylene polymer) in an atmosphere of H2
and O2 at room temperatures. In all cases the EMF is determined by the oxygen
concentration in the locality of either electrode, which is generated by the selective
oxidation of hydrogen on the anode and reduction of oxygen on the cathode. This
arrangement has also found utility in direct methanol fuel cells, eliminating the
problem of methanol crossover [11]. When mixed-reagents are used the electrolyte
27
need not be impervious to the fuel, oxidant or reaction products, so the weakness of
ceria-based SOFCs can be overcome.
2.2.3 Single-Chamber SOFC
Hibino et al. demonstrated single chamber SOFCs for YSZ, LSGM, GDC and SDC
electrolyte systems in the planar, tri-layer cell configuration shown in Figure 2.4 [46-
48]. Hibino’s intermediate temperature SDC single chamber cell gave far better
performance with ethane and propane, which is understandable, given the presence of
additional C-C bonds, but importantly showed a degree of efficiency when operating
directly from methane. This shows that such cells can operate directly from LPG or
butane [49] as well as natural gas and not suffer from carbon deposition that would
occur if internal reforming were attempted.
Hydrocarbon + air feed
Solid oxide electrolyte
Ni-SDC electrolyte
SCC electrolyte
Figure 2.4 Single Chamber Fuel Cell schematic, adapted from [49]
A mixture of 30% CH4, 15% O2 with balance N2 flowing over a 150 μm electrolyte
supported cell (90%NiO-SDC anode, 10%SDC-SCC cathode) gave a power output
around 50 mW/cm2 over a 350-600°C temperature range [50]. Again, a massive
performance boost was observed when precious metals were included in essentially
the same anode structure. The addition of 7wt% PdO to the 70wt%NiO-SDC cermet
28
precursor gave 644 mW/cm2 of power when fabricated on a 500 μm electrolyte pellet
in a 1:1 CH4:O2 gas mixture at 550°C [51]. Ideally noble metal catalysts should not
be central to the development of intermediate temperature SOFC technology.
However, direct methane operation under 2 mg/cm2 is comparable to the utilisation of
Pt in PEMFC’s, which also require trace amounts of Ru and/or Rh which are vastly
more expensive.
A study of anode versus electrolyte support was reported for the YSZ system by
Jasinski and coworkers, with a Ni anode and LSM cathode. They obtained 4.1
mW/cm2 at 730°C for electrolyte support and 3.4 mW/cm2 at 539°C in a gas mixture
of 10% propane, 90% air [52]. This was followed by testing in methane atmospheres
and an anode supported Ni-YSZ cell and LSCF cathode delivered 120 mW/cm2 at
750°C in 17% methane, 83% air [53]. They describe a low-temperature processing
method, with sintering temperatures below 1000°C and employed a polymeric YSZ
precursor deposition method developed by Chen [54, 55] that will be discussed in
Section 2.3.1.1.
Whilst the low performance and fuel utilisation limits have caused some to consider
single chamber fuel cells unfeasible for large scale power generation, a healthy
academic and technical interest in the field of microscale power generation exists [56].
Hibino et al. patented the fundamentals of their definitive work [50]. More recently,
researchers at HP have patented designs for SOFC single chamber reactors for what
could be considered a fuel battery [57]. This is similar in form and function to the
single chamber fuel battery patented by Shinko Electric Industries [58], who reported
an output of 3.63 mW from an unspecified cell size, employing a Li-doped NiO anode
29
which they presume does not reduce in the methane-air atmosphere and an SDC
electrolyte. Both patents focussed on maintaining dimensions less than the
extinguishing diameter, 0.1-3.0 mm for methane-air, within the cell chamber to
prevent explosive combustion.
The simplicity inherent to a single-chamber fuel cell design allows rapid screening of
cell performances as shown for both Ni-YSZ anode support[52] and YSZ electrolyte
support [53]. In both theUS [59-61], Europe [62] there is much interest in cell design
and structure-performance relations of single-chamber SOFC technology and this
thesis intends to contribute to this knowledge.
Single chamber SOFCs may also have utility in electrocatalysis, where the selectivity
of cathodic or anodic reactions is controlled by the potential in a given atmosphere.
The electrocatalytic reforming of CO2 by methane has been demonstrated in an SOFC
system with an MgO-modified NiO anode [63]. The system catalyses the reaction in
Equation 11 and extracts currents from the oxidation of CO and H2 equivalent to 16
mW/cm2 at 800°C in a YSZ electrolyte system.
CH4 + CO2 2CO + 2H2 Equation 11
Similarly, the electrocatalytic oxidative coupling of methane begins to realise the
potential for its activation by charged species such as oxygen ions [64]. A polarised
Ag-Sm2O3 anode layer inside a YSZ tube was synthesized 85% purity ethylene [65]
from CH4, which can be mechanistically visualised as the oxidative coupling of CH3
fragments and the anodic reaction of H2. The reactor design involves trapping the
30
larger molecular weight hydrocarbon products in molecular sieves to recycle the gas
for further reaction.
For either the electrochemical generation of power, or the electrocatalytic production
of hydrocarbon from methane, the single-chamber SOFC is unique as it provides
means to investigate the interaction of processing parameters on performance.
Furthermore, nascent nature of the single chamber SOFCs makes them an attractive
means to demonstrate the utility of mechanochemically synthesized nanopowders for
fabricating ultra-fine microstructures.
2.3 Processing Methods for Improved SOFC Performance
The methods employed for the fabrication of SOFCs have been covered in a number
of the references cited in previous sections [15]. In general, electrolyte powders are
commercially supplied after calcination and spray drying. Hence the SDC or YSZ
powders consist of micron sized agglomerates with primary crystallites in the sub-
micrometer range. To obtain ~10 μm electrolytes, and electrodes with high internal
surface areas, these commercial starting materials are unsuitable and to meet these
conditions they must obviously be produced in a finer state. Whilst vapour deposition
methods can deposit ultra-thin films and have produced nanoscale powders for
decades, the costs associated with these processes do not assist the commercialisation
of SOFC technology. Less costly synthetic routes are needed for the production of
nanoscale powders and the mechanochemical processing method is one such
candidate, amongst a variety of liquid phase methods.
31
2.3.1 Toward Nanostructured Interfaces in Solid Oxide Fuel Cells
Nanostructured materials are, by definition, those with characteristic dimensions of
less than 100 nm. A number of synthetic methods and processing techniques will be
discussed, and it will be evident that having nanoscale powder(s) on its own is a
necessary but not sufficient criterion for obtaining a nanostructured ceramic. If the
powder is to be compacted and densified, the large shrinkage associated with the
sintering may lead to cracking. When tape cast, there may be problems in obtaining a
workable slip due to the large amounts of organic dispersants and binders required to
keep a suspension stabilised [66, 67].
2.3.1.1 Microstructural Refinement of Electrolytes
Researchers from Nanomaterials Corporation claim that 5.8 nm nanoparticles of 9
mol% YSZ can be pressed and sintered to full density after 17 hrs at 1200°C [68].
This electrolyte had 83 nm grains (by XRD) and in a Ag-YSZ//YSZ//Ag-YSZ cell
they claimed almost an order of magnitude increase in conductivity over micron-sized
YSZ specimens. These results are possibly due to the fine microstructure of the Ag-
YSZ composite used in contact with their measurement leads, but there is an absence
of clear microstructural evidence supporting this. There is no evidence for such
improved conductivity in 1-5 mol% Co3O4-doped Gd0.2Ce0.8O2-x which had 120 nm
grains when densified. Fabricated from 20 nm GDC it could be sintered to complete
density at 900°C by a liquid phase sintering mechanism as the cobalt oxide melts [69].
Furthermore, the theoretical basis for enhanced ionic conductivity with decreasing
grain size depends crucially on having no impurities in the grain boundary phase,
which is the very place where impurities will concentrate [6].
32
The traditional method of improving sinterability is to add compounds that facilitate
liquid-phase sintering, more often than not these are SiO2 based, which is deleterious
to ionic conduction and hence unsuitable for application in fuel cells. Strict control of
silicon impurities to below 50 ppm has been reported as necessary for the fabrication
of SOFC electrolytes [32]. Rahaman obtained completely dense ceramics from
hydrothermally synthesized nanocrystalline 6 at% CaO doped CeO2 powder [70],
observing the inhibition of grain boundary diffusion due to calcium-doping. Sintering
at 1350°C gave a ceramic with sub-100 nm grains. Since Ca was far more effective at
producing small grains than Mg, Rahaman assumes that the dopants must concentrate
at the grain boundaries according to space-charge theory, and in doing so presumably
limit their mobility. As the level of Ca increased beyond 6at% the sintered grain size
increased, so the phenomena is not relevant to the more highly doped such as SDC.
The results of Rahaman are related to the findings of Chen, who states that
densification with limited grain growth relies on suppressing grain-boundary
migration whilst grain-boundary diffusion is kept active [5]. Chen achieves high
density 1mol% Mg-Y2O3 with 60 nm grains using a 2-stage sintering profile where the
temperature reaches 1310°C before holding at 1150°C until densification is complete.
Chen’s powders are prepared via the HMT (hexamethylenetetramine) method [71],
yet another homogeneous precipitation technique yielding crystallite sizes from 10-
100 nm in size. The purpose of the first sintering step is to ensure that the density is
above 70% for at this point the pores in the ceramic become sub-critical and will
shrink via capillary action without further grain growth [72]. However, Chen does not
33
report these phenomena for highly doped systems, so its applicability to SDC or YSZ
systems is questionable.
Rahaman has also produced thin, dense films upto 1.4 μm thick from well-dispersed
6at%Y-CeO2 [73]. Whilst the grain sizes are 200 nm it remains to be seen whether
the same can be achieved for SDC. Importantly, Rahaman correlates the film’s
homogeneity with the nature of the dispersion from which it was spin-coated. In
acidified water the hydrothermally precipitated oxide is dispersed with 10wt%
polyvinylpyrrolidone (PVP) and, in what is an excellent example of electro-steric
dispersion, report a well-dispersed suspension that yields a defect-free film [74].
Chen has also developed thin ceria films by the deposition of polymeric SDC
precursor solutions on support structures of YSZ or LSCF [54]. The suspension is an
in-situ, acid-catalysed polymerisation of ethylene glycol and allows crystallisation at
320°C and hence the 500 nm membranes can be densified at 800°C. In subsequent
years this process was incorporated in a patent [55], which discloses that increasing
the solids content of the polymer by suspending a colloidal powder, i.e. SDC
nanoparticles, prior to its deposition by spin coating.
In a similar vein, Visco et al. disclose the deposition of commercially available GDC
(unspecified particle size) suspended in alcohol on to a pre-fired NiO-GDC anode
precursor [75]. The substrate was held at 100-150°C during spraying to facilitate
rapid drying. Sintering at 1450°C for 4 hrs gave 5-20 μm of dense GDC, SDC or YSZ.
Visco et el. suggest that the membrane and substrate should have identical shrinkages,
failing this the shrinkage of the substrate should be greater than the membrane so that
34
any excess curvature can be alleviated without introducing cracks in the membrane.
The best way to control shrinkage rates is to control the particle size of the starting
materials and considerable flexibility is gained when the particles are in the nanoscale
range. This was not possible for Visco et al. who were restricted by the nature of the
commercially supplied oxides.
2.3.1.2 Microstructural Refinement of Electrodes
The previous section covered a number of homogeneous precipitation methods for the
production of nanopowders and in general these techniques are applicable for the
fabrication of anodes, anode precursors or cathodes. Seabaugh et al. reported the
fabrication of all popular SOFC electrolytes and electrodes from nanopowders
synthesized via their precipitation technique [76]. The process is continuous and
following precipitation the material is subject to the following operations: Filtration &
Washing; Drying & Sieving; Calcination; Milling and a final stage of Drying/Sieving
[77]. Perhaps such a large number of steps results in the prohibitive cost, for which 1
kg of 140 m2/g SDC was quoted at US$1,650. Whilst the patent claims to allow the
production of nano-composite electrodes their existence is unclear from the images
provided, and whilst the material has a high surface area suggesting sub-100 nm
dimensions the crystallites are clearly agglomerated into 500 nm particles. This
shows that there is significant potential to define more exactly how fine a structure has
been obtained at the end of a particular process.
Vapour synthesis methods have been used for many years for the production of
nanopowders, the combustion chemical vapour condensation (CCVC) method of
Maric et al. [78] shows much promise for the synthesis of electrolyte and electrode
35
materials for SOFCs. This method is simpler than precipitative techniques, the excess
oxygen present during flame-induced combustion means that the crystallites will
oxidise completely so no calcination or processing is required aside from collecting
the powder. The researchers from MicroCoating Technologies are able to synthesize
NiO-SDC, SCC-SDC electrode materials and SDC, with surface areas from 112-175
m2/g. Sintering occurs for SDC at 1150°C and complete densification is claimed and
appears evident. When an SOFC is fabricated from such materials, the ceria does not
appear to suffer from reduction as mentioned in Section 2.2.2, delivering 630 mW/cm2
at 600°C from moist hydrogen on a 450 μm thick electrolyte. Whether this
exceptional performance is a result of the ultra-fine microstructure demonstrated is
unclear.
The glycine-nitrate precipitation/combustion technique employed by Xia and Liu in an
internally reforming anode supported fuel cell (Table 2.2) includes highly porous
foam particles of GDC, as well as NiO-GDC [79] anode substrate precursor or SCC
[80] for the cathode. For the Ni-GDC a grain size of 500 nm is reported after sintering
at 1250°C and reduction. Larger grains, obtained after sintering at 1450°C, result in
an increase in anode resistivity and poor fuel cell performance, associated with loss of
open porosity. These vapour or liquid phase precipitation techniques have been
classified as bottom-up strategies, and for laboratory scale experiments they are
sufficient. For the manufacture of larger quantities, the loss of product quality in
terms of the degree of agglomeration is unavoidable and for hydrothermal or vacuum
methods the cost of reaction vessels is high.
36
Most top-down strategies for powder synthesis will produce nanostructured powders,
not necessarily nanoscale powders. The doped ceria and bismuth phases synthesized
and characterised by Vitlov-Audino and Lincoln are obtained by the mechanical
alloying of the respective oxide phases [81], the ionic conductivities of the sintered
ceramics are not reduced by impurities originating from the stainless steel milling
media or hardened steel vial, which is a common criticism of such techniques. Pure
comminution or mechanical alloying are unable to reduce the primary particle size
below ca. 100 nm, over a realistic timescale, despite producing a nanoscale grain
structure for a range of brittle and ductile materials [82]. To break these agglomerates
and maintain fully separated crystallites an alternative technique is needed and
mechanochemical milling, as shown in Figure 1.3, is able to accomplish this, and
produce nanopowders.
2.3.1.3 Mechanochemical Synthesis of Nanopowders
The mechanochemical process is the dry milling of desired reagents, for example
nickel chloride and metallic sodium, in the displacement reaction shown in Equation
12.
NiCl2 + 2Na + xNaCl Ni + (2+x)NaCl Equation 12
When no diluent is present, i.e. x = 0, the reaction occurs via combustion due to the
large enthalpy change and the Ni product phase consists of sub-micron agglomerates
[83]. When x = 1.5, then the salt is of sufficient volume fraction to separate the Ni
product phase and allows the reaction to proceed steadily. The continual grain
refinement forms a composite structure of 10-20 nm crystallites of Ni within a NaCl
37
matrix [84]. Once the diluent phase is dissolved and product recovered, the Ni exists
as discrete, nanosized crystallites, i.e. a nanopowder.
The process is similarly applicable to oxides such as ZrO2 [85], Mg-PSZ, Y-TZP [86]
and YSZ [87] though in these cases the reaction is not completed during milling, a
mild heat-treatment is necessary to form the solid solution nanoparticles. Dodd et al.
go on to show the enhanced sinterability of nanocrystalline YSZ [87], whose sintering
onset is enhanced relative to commercial YSZ by 400°C. The mechanochemically
synthesised NiO-ZrO2 composite nanopowders previously mentioned [88] have not
been investigated further.
Both Gd2O3 [89] and CeO2 [90] nanoparticles have been synthesized
mechanochemically, the general form of these reactions is shown in Equation 13,
where Ln = Gd, Ce.
LnCl3 + 3NaOH + xNaCl Ln(OH)3 + (3+x)NaCl Equation 13
Calcination of the as-milled mixture at 500°C yields ~1 μm particles of single-phase
oxide when x = 0. When x = 11-12 the process yields 10-20 nm nanoparticles, whose
size can be controlled by the intensity of the calcination. Because the reagents in the
milling step are hydroxides or chlorides the degree of iron contamination in the
product is minimal. Furthermore, the heat treatment occurs while the nanoparticles
are separated within their NaCl matrix so particle growth is minimised agglomeration
is physically impossible. Once the washing procedure has removed the salt phase the
nanoparticles tend to exist as loosely flocculated structures. In comparison with
38
vapour and liquid synthesis methods summarised in table, mechanochemical
synthesized nanoparticles have a much narrower size distribution and a much more
controllable primary particle size, as shown in Table 2.3.
Table 2.3 Comparison of nanoparticle characteristics according to synthesis method
Particle
characteristics for
synthetic method
Mechano-
chemical
Liquid Vapour
Minimum primary
crystallite size
5nm 5nm 5nm
Particle size
distribution
Narrow Wide Wide
Degree of
agglomeration
Low Intermediate High
Purity Controllable High Very high
The mechanochemical method allows a unique and highly flexible means to fabricate
materials for the fabrication of an intermediate temperature, direct methane solid
oxide fuel cell. The most relevant materials for this purpose are the SDC electrolyte,
the NiO-SDC anode precursor and the cathode perovskite (SmSr)0.5CoO3. To assess
the performance of these materials the single-chamber fuel cell represents the simplest
laboratory scale method and is technically novel. The underlying philosophy is to
obtain and maintain as fine a microstructure as possible in the fuel cell electrodes with
a thin electrolyte. The nanoparticles characteristics obtainable from the
mechanochemical method are ideally suited to approach this challenge and to seek
improved SOFC performance.
39
3 Experimental
3.1 Powder Handling
Air-sensitive or toxic materials were handled within an MBRAUN argon glove-box.
Air-sensitive powders requiring XRD analysis were mounted on an air-tight sample
holder in the globe box.
3.2 Powder Synthesis
3.2.1 Reagents
The basic reagents for all millings: SmCl3, Sm2O3, CeO2, Ce(OH)4, NiCl2, Ni(OH)2
and NaOH; were supplied by Sigma-Aldrich. The graphite used was ThermoPURE
LBG-73, supplied by Superior Graphite Co.
“Ceria pre-mill” was obtained from Advanced Powder Technology, Ltd. (Australia) as
65wt% Ce(OH)4 as-milled in NaCl. It was prepared by the attrition milling of
Ce(OH)4 with NaCl for 30 mins at 400 rpm and if heat-treated for 30 mins at 500°C it
produces 5 nm CeO2.
Samples of SDC and SDC-SSC, a 400g sample of nanoscale, coprecipitated 21.1 wt%
Sm2O3-CeO2 were obtained from Advanced Materials Resources (UK). A 20 g
sample of nanoscale, combustion-chemical vapour deposited 10 wt% Sm0.2Ce0.8O1.9-
(SmSr)0.5CoO3 was purchased from MicroCoating Technology (USA).
40
3.2.2 Mechanochemical Processing
The majority of preparations were performed using a SPEX 8000 mixer/mill with a 5
g charge and 100 g of 9.5 mm stainless steel grinding media. The action of the
Mixer/Mill is illustrated in Figure 3.1. The agitation is achieved by motion in three
mutually perpendicular directions. The main component is a swing through about 3.2
cm at the end of a 8.9 cm arm through an arc of approximately 40°. Within a single
cycle, there is an additional vertical displacement of 0.48 cm and a horizontal
displacement of 1.43cm. This movement is cycled at about 1200 rpm and in the 40
mm wide and 60 mm long milling container (hardened steel) there are multiple ball-
ball and ball-wall collisions. The energy associated with these impacts fractures and
welds materials and in most instances induces chemical reactions. Mixtures were
typically loaded in the glove-box, and at the completion of ca. 4 hrs milling, were
transferred to a porcelain crucible for heat-treatment in air, if necessary.
Figure 3.1 Stylised SPEX mill operating
41
For the preparation of larger quantities a custom-built 1.8 L attrition mill, with a
capacity of 250 g of charge and 6 kg of 6 mm stainless steel grinding media was
utilised. This was operated with a water cooled jacket at speeds of 300-400 rpm.
After sieving to separate the as-milled powder from the milling media, additional
NaCl was added and milled for a short period to assist in product removal from media
and mill lining.
Removal of the reaction by-product and soluble salt phase was effected by washing
with distilled water. Heat-treated samples were transferred to polystyrene or
polypropylene bottles that were filled with distilled water, mixed thoroughly and
centrifuged. The supernatant liquor was decanted and the pH and salinity levels were
logged. The residue was dislodged, diluted and subject to ultrasonic agitation until no
large clumps were visible. As the washing cycles progressed a stable suspension
would form that would never entirely settle even after prolonged centrifuging. In
these cases, drops of 3 mol/L ammonium acetate were added until flocculation was
observed. This serves to increase the ionic strength of the solvent, thus encouraging
flocculation, but not adding impurities to the ceramic since this reagent is volatile. In
systems where the flocculant was not added, the salinity was below 10 ppm after 5
washes. When transferring the slurry to an ethanol system, the sonication and
centrifuge procedures were performed twice. Residues were air dried at 60°C or
150°C before final crushing in an agate mortar and pestle.
3.3 Fabrication Techniques
The fabrication of fuel cells was an evolutionary process and since their performance
is so heavily governed by the processing techniques applied they are treated
42
specifically in Chapter 7. In general, the 55 wt%NiO-SDC anode precursor was
compacted in a 12 mm die. This was coated with SDC using a Paasche VL-150
airbrush and the bilayer was sintered. The cathode was applied by airbrush over an 8
mm mask onto the electrolyte and fired at 900°C for 2 hrs. Spin-coating was
performed on a custom-built instrument operating at 3000 rpm.
3.4 Characterisation Techniques
3.4.1 X-ray Diffraction (XRD)
XRD analyses of bulk powder or sintered samples were performed on either a
Siemens D5000 or D500 diffractometer with Cu-Kα radiation. Diffraction patterns
were obtained with a step size of 0.04° and a rate of 3.4 sec/step. Scans for lattice
parameter and crystallite size determinations were run from 25-91° 2θ at 30 sec/step.
Peak positions and profiles were extracted using WinFit!(Beta version 1.2.1) [91]
using Pearson VII line profiles. The Scherrer equation [92] was used to calculate the
average crystallite size by determining the peak full-width at half maximum (FWHM)
for well resolved peaks. The crystallite sizes obtained were averaged for each
diffraction pattern. The instrumental line broadening was measured using 99.9%
Aldrich CeO2 as a standard after it had been calcined for 20 hrs at 1200°C. Accurate
lattice parameters were determined using the Nelson-Riley extrapolation [93], which
involved a graphical procedure. The instrument precision was monitored periodically
using an external quartz (SiO2) standard. Internal standards were used periodically for
powder samples but no significant deviations were observed when compared to
externally standardised experiments.
43
3.4.2 Specific Surface Area Measurement
Micromeritics Gemini 2360 or TriStar Surface Area Analysers were used to determine
specific surface areas using a five-point BET method. The equivalent spherical
particle size, referred to hereafter as the “BET particle size”, was calculated from the
equation D = 6/Sρ, where D is the particle diameter, S is the specific surface area and
ρ is the density.
3.4.3 Electron Microscopy
Particle morphologies and size distributions were assessed by transmission electron
microscopy using a Phillips 430 at 300 kV or a JEOL 2000FX at 80 kV. Samples of
slurry were dispersed in alcohol, using an ultrasonic bath before, pipette transfer onto
holey-carbon coated copper grids. High-resolution and energy filtered transmission
electron microscopy was performed on a JEOL 3000F microscope equipped with an
EDS detector and a Gatan Image Filter system. Energy filtered data was acquired and
manipulated using Gatan Digital Micrograph 3 software.
Scanning electron microscopy was performed on a JEOL 3000F (high resolution, 3
kV) or a 6400 (15 kV), or a LEO VP-FEGSEM from 1-15 kV.
3.4.4 Thermal Analysis
Simultaneous thermogravimetric-differential thermal analysis (TG-DTA) and
differential scanning calorimetry (DSC) were performed on a TA Instruments
SDT2960 in alumina pans under 100 mL/min of flowing air, or other specified gas.
44
3.4.5 Dilatometry
Shrinkage profiles were recorded on a Netzch 402EP dilatometer/furnace interfaced to
a PC and acquired using LabView 4 (National Instruments) software. Specimens
were typically compacted in a 5 mm tool steel die and were placed inside an alumina
collar with alumina spacers to limit cross contamination, particularly for Ni-
containing specimens. Data was represented as relative shrinkage (ΔL/Lo) vs
temperature or time, and the differential rate of shrinkage with respect to temperature,
d/dT(ΔL/Lo). Accurate densities were determined from compacted specimens by their
geometry and also the Archimedes method.
3.4.6 Gas Permeation Measurements
A cylindrical sample holder was fabricated from brass to allow flow through the
thickness of pressed, sintered and reduced cermet specimens. The pressure drop was
measured from the gauge and the flow rate recorded with a soap-film flowmeter as
illustrated in Figure 3.2. For a specimen with dimensions of A [m2] * l [m], the
permeability K [mol.m/m2Pa.s] for a flow rate r [mol/s] and pressure drop δp [Pa] is
given by, K = rl/(δpA).
P
Specimen: A [cm2], l [cm]
Flowmeter: r [mol/s]
Specimen holderP
Specimen: A [cm2], l [cm]
Flowmeter: r [mol/s]
Specimen holder
Figure 3.2 Schematic of permeability test setup
45
The formula quoted in Australian Standard physical test method for the permeability
to gases for refractories and refractories (AS 1774.7) was also used.
3.4.7 Conductivity and Fuel Cell Performance Apparatus
3.4.7.1 Total Conductivity Measurements
For the measurement of electrical conductivity the four-point DC method [94] was
implemented. Rectangular specimens of around 1.0 mm2 cross sectional area,
sectioned using a diamond saw from the sintered and polished cylindrical specimens,
were typically 10 mm in length. Either Pt or Ag paint was applied in four strips, the
innermost separated by a sticky-tape mask that was removed after the paint had dried.
Ag wire was tied around each contact and more paint was applied to ensure contact.
The separation between the innermost contacts was measured with an objective
micrometer as shown in Figure 3.3.
Figure 3.3 Digital photograph of a typical sample prepared for 4-pt conductivity measurements.
Electrical contact achieved with Ag wire and Pt paste
46
Constant current was applied across the outermost contacts with a Keithley 220
Programmable Current Source and the potential drop across the innermost terminals
measured with a Keithley 182 Sensitive digital voltmeter. The polarity of the current
was reversed to account for the thermoelectric effect and a number of readings were
taken to obtain an average.
3.4.7.2 Design and Evaluation of Testing Rigs
It was necessary to design and fabricated an apparatus that could be used for both
conductivity and fuel cell measurements, both of which require controlled
atmospheres and temperatures up to 900°C. The operation of the fuel cell in single
chamber mode simplifies the design, since only 2 of 4 leads need to be used. A
quartz-glass tube was fabricated with four tungsten carbide (WC) rods protruding into
the central region, as shown in Figure 3.4. Since WC was found to slowly oxidise at
operating temperatures, Pt wire was welded onto the ends to facilitate contact with
sample leads. The external ends were coated with Ag-solder to assist in electrical
contact with instrument leads. A K-type thermocouple was placed as close as possible
to the WC/Pt join.
47
Figure 3.4 Flow cell used for conductivity and fuel cell performance measurements. The thermocouple
is located at A, leads to specimen are crimped onto Pt wires at B
Conductivity specimens were allowed to rest on the glass base, fuel cell trilayers were
placed between plates of porous alumina and held with Ag leads on either face with a
small steel thumbscrew/clamp as shown in Figure 3.5.
Figure 3.5 Fuel cell clamped between porous alumina with attached leads
3.4.7.3 Fuel Cell Performance Measurements
The atmosphere was controlled using a three-way gas switch connected to cylinders of
argon, air, 10%H2-Ar or 6%CH4-6%O2-88%Ar pre-mix (Fuel Cell Gas, or FCG)
48
supplied by Air Liquide. For operation from mixtures of methane and air, two non-
return valves were attached to a Y-joint on one side of the flow cell. This allowed the
atmosphere to be switched from inert to methane before diluting the methane to the
desired flowrate with air, in accordance with local safety regulations, since increasing
the methane concentration from lean to rich crosses the explosive limit [95].
Open cell voltages were measured with a digital voltmeter; fuel cell power outputs
were measured according to the circuit shown in Figure 3.6. Galvanic current
interruption measurements were captured using a HP 54522A Digital Oscilloscope
according to the method of Badwal [96].
Figure 3.6 Circuit diagram for measuring fuel cell output. Battery element is the fuel cell test specimen
49
4 Synthesis and Characterisation of Nanopowder
Precursors
4.1 Sm-Doped CeO2
4.1.1 Mechanical Alloying of Sm2O3 with CeO2
It is known that that ion-conductive (Sm2O3)0.2(CeO2)0.8 can be obtained by the
mechanical alloying of CeO2 with Sm2O3 for 8 hours [81]. This procedure was
replicated at the specified stoichiometry of 20 mol% Sm2O3. After SPEX milling for
4hrs a grey/brown powder was obtained with a lattice parameter, ao = 5.453 Å, a
crystallite size Dx = 14nm and BET particle size of 56 nm, which suggests that the
primary crystallites exist as agglomerates of a few crystals. The addition of a diluent
phase such as NaCl assists in reducing such agglomeration [86] and it was similarly
effective in this instance. Diluting the oxide to 10vol% in NaCl and milling for a
further 4 hrs reduced the BET particle size to 21 nm, but TEM examination in showed
that, even after milling with 90 vol% NaCl, the powder consists of 7-70 nm
crystallites agglomerated in ca. 200 nm clusters (Figure 4.1).
Figure 4.1. TEM Micrograph of Mechanically Alloyed (Sm2O3)0.2(CeO2)0.8
50
At a stoichiometry of 20.2 wt% Sm2O3 a lattice parameter of 5.4412 Å was obtained
which is consistent with the findings of Eguchi [30]. The lattice parameter was
observed to increase to this value over the 8 hour milling period, indicating the
formation of a solid solution in agreement with other research shown in Table 4.1;
chemical (precipitation/sol-gel), thermal (extended mixed oxide powder calcination)
or mechanochemical.
Table 4.1. Comparison of Literature and Industry Values for the Lattice Parameters of the Sm-CeO2
system
Source Method Composition, x in
SmxCe1-xO2-x/2
Lattice
Parameter, Å
Eguchi[30] Thermal 0.20 5.441
This work Mechanical alloy 0.20 5.4412
Vitlov-Audino[81] Mechanical alloy 0.33 5.4386
Eguchi[30] Thermal 0.10 5.4312
Eguchi[30] Thermal 0.33 5.450
This work Mechanochemical0.33 5.4530
Huang[97] Chemical 0.20 5.4313
Balazs[98] Thermal 0.20 5.432
Advanced Materials
Research (Ind.)
Chemical 0.21 5.4369
Rhodia (Ind.) Chemical 0.20 5.430
When these lattice parameters are plotted against the samarium content as in Figure
4.2, significant scatter is evident. This suggests that caution should be exercised when
comparing lattice parameters obtained from different studies, particularly when the
method use to determine the parameter was not explicit in the literature. Ideally the
Nelson-Riley extrapolation should be performed, for this is the only means that the
systematic errors can be eliminated and an accurate lattice parameter obtained. All
51
lattice parameter determinations were performed by using the same procedure, from
diffraction patterns obtained from the same diffractometer.
5.405
5.41
5.415
5.42
5.425
5.43
5.435
5.44
5.445
5.45
5.455
5.46
0 0.05 0.1 0.15 0.2 0.25 0.3 0.35[x], in Ce1-xSmx
a, [A
ngst
rom
]
Figure 4.2 Lattice parameter vs Sm content in SDC
Given the hardness of the reagents and sample discolouration it is likely that the
mechanical alloying method caused significant abrasive wear from the milling media
and vial. The degree of iron contamination was found to be 2.3 wt% by ICP. It is not
known to what degree the iron may incorporate itself in the ceria solid solution and
influence the lattice parameter or sintering properties, Fe contamination was deemed
too high to consider this synthetic route appropriate.
4.1.2 Mechanochemical Synthesis of Sm0.2Ce0.8O1.9-x
Previous research [89, 90] suggests that nanoparticulate rare-earth hydroxides are
obtainable by the mechanochemical reaction of the rare-earth chlorides with a
stoichiometric amount of soluble hydroxide and a suitable amount of a soluble diluent
52
phase, typically a 90% volume fraction of NaCl. Therefore, Sm2O3 should be the
recoverable from the mechanochemical reaction between SmCl3 and NaOH. Brief
experiments confirmed showed that in the XRD there were a number of unidentifiable
reflections appearing in the pattern, presumably due to other reactions between SmCl3,
H2O and CO2. The formation of nanoscale Sm(OH)3 crystallites within a NaCl matrix
is likely and it was believed that this hydroxide would not be as abrasive when milled
with 5 nm Ce(OH)4.
To achieve the formation of nanoscale powders it has been shown necessary to keep
the product phase between 10-20 vol%. There is no data for the density of Sm(OH)3
so calculations were based on the density of Sm2O3. The reactants and ultimate
products formed are shown in Equation 14.
0.2SmCl3 + 0.8Ce(OH)4 + 0.66NaOH +
1.54NaCl
Sm0.2Ce0.8O1.9 + 7.26NaCl +
0.06NaOH Equation 14
The order of reagent addition to the mill was varied according to the following
regimes:
• Mill SmCl3/NaOH/NaCl, then add Ce(OH)4/NaCl and balance of media
• Mill SmCl3/NaOH/Ce(OH)4, then add NaCl and balance of media
• Mill all reagents together
There were no differences in surface area, lattice parameter or crystallite size found
between the powders obtained using these schemes, so the simplest was chosen. XRD
analysis of the as-milled reaction mixture showed the absence of discrete samarium
phases such as Sm(OH)3, indicating that it was amorphous or poorly crystalline.
53
Heat treatment at 550°C for 30 minutes yielded a powder with a lattice parameter of
5.4287 Å with crystallite and BET particle sizes of 6nm. Extending the milling time
had no observable effect on the lattice parameter, which was deemed to be too far
below the literature value of 5.441 Å to confirm the formation of a complete solid
solution. Heat treating the as-milled powder a further 2 hours at 700°C caused the
lattice parameter to increase to 5.4342 Å with a slight rise in crystallite and particle
sizes. These dependencies are displayed graphically in Figure 4.3. Chemical analysis
by ICP showed an iron contamination of less than 1000 ppm in the product heat
treated for 2 hours at 700°C, having a specific surface area of 48.7 m2/g, consistent
with a 17 nm particle and crystallite size.
4
6
8
10
12
14
16
18
5.424
5.426
5.428
5.43
5.432
5.434
5.436
5.438
5.44
500,1 600,1 700,1 700,2
BET
Par
ticle
Siz
e [n
m]
XR
D C
ryst
allit
e S
ize
[nm
]
Latti
ce P
aram
eter
[Ang
stom
]
Anneal Temp. [deg C], Time [hrs]
Figure 4.3. Development of mechanochemically milled SDC lattice parameter and particle/crystallite
sizes with annealing temperature and time
54
The dependence of the lattice parameter on the intensity of the heat treatment suggests
that the milling serves primarily to intimately mix the cerium and samarium hydroxide
particles, but is not sufficient to form a complete solid solution. Thermal treatment
then allows the increased diffusivity to facilitate formation of a homogeneous solid
solution.
Initial TEM investigations showed that the crystallites are well separated and have a
range of particle sizes between 10-30 nm. Figure 4.4 shows a high resolution image
of two particles hanging ever the edge of the amorphous carbon film. The small
particles size and the clear separation of Ce and Sm M-edges and render the sample
ideal for energy filtered TEM analysis [99].
Figure 4.4. TEM Lattice image of mechanochemically synthesized SDC
In Figure 4.5 the Cerium, Samarium and Oxygen elemental maps show a density
distribution that is both uniform and comparable to thickness variations on the
thickness map (not shown). This suggests that the distribution of cerium and
samarium throughout the bulk nanoparticles is constant, at an amount, determined by
55
EDS at 20 mol% Sm. There was no evidence for Fe, Na or Cl contamination based on
EELS analysis.
Figure 4.5. EFTEM Elemental maps and reference image of SDC Nanopowder. From bottom left: Ce-
M edge, Sm-M edge, O-K edge. All micrographs have the same field of view.
As shown previously mechanical alloying techniques, on their own, are not
appropriate for synthesizing nanoparticulate samarium-doped cerium oxides due to
iron contamination. The application of mechanochemical milling techniques to the
reaction between SmCl3, NaOH and nanoscale Ce(OH)4 in a NaCl matrix, yields an
intimate mixture of products, allowing the formation of a solid solution at moderate
heat treatment temperatures and with negligible Fe contamination compared to other
milling methods. The resulting small particle size suggests that this material may be
useful for the fabrication of thin film electrolytes and nanostructured electrodes.
56
4.1.2.1 Mechanochemical Synthesis of SDC Nanopowders in an Attrition
Mill
To obtain larger quantities of SDC it is necessary to perform the mechanochemical
process in the volumes afforded by an attrition mill. As discussed in Section 3.2.2, the
1.8 L horizontal attrition mill was used with 6 mm media. The grinding characteristic
of an attrition mill differs from the attrition/impact mixed mode that is delivered by a
SPEX mill. The impact of these differences becomes evident in the microstructures,
physical characteristics and impurities of the products. The mechanochemical
synthesis of SDC and the treatment of commercial nanocrystalline SDC will be
discussed.
4.1.2.1.1 Attrition Milling of Commercial SDC
For comparative purposes, chemically precipitated SDC, actually 21.2 wt%Sm2O3-
CeO2 or Sm0.21Ce0.79O1.9, supplied by Advanced Materials Resources was milled in an
attrition mill at a 15 vol% relative to NaCl for 80 mins at 400 rpm (68.5 g SDC, 181.5
g NaCl, 6 kg media). As a reference the same proportions were milled in a SPEX vial
for 4 hrs. Figure 4.6 shows the change in powder morphology as received (left), to
SPEX (a) and Attrition milled (b) as a powder with an XRD crystallite size of 5 nm.
The as-milled powders were heat treated for 30 mins at 500°C.
57
Figure 4.6. AMR SDC as received and after (a) 4 hrs SPEX milling, (b) 80 mins attrition milling; both
received heat treatment for 30 mins at 500°C
The XRD crystallite size was the same in all specimens and there was nothing to
distinguish the two processing methods from one another or from their starting
material. Attrition milling was slightly less effective in reducing the BET particle size,
achieving 8.5 nm particles compared to 6 nm from SPEX milling. Notable to both is
the persistence of 250 nm+ dense agglomerates, together with some looser, flat
agglomerates of ca. 65 nm diameter. Neither the SPEX nor attrition milling processes
were sufficient to disintegrate these hard agglomerates. The material, presumably
synthesized via a nitrate sol-gel route had received calcination at 400°C and this
temperature appears to have induced oxide bridging between crystallites, forming the
micron sized agglomerates. These harder agglomerates are unable to be broken down
further by SPEX milling action.
58
Heat treatment for 2hrs at 700°C yielded 27 nm crystallites with a BET particle size of
40 nm, which is unsatisfactory when compared to the 17 nm particle and crystallite
sizes obtained above. The transmission electron micrograph in Figure 4.7 shows 6-40
nm primary crystallites in various states of agglomeration. The presence of dense sub-
micron agglomerates in the upper-right of the image is undesirable and suggests that
the material does not respond well to size reduction by mechanical attrition.
Figure 4.7 SPEX Milled AMR SDC heat treated 2 hrs at 700°C
4.1.2.1.2 Mechanochemical Synthesis of SDC in an Attrition Mill
Agglomeration could not be eliminated by attrition milling of commercially supplied
SDC, so to obtain a truly nanoscale product it was necessary to perform the reaction
from Equation 14 at a larger scale in a horizontal attrition mill. Since the mill can not
be perfectly sealed there was a risk that the Sm source, SmCl3, being air/moisture
sensitive could decompose by reaction with atmospheric moisture during milling and
preclude the formation of a solid solution. Furthermore, moisture absorption may
manifest itself in excessive powder sticking to mill walls or media, resulting in a
lower yield and possible chemical inhomogeneity. However, since the attrition mill
59
rapidly generates heat, it is likely that the reaction between NaOH and SmCl3 will
occur before hydration could occur.
A milling of 95.89 g of 65wt% Ce(OH)4/NaCl and 125 g of NaCl was loaded into an
attrition mill with 6 kg of 6.4mm steel balls. In the glove box, 9.9 g of NaOH was
mixed with 19.23 g of SmCl3; this was promptly transferred and milled for 60 mins at
400 rpm. The product was recovered from the mill and stored in a sealed
polyethylene (Zip-Lock) bag in an oven at 35°C. There was no excessive sticking of
powder to the mill walls, attrition shaft, rotors or media hence atmospheric moisture
was not a problem. Heat treatment of a sample for 2 hrs at 700°C gave a single cubic
ceria phase with a lattice parameter of 5.4193 Å and BET and XRD sizes of 17 and
15 nm, respectively. The lattice parameter was too low, suggesting that an extensive
solid solution was not formed. Furthermore, a number of unassignable diffraction
peaks at ~14° and 23° 2-theta appended in the x-ray diffractogram, shown in Figure
4.8 suggests the presence of unreacted/undissolved Sm based compounds.
Figure 4.8 X-ray diffractogram of 60 min attrition milling of SDC precursors, anomalous peaks
suggesting unreacted SmCl3 indicated with arrows.
60
A lattice parameter of 5.4327 Å was obtained by re-milling the attritioned powder in a
SPEX vial, and heat treating for 2 hrs at 700°C. Subsequent heat treatment of the
attrition milled powder did not give an increased lattice parameter – an additional 2.5
hrs heat-treatment at 750°C still gave a = 5.412 Å and the anomalous low-angle peaks
persisted. It appears, therefore, that more milling was required.
A second attrition milling was performed for 80 min; heat treatment for 2hrs at 700°C
gave a product with lattice parameter of 5.4316 Å, XRD crystallite size of 13 nm and
BET particle size of 14 nm. The TEM micrograph in Figure 4.9 shows 5-35 nm
crystallites in loose 25-130 nm clusters.
Figure 4.9. TEM Micrograph of mechanochemically synthesized SDC, milled twice (80mins +
80mins) in a horizontal attrition mill, heat treated at 700°C for 2hrs
After the first 80 mins of attrition milling the SDC product contained 430 ppm Fe, the
second 80 minutes of milling increased this to 990 ppm Fe. By comparison, the SPEX
re-milling of 80min attrition milled powder increased the level to 12800 ppm Fe.
Magnetic separation reduced the iron content in the attrition milled material to 860
61
ppm. Evidently the less intense attrition grinding action leads to reduced abrasion
from the media and the mill than the impact-type action in a SPEX mill, but the iron
contamination introduced from either method can easily be minimised by magnetic
separation.
4.2 Mechanochemical Synthesis of NiO-SDC Composite
Powders
Given the contamination observed for the mechanical alloying of CeO2 and Sm2O3, it
was not prudent to pursue the milling of NiO with SDC to synthesise nanoparticulate
NiO-SDC. Hydroxides precursors were used during SPEX millings instead. The
optimum cermet concentration tends to lay around 50 vol%Ni-SDC, so three NiO-
SDC compositions were selected to cover this composition range so that accurate
conductivity measurements could be performed.
Mixtures to yield nanopowder compositions of 35 wt%, 50 wt% and 65 wt% NiO-
SDC (subsequently referred to as x% NiO; x=35, 50, 65) were mechanochemically
synthesized in SPEX vials from mixtures of Ni(OH)2, SmCl3, NaOH, Ce(OH)4 and
NaCl. Figure 4.10 shows the diffractogram for 50%NiO, typical for any of the
materials in the series. Table 4.2 provides a summary of characterisation data for the
NiO/SDC composites showing the XRD crystallite size for NiO and SDC, the SDC
lattice parameter and BET surface area. Since the specific densities of SDC and NiO
are similar at 7.35 and 7.45 g/cm3 respectively, the mean equivalent spherical particle
size was calculated using the composite’s theoretical density.
62
Figure 4.10 XRD Pattern of 50 wt%NiO-SDC nanocomposite, NiO peaks indicated with arrows
Table 4.2. Summary of BET and XRD characterisation data for NiO/SDC nanopowders
Composition Theor.
density,
g/cm3
BET
Surface
Area, m2/g
Mean
Particle
size, nm
NiO
Crystallite
size, nm
SDC
Crystallite
size, nm
SDC
Lattice
Parameter,
Ǻ
35%NiO 7.38 39.32 20.7 12.5 14.4 5.4340
50%NiO 7.40 32.96 24.6 13.6 11.4 5.4315
65%NiO 7.41 27.69 29.2 16.2 12.3 5.4298
The fact that the BET particle sizes are approximately twice as large as the crystallite
sizes suggests that slight agglomeration occurred, which was confirmed by TEM
examination. The degree of agglomeration was also seen to increase slightly with
increasing NiO content and its crystallite size. Some inter-particle agglomeration
during heat-treatment and drying was evident. Figure 4.11 shows a series of energy
63
filtered TEM elemental maps indicating separation of Ni and Ce rich regions. These
regions correspond to individual NiO and SDC crystallites.
64
Figure 4.11 Energy filtered transmission electron micrographs of 50%NiO-SDC nanopowder: (Top-
left) Unfilitered, (Top-right) Oxygen map, (Bottom-left) Cerium map, (Bottom-right) Nickel map
The size of the crystallites correlates well with the data in Table 4.2. The elemental
map for Sm was too noisy due to be meaningful, however results in Section 4.1.2
showed that complete solid solution exists based on lattice parameters. This
nanoscale dispersion of NiO and SDC crystallites was observed in the elemental maps
for all the three compositions studied.
4.2.1 Mechanochemical Synthesis of NiO-SDC in an Attrition Mill
While the optimum Ni content for the cermet remained unknown, a batch of 35
wt%NiO was attrition milled for compaction and sintering experiments. The reagents
(SmCl3, NaOH, Ni(OH)2, Ce(OH)4 and NaCl) were attrition milled at 400 rpm for 140
65
mins. Heat treatment of the as milled material for 2 hrs at 700°C yielded a black
powder with a SDC lattice parameter of 5.4216 Å, 13 nm SDC crystallites, 17 nm NiO
crystallites and a mean BET particle size of 23 nm. The morphology of these crystals
was mixed, predominantly 9-16 nm equiaxed crystals. Despite the comparatively low
lattice parameter, there were no anomalous reflections at low angles in the
diffractogram, suggesting that there were no unreacted Sm compounds. Heat
treatment at 1000°C for 10 hours gave a lattice parameter of 5.4327 Å indicating that
the desired solid solution formed during sintering.
Concerns that 140 mins may be too long a milling time and may encourage the
formation of hard agglomerates, as reported in the La0.7Ca0.3MnOz system [100], were
assessed by performing a 1hr mill in the SPEX. Heat treatment was identical but the
product had anomalous low-angle (14° and 22° 2θ) XRD peaks and a lattice
parameter of only 5.4184 Å. A BET mean particle size of 19 nm correlates with the
17 nm SDC and 23 nm NiO crystallite sizes from XRD but TEM showed a wide range
of particle sizes, 8-40 nm crystals in 50-300 nm agglomerates. The density of the
large agglomerates as shown in Figure 4.12 and the diffraction data suggests that a 1
hour SPEX milling is insufficient. Attrition milling for 140 mins or 4 hours in a
SPEX mill were necessary to obtain a product that was nanoscale and had the correct
lattice parameter.
66
Figure 4.12 35 wt%NiO-SDC milled for 60 mins in SPEX showing wide particle size distribution and
hard agglomerates
As will be seen in Section 6.2.2, a composition of 50 vol%Ni-SDC was chosen for the
anode substrate. The relevant oxide precursor is 55 wt%NiO-SDC, which cannot
readily be distinguished microstructurally from any of the previous cermet precursors.
The lattice parameter, crystallite size and surface area were monitored over the course
of milling to ensure a complete solid solution is formed in the SDC, a concern due to
the higher dilution of the SmCl3 in this system. A milling time of 240 mins was
necessary to achieve the lattice parameter of 5.421 Å, as shown in Figure 4.13, after
heat treatment for 2 hrs at 700°C.
67
5.408
5.41
5.412
5.414
5.416
5.418
5.42
5.422
10
12
14
16
18
20
22
24
0 50 100 150 200 250
SDC Lattice Parameter [Angstroms]SDC Crystallite Size [nm]
NiO Crystallite Size [nm]
[Ang
stro
ms]
[nm]
Milling Time at 400rpm [mins]
Figure 4.13 Plot of lattice parameter and crystallite size against attrition milling time for the
55wt%NiO-SDC reaction mixture. Heat treated for 2 hrs at 700°C.
The specific surface area of the recovered powder was relatively constant throughout
at 41 m2/g, which suggests that the equivalent 20 nm spherical particles consist of
lightly agglomerated 12-16nm crystallites.
4.3 Synthesis of (SmSr)0.5CoO3
The doped perovskite (SmSr)0.5CoO3 (SCC) is a popular cathode material for
intermediate temperature SOFCs. Prepared conventionally by the calcination of
Co3O4, Sm2O3 and SrCO3 at 1200°C for 6 hrs, the material has a crystallite size of ca.
50 nm and a cubic lattice parameter of 3.809 Å from the diffraction pattern shown in
68
Figure 4.14. This lattice parameter was within experimental error of the literature
value of 3.795 Å [101].
0
20
40
60
80
100
20 30 40 50 60 70 80 90
Inte
nsity
2-Theta
Figure 4.14 X-ray diffractogram of (SmSr)0.5CoO3
Calcination at 1200°C gave small BET surface areas and TEM examination revealed
large agglomerates of micron sized crystallites that were not electron transparent.
When the material was synthesized mechanochemically from chloride reactants
according to Equation 15, the desired phase was not formed until the temperatures
exceeded 1000°C. This resulted in severe agglomeration since the molten NaCl phase
allows inter-particle contact and growth.
SrCl2 + 2CoCl2 + SmCl3 + 9NaOH +
9.2NaCl + 3/4O2
2(Sr0.5Sm0.5)CoO3 +18.2NaCl +
9/2 H2O Equation 15
69
The as-calcined material from Equation 16 was milled with NaCl to attempt to reduce
its particle size, but it was not possible to obtain particles less than 4 microns in size.
It was apparent that the SCC was thermodynamically unstable since CoO and SrCoOx
phases were observed in its XRD pattern after milling.
2SrCO3 + Sm2O3 + 4CoO 2(Sr0.5Sm0.5)CoO3 Equation 16
Since the mechanochemical route failed to produce nanoscale or ultra-fine material,
SCC required for future experiments would have to be obtained by following literature
methods [101] and hence would have micron sized particles. Alternatively,
researchers at MicroCoating Ltd. supply ultra-fine SCC-SDC composites via a
combustion flame vapour synthesis route [78], and given reports of its suitability in a
fuel cell (Section 2.3.1.2) it reduces the number of variables for this project and hence
was also tested as a cathode.
4.3.1 Characterisation of Commercial 10%SDC-SCC
The material received differed from the 154 m2/g (SmSr)0.5CoO3 shown in Figure
4.15(left); the micrograph on the right [102] shows that the cathode material consists
of 7-17 nm crystallites agglomerated into ~750 nm particles with some reaching 1.5
μm. The black powder was exceptionally fine and free flowing having a surface area
of 65 m2/g.
70
Figure 4.15 Transmission electron micrographs of SCC (left) and 10%SDC-SCC (right) supplied by
MicroCoating Ltd.
The XRD pattern of the as received material, shown in Figure 4.16, could be
positively indexed for two SDC reflections only, which give a crystallite size of 46 nm
and lattice parameter of 5.456 Å. After heat treatment for 3 hrs at 850°C the cubic
(SmSr)0.5CoO3 phase was observable in Figure 4.16 and had a lattice parameter of
3.73Å.
The un-indexable phases indicated by question marks were presumed to be product(s)
from the reaction with the atmospheric moisture or CO2. The powder was not stored
under an inert atmosphere and whilst these compounds could convert back to the
desired parent compound by sintering in this instance, at later stages the
decomposition products such as cobalt oxide would persist and lead to problematic
cathodes, to be discussed in following sections. The stability of such a material with
fuel present in the atmosphere of a single-chamber SOFC has yet to be determined.
Heat treatment at 1200°C for 3-6 hrs recovered the doped perovskite and reduced the
large surface area but gave an ostensibly more stable and consistent material.
71
Figure 4.16 MCT 10%SDC-SCC as received (bottom) and after calcination at 350°C for 2hrs (top).
Squares denote SCC reflections, circles SDC, question marks denote un-indexable reflections
72
5 Processing of Mechanochemically Synthesized
Nanopowders
An initial investigation was undertaken to assess whether it was possible to fabricate
components from mechanochemically synthesized nanopowders using tape casting
techniques. Tape casting is the preferred method for the fabrication of the supporting
component in a planar solid oxide fuel cell. It is a far more economical process than
compaction as no dies are required; hence it is readily scalable for the fabrication of
large ceramic sheets. Since drying of nanopowders induces agglomeration, tape
casting may be a viable method to densify mechanochemically synthesized
nanomaterials without drying.
5.1 Slip Characterisation and Tape Casting of CeO2
Nanopowder
5.1.1 Slip Formulation
Ethanol was chosen as the solvent based on its compatibility with other organic
additives, low cost, ease of use and disposal. The ideal dispersant was selected by the
addition of 5% by weight, relative to the oxide equivalent of the five candidates
(DisperByk®-107, DisperByk®-111, BYK®-P-105, Hypermer KD-1 and Menhaden
fish oil) to samples of a 40 wt%, 25 nm CeO2/ethanol slurry. The slurry had an initial
viscosity of 400 mPa.s at a shear rate of 100 s-1. The slurries were stirred thoroughly
then ultrasonicated for 5 minutes in sealed plastic containers. Visual inspection was
sufficient to deduce that DisperByk®-111 gave the most dramatic reduction in
viscosity, significantly lower than the other four dispersants. A rheometry titration
73
was performed on the 40wt% CeO2 slurry with DisperByk®-111 (DB111), which the
supplier’s notes state is a polyester with a molecular weight around 1500. Mixing
between sequential additions of dispersant was achieved by roll milling for 30 mins.
Figure 5.1 shows that beyond a dispersant loading of 3.5 wt%, relative to the mass of
dry oxide, there was no significant decrease in the slip’s viscosity. This suggests that
the surface of the particles was saturated with dispersant molecules and any further
addition would not improve the degree of dispersion.
0100200300400500600700800900
0 2 4 6
DB111 Concentration [wt%]
Visc
osity
[mPa
.s]
8
Figure 5.1 Rheometry titration of 40wt% 25nm CeO2/ethanol with DisperByk®-111
This dispersant was so effective that addition directly to the wet centrifuged cake
yielded a low viscosity 58.5wt% CeO2 slip. This is an advantage inherent to the
nature of the mechanochemically synthesized starting material, as it is obtained from a
highly dispersed state in NaCl. Concentration during washing to a wet state means
that no hydroxyl bridges can form hard agglomerates, as is reported to occur during
the drying of other nanopowders [103, 104]. The pseudoplasticity of the slip is
evident by the shear thinning behaviour shown in Figure 5.2 which can be
74
characterised by a pseudoplasticity index, n = 0.29 in τ = Kγn where τ is the shear
stress, K is a consistency factor and γ is the shear rate.
Figure 5.2 Shear stress and viscosity vs shear rate for 58.5wt% 25nm CeO2/ethanol suspension
dispersed with 3.5% DisperByk®-111,
The suspension is therefore highly pseudoplastic as its stress-strain behaviour deviates
strongly from that of a Newtonian fluid. In comparison with published data regarding
the dispersion of 137 nm yttria-stabilised zirconia with menhaden fish oil (MFO) and
phosphate ester (PE) in ethanol/methyl ethyl ketone azeotrope [105], Table 5.1 shows
that on the basis of solids loading and viscosity, DisperByk®-111 has excellent
performance for the dispersion of ceria nanoparticles.
75
Table 5.1 CeO2 nanopowder slip properties compared to literature source [105]
System Solids
Loading,
wt%
Dispersant
Loading,
wt%
Viscosity
at 100s-1,
mPa.s
MFO 51 2 1160
PE 51 1 1110
This work 58.5 3.5 160
There are no rigorous scientific procedures to quantitatively determine the amounts of
binder and plasticizer that are required to obtain a slip suitable for tape casting.
Literature data on the tape casting of zirconia [22] showed typical binder contents in
the green tape of 17.1 vol% with the plasticiser content of 38.7 vol%. Starting from
these values, which did not yield coherent tapes, a suitable film-forming slip was
obtained. The formulation listed in Table 5.2 was used for the tape casting of
subsequent specimens. Upon addition of the binder to the dispersed slip the viscosity
increased drastically, making it difficult to achieve consistent mixing.
Table 5.2 Typical slip formulation
Component
Wt%
in slip
Vol%
in green
body
25nm CeO2 47.0 15.2
Butvar B79 7.0 69.0
Santicizer 160 3.8 7.9
PEG-400 1.7 3.4
DisperByk-111 1.8 4.6
Ethanol 38.6 0
Both silicone and wax coated PET substrates were tested and detaching the green
tapes intact from either surface was equally challenging. The addition of polyethylene
76
glycol, a Type-II plasticizer [66], assisted in delamination from the substrates. After
14 mm diameter blanks were punched from a 175 μm thick green tape it was possible
to detach the tape from substrate. The green density was 3.2±0.1 g/cm3, 45% of
theoretical. The microstructure of the green tape can be seen in Figure 5.3 and no
features larger than ca. 300 nm are discernible in its structure.
Figure 5.3 Electron micrograph of green 25nm CeO2 tape
It is reasonable to conclude that the degree of suspension in the slip is high, that is the
crystallites are loosely flocculated, allowing the plastic phase of the binder and
plasticiser to surround them fully. TEM of a diluted slurry sample shows (Figure 5.4)
ca. 150 nm flocculates of 6-37 nm crystallites of CeO2. A degree of is inevitable in
the highly concentrated conditions of a tape casting slip [105]. This applies even
more strongly to nanoparticles as they are more susceptible to flocculation by van der
Waals attractive forces [106].
77
Figure 5.4 Transmission electron micrograph of a diluted slip sample
5.1.2 Thermal Analyses and Sintering Behaviour
Thermal analysis (Figure 5.5) shows multi-stage mass loss of 22% from 25°C to
375°C with 3 maxima visible on the trace for rate of mass loss with respect to
temperature. These are due to various evaporation and decomposition stages of the
organic phases present in the binder.
78
Figure 5.5 TGA of green 25 nm CeO2 tape. Squares = weight loss, Circles = derivative weight loss
with respect to temperature
This mass loss can be correlated with a shrinkage of 5%, shown by dilatometry in
Figure 5.6, between room temperature and 310°C. The onset of sintering can be
observed at 600°C with a maximum shrinkage rate at 1000°C. The oscillations in the
trace at low temperatures are due to instrumental effects.
79
Figure 5.6 Dilatometry curves for green 25 nm CeO2 pellets. Squares = temperature program,
diamonds = shrinkage vs. time, circles = shrinkage vs. temperature
The dependency of grain size on sintering was assessed by powder x-ray diffraction
measurements. Figure 5.7 shows that the grain size after sintering for 1 hr at 1000°C
was 72 nm which suggests that to maintain nanoscale grains the sintering temperature
must be kept between 1000-1100°C.
80
Figure 5.7 XRD Crystallite size vs. sintering time at temperature for tape cast CeO2
The sintering profile for the green tapes was chosen with a slow ramp to 350°C to
allow the organics to burn out, then a ramp to 1050°C with a 3 hr dwell. The green
tapes were placed between porous alumina spacers to prevent curling which was
otherwise problematic. The microstructure for a sample sintered to 67% density under
these conditions shown in Figure 5.8(left) with a mean grain size of 450 nm. Figure
5.8(right) shows a 122 μm thick cross-section of the same specimen.
Figure 5.8 Electron micrograph of the sintered grain structure on the material’s surface (left); sintered
tape cross-section fracture surface (right)
81
Despite the intermediate density and relative thinness of the sintered specimens, they
were easily manipulated. Whilst it is difficult to envisage 100% dense nanostructured
electrolytes being produced via this method at present, the prospects for tape cast
anodes are promising. Porosity is detrimental to the performance of an ionically
conductive electrolyte, but it is essential for the performance of the anode where gas
diffusion must be facile. For an electrolyte in a single chamber fuel cell, the issue of
gas leakage does not exist so porosity is only detrimental if it is large enough to
influence the conductivity. Tape casting may be suitable, therefore, for the fabrication
of nanostructured anodes and electrolytes in a single chamber SOFC. However, for
the scale of this investigation, the material losses incurred during the slip processing
associated with high viscosity slurries led to the fact that tape casting was not chosen
as the processing method as compaction would appear far simpler.
5.2 Compaction of Nanopowders
This section discusses the compaction of various nanopowders, their densification
properties as measured on a dilatometer and the microstructures and densities obtained
after various sintering schedules.
5.2.1 SDC Electrolyte
Dilatometry measurements (Figure 5.9) shows two maxima in the sintering rate, at
830°C and 1140°C. The latter was used as the actual sintering temperature because it
was associated with a much higher shrinkage rate and yielded samples with structural
integrity; specimens removed after sintering at 830°C typically disintegrated during
removal from the furnace.
82
-25
-20
-15
-10
-5
0
-0.14
-0.12
-0.1
-0.08
-0.06
-0.04
-0.02
0
0.02
400 600 800 1000 1200 1400
Shr
inka
ge %
DL/
Lo
Der
ivat
ive
(T) S
hrin
kage
Temp [deg C]
Figure 5.9 Dilatometry curve for MCP-SDC
According to the rationale of Chen [5], sintering with a “spike” of 1200°C and a long
hold at 1000°C yields dense, nanostructured ceramics. This was not found to be the
case; the analysis giving a density of 72% although the XRD grain size was 28 nm.
Ultimately, 91% dense SDC was obtained after sintering for 12 hrs at 1300°C, or 88%
density after 4 hrs at 1350°C. The grain sizes, too large for a significant XRD
crystallite size determination, are visible in Figure 5.10 and have an average size of
980 nm.
83
Figure 5.10 SEM fracture surface of sintered MCP-SDC, 91% dense
By comparison, the AMR-SDC as received did not show good sinterability, the
density only reaching 67% after sintering at 1500°C, despite starting with a green
density of 55%. It was known that this material had received heat treatment at 380°C
and between room temperature and 800°C thermogravimetry showed a continual mass
loss of 11.3%. This, in combination with an uneven density distribution due to the
presence of large, hard agglomerates prevented complete densification.
From experiments described in section 4.1.2.1.1 it was clear that attrition milling of
AMR-SDC was unable to completely break down the agglomerates that existed in the
material as recieved. However, the milled and heat-treated nanopowder was far more
sinterable, achieving 93% density after sintering as shown in the dilatometer
measurement in Figure 5.11.
84
-20
-15
-10
-5
0
-0.1
-0.08
-0.06
-0.04
-0.02
0
0.02
0 200 400 600 800 1000 1200 1400 1600
Shr
inka
ge [%
DL/
Lo]
Der
ivat
ive
(T) S
hrin
kage
[d/d
T(D
L/Lo
)]
Temp [deg C]
Figure 5.11 Dilatometry curve for AMR-SDC, heat-treated for 2 hrs at 700°C
Tabulated data (Table 5.3) from Figure 5.9 and Figure 5.11 show that AMR-SDC has
a higher maximum shrinkage rate and a lower total shrinkage than MCP-SDC.
Despite having a smaller primary crystallite size, the persistence of large agglomerates
in attrition milled AMR-SDC suggests that this material would be inferior to
mechanochemically synthesized MCP-SDC for the fabrication of nanostructured
anodes or thin-film electrolytes.
85
Table 5.3 Densification data for AMR-SDC and MCP-SDC
Dilatometer Measurement AMR-SDC MCP-SDC
Temperature of Max.
Shrinkage Rate [°C]
1350 1140
Max. Shrinkage Rate
[%/°C]
-0.1 -0.12
Total Shrinkage (%) 20 24
Sintered Density (%) 93 88
5.2.2 NiO-SDC Anode Precursor
As discussed in section 2.1.1.1, the microstructure of the anode is of central
importance in the performance and optimisation of a SOFC. The high surface area of
nanoscale powders provides a large thermodynamic driving force and imparts a high
sinterability, but it is this fine grain structure that is desirable and should be retained
within the anode precursor’s sintered structure. Densification of nanopowders, whilst
maintaining a nanoscale grain structure, can be achieved by preventing grain-
boundary migration and enhancing grain-boundary diffusion, as shown by Chen [5].
This can sensibly be achieved by fabricating a composite structure in which neither
phase has a high diffusivity or solubility with each other, for example NiO and SDC.
Given the highly dispersed state of mechanochemically synthesized SDC and NiO
nanocomposites, they will naturally inhibit the growth of their interphase boundaries;
hence it should be much easier to obtain a nanostructured composite following an
appropriate sintering schedule. Figure 5.12 shows the dilatometry curves for
compacts pressed at the three compositions, heated at 300°C/hr to 1500°C and held
for 3 hrs.
86
Figure 5.12 Dilatometry curves for compacted NiO/SDC composites
The temperatures corresponding to the maximum sintering rates were extracted from
the derivative of shrinkage rate with respect to temperature. Table 5.4 presents this
data together with the green and sintered densities and the post-sintering crystallite
sizes, as determined by XRD. The compaction pressure was 250 MPa and the dwell
time at the sintering duration was 3 hours at the specified temperature. These
specimens were used for subsequent cermet formation and conductivity measurements.
Table 5.4 Sintering data for NiO/SDC compacts
Composition Temp. at
max.
sintering
rate, °C
Sintering
Temp., °C
Green
density, %
Sintered
density, %
NiO Cryst-
allite size,
nm
SDC Cryst-
allite size,
nm
35%NiO 1000 1000 50 66 46.9 34.6
50%NiO 1175 1200 48 86 41.6 35.5
65%NiO 1210 1200 73 87 42.7 34.8
87
It was apparent that a range of densities appropriate for cermets are achievable while
nanoscale grain size has been maintained. Polished and unpolished specimens were
examined under the SEM, shown in Figure 5.13. The dark and light regions on the
polished surface are due to backscattered electrons collected by the secondary electron
detector, such contrast occurring because the surface was flat. As CeO2 has 2.3 times
the atomic weight of NiO it is easy to visualise the generation of a backscattered
electron signal, even at the 3kV accelerating voltage used in the FESEM. In these
images the light regions are SDC rich and the dark regions are NiO rich.
Figure 5.13 Sintered 35%NiO composite pellet micrographs: (Left) Polished surface, (Right) Fracture
surface
Ion-thinned specimens were examined using EFTEM techniques and Figure 5.14
shows that the powder was a nanostructured composite of 68 nm NiO and 62 nm SDC
grains. The nanoscale separation of NiO and SDC grains as seen in Figure 4.11 was
preserved following sintering.
88
Figure 5.14 TEM reference image and EFTEM Elemental Map 75% dense 35wt%NiO-SDC: NiO =
Blue, SDC = Red
The 55wt% NiO-SDC synthesized previously (Section 4.2.1) had a sintering
behaviour indistinguishable to that of 50%NiO, having a maximum rate of shrinkage
at 1200°C. After sintering for 3hrs at this temperature the composite had a density of
85%, the SDC crystallites had grown to 61 nm and NiO to 77 nm. After sintering it
was found that the SDC lattice parameter had not changed significantly.
5.2.3 SDC-SCC Cathode
The 10%SDC-SCC cathode material supplied by MCT required densification for
conductivity determinations and fuel cell fabrication. The supplier specified a
sintering temperature of 900°C [102] which exceeds the maximum shrinkage rate
shown in Figure 5.15 by about 100°C. A second minimum in the differential
shrinkage rate can be seen around 900°C so the specified heating schedule is justified.
89
-30
-25
-20
-15
-10
-5
0
-0.14
-0.12
-0.1
-0.08
-0.06
-0.04
-0.02
0
0.02
500 600 700 800 900 1000 1100 1200 1300
Shrin
kage
[%D
L/Lo
]
Der
ivat
ive
(T) S
hrin
kage
[d/d
T(D
L/Lo
)]
Temp [deg C] Figure 5.15 Dilatometry curves for 10%SDC-SCC
The continuous shrinkage upto 1200°C is presumably accompanied by rapid grain
growth, given that the material synthesized at this temperature consisted of micron
sized grains. Thermogravimetry under 10mL/min of air showed a 10.5% mass loss to
800°C with melting occurring at 1337°C. For the fabrication of cathode layers the
combustion-chemical vapour condensation method of MCT appears to deliver a
nanomaterial that densifies at temperatures that will not cause densification of the
anode’s microstructure.
90
6 Electrical and Microstructural Characterisation of
Fuel Cell Components
6.1 SDC Electrolyte
Comparison of the conductivities of electrolyte candidates, AMR-SDC and MCP-SDC,
shows the effect that the synthesis and processing methods have on the material.
Figure 6.1 shows the temperature dependence of the conductivity for the two
electrolytes as measured by the four-point DC method in flowing air. Both showed
behaviour consistent with ionic conduction according to Equation 9.
Figure 6.1 Temperature dependence of the conductivity of doped ceria specimens
The conductivity of MCP-SDC is visibly higher than that of AMR-SDC over the
temperature range tested, and at 600°C it is more than double. In perspective this
value is still more than half the highest value reported for this material in the literature,
and as shown in Table 6.1, MCP-SDC is within range of reported data. The fact that
the activation energies were comparable to literature sources suggests that the lower
91
ionic conductivities measured were caused by incomplete densification. This is
because a small amount of porosity will not hinder the bulk charge transfer
mechanism, but would increase the internal resistance.
Table 6.1. Comparison of Conductivity Parameters for the Sm-CeO2 System
Source Activation
Energy,
eV
Conductivity
at 600°C,
S/cm
Eguchi[30] 0.78 0.02
MCP-SDC 0.81 0.009
Huang[97] 0.97 0.005
AMR-SDC 0.79 0.004
The objective in this study was to of fabricate thin film electrolytes for intermediate
temperature SOFCs. As discussed by Steele [9] the electrolyte should contribute, at
most, a resistance of 0.15 Ωcm2 to the total internal resistance of the fuel cell. For
GDC at 500°C a maximum allowable electrolyte thickness of 15 μm is reported.
Based on the conductivity of MCP-SDC at 600°C the maximum allowable electrolyte
thickness was 13.5 μm.
6.2 Ni-SDC Anode
In this section the electrical properties of the oxide precursors and their conversion
into a cermet by exposure to a reducing atmosphere are discussed.
92
6.2.1 Cermet Precursor
The conductivity of sintered NiO-SDC composites as fabricated (Section 5.2.2) were
measured over a temperature range from 400-800°C as a reference point for their
subsequent reduction. As the NiO content in the composite increased, the
conductivity increased dramatically and a percolation threshold was passed at around
40 vol%NiO as shown in Figure 6.2. Since the densities of NiO and SDC are similar
it has been assumed that wt%NiO = vol%NiO. When corrected for sample porosity
the percolation threshold is closer to 30 vol%NiO which is in agreement with the Ni-
YSZ from the literature [26].
0.0001
0.001
0.01
0.1
1
0 20 40 60 80 100
400oC [S/cm]500oC [S/cm]600oC [S/cm]700oC [S/cm]800oC [S/cm]
Con
duct
ivity
[S/c
m]
wt%NiO
Figure 6.2 Cermet precursor conductivity as a function of NiO-SDC composition showing percolation
of acceptor-doped NiO.
93
Pure NiO is a p-type semiconductor with cationic vacancies and electron holes as
primary defects. Its conductivity can be enhanced substantially from a baseline of
0.01 S/cm at 600°C by doping with acceptors such as Na and particularly Li, for
example, 2.4 mol% Li-doped NiO has a conductivity of around 200 S/cm [107]. The
obvious acceptor dopant is Na, originating from NaCl or NaOH present during milling.
The typical Na residual level for attrition milled 55 wt%NiO-SDC was 370 ppm as
determined by ICP. For SPEX milled composites the Na contamination is higher, but
still less than 1000 ppm. Pure Ni(OH)2 milled with the same portions of NaOH and
NaCl had a Na concentration of 1920 ppm after washing and drying, corresponding to
0.6 mol% Na-NiO. Since the conductivity of 0.5 mol%Na-NiO is around 1 S/cm the
conductivites obtained for the cermets with high NiO compositions are thus
rationalised.
The temperature dependence of the conductivity, plotted in Figure 6.3, shows the
conduction mechanism changing from ionic to electronic as observed in the NiO-YSZ
system [108]. Attempting to extract activation energies for ionic conduction in SDC
and the electron-hole conduction in NiO is complicated by the fact that ionic and
electron-hole conductivities have diverging temperature dependences. Ionic
conduction is best described by a diffusion modified Arrhenius equation (Equation 9)
and has a strong temperature dependence, σiT = (A/T).exp(-Ea1/kT). Conversely,
electron-hole conduction is described by the Arrhenius equation, σe = B.exp(-Eae/kT),
and has a weaker temperature dependence. The magnitude of the conductivity for
doped NiO and SDC were comparable over this temperature range, so it was not
possible to separate the ionic and electronic contributions to the total conductivity. An
apparent activation energy can be calculated from the measured slopes and it has a
94
value of 0.8eV from 0 wt% to 20 wt%Ni, decreasing with increasing NiO content to a
value of 0.16eV above 50wt% Ni. This corresponds to the change in majority charge
conduction mechanism assocated with the percolation of the NiO phase.
0.0001
0.001
0.01
0.1
1
0.9 1 1.1 1.2 1.3 1.4 1.5
0wt%NiO20wt%NiO35wt%NiO55wt%NiO65wt%NiO50wt%NiO80wt%NiO
Con
duct
ivity
[S/c
m]
1000/T[K]
Figure 6.3 Conductivity vs reciprocal temperature for NiO-SDC cermet precursors
In the single chamber SOFC atmosphere, NiO can be kinetically stable and may
remain at the core of a particle after the exterior has been reduced. The contribution
that electron-hole conductivity makes to the total conductivity is unlikely to have a
negative impact to the performance of the anode.
95
6.2.2 Cermet Formation
Exposure of NiO-SDC ceramics to an atmosphere of 10%H2-Ar at 600°C gave a rapid
increase in the conductivity. Figure 6.4 shows the reduction is largely complete after
10 minutes annealing. The transient increase in conductivity for 35 wt%NiO-SDC is
because the NiO particles are initially reduced on their surfaces, and a slight increase
in conductivity is seen before further volume reduction causes the Ni surfaces to
separate [26].
0.001
0.01
0.1
1
10
100
1000
100000 5 10 15 20 25 30
Annealing Time [mins]
Con
duct
ivity
[S/c
m]
65wt%NiO
50wt%NiO
35wt%NiO
55wt%NiO
Figure 6.4 Cermet formation via reduction of NiO-SDC composites in 10%H2/Ar at 600°C
Reduction was accompanied by a colour change in the composite from light grey-
green to a dark grey and the introduction of porosity to the microstructure. Figure 6.5
shows electron micrographs of an incompletely reduced cermet from 65 wt%NiO-
SDC. A boundary can be seen in the lower-right section of both images, appearing
96
denser and more continuous in the secondary electron image and lighter in the
backscattered image. This suggests that the lower-right region was unreduced.
Figure 6.5 Electron micrograph of 65wt%NiO-SDC ceramic showing the interface between Ni and
NiO(left), backscattered electron image of the same region(right)
Individual Ni/NiO and SDC grains were more readily visible on the backscattered
micrograph (Figure 6.5(right)). The light SDC regions are uniformly dispersed with
grey Ni/NiO and black pores. The grain sizes were 158 nm for SDC and 196 nm for
Ni, somewhat larger than the obtained XRD derived crystallite sizes of 54 nm and 69
nm.
As the NiO content in the ceramic increases, the conductivity changes by several
orders of magnitude. In Figure 6.6 the conductivity passes the percolation threshold at
around 30vol%Ni and plateaus at around 50vol%Ni for 1068 S/cm. This value is
13.4% larger than the 942 S/cm quoted for Ni-GDC at 682°C [109] and comparable to
the value for Ni-YSZ of 1311 S/cm at 677°C [26].
97
0.001
0.01
0.1
1
10
100
1000
100000 20 40 60 80
Volume fraction of Ni in Ni-SDC Cermet [%]
Con
duct
ivity
at 6
00°C
[S/c
m]
100
Figure 6.6 Conductivity of Ni-SDC cermets at 600°C as a function of Ni content
The conductivity as a function of temperature for the anode substrate precursor, 55
wt%NiO-SDC and its corresponding 50 vol%Ni-SDC cermet are shown in Figure 6.7.
The activation energy for the precursor was 0.77 eV, close to that of pure SDC; hence
the prevailing conduction mechanism is ionic. The metallic conduction of the cermet
was characterised by an activation energy of 4.8kJ/mol, slightly lower than the
literature value of 5.38 kJ/mol [26].
98
0.1
1
10
100
100
1000
104
0.0009 0.001 0.0011 0.0012 0.0013 0.0014 0.0015
sT [K
S/c
m] 5
5wt%
NiO
/SD
C in
air
s [S/cm
] 50vol%N
i/SD
C as reduced in A
r
σT [K
.S/c
m]
σ [S/cm
]
1/T [1/K]
Figure 6.7 Temperature dependence of conductivity for 55 wt%NiO-SDC and 50 vol%Ni-SDC
Figure 6.8 shows the ultra-fine, porous microstructure obtained after reduction. The
grain size was similar to that found for previous cermets. Whether a density of 75%
will allow facile diffusion of methane, water and carbon dioxide molecules necessary
for efficient fuel cell operation will be assessed during fuel cell testing, later in this
thesis.
99
Figure 6.8 Scanning electron micrograph of 50vol%Ni-SDC cermets
6.2.2.1 Pore Forming Additives
As discussed in section 2.1.1.1, it is essential for the anode to have porosity to allow
facile gas diffusion so that concentration polarisation is minimised. The degree of
porosity must however be balanced with the mechanical strength of the anode. Corbin
and Apte suggest that graphite particles larger that 50 μm leave pores without
contributing to the overall shrinkage of the specimen [22]. Commercially supplied
graphite (ThermoPURE, Superior Graphite Co.) with a mean particle size of 68 μm
was chosen to form porosity within 55 wt%NiO-SDC.
The composites were prepared from milled, heat-treated, washed, magnetically
separated, ethanol washed and centrifuged NiO-SDC cake and combined with 3.5
wt% DisperByk-111. Graphite was added to achieve 9.1, 13.1 and 16.8 vol%
specimens and additional ethanol was added with 3 mm YSZ milling media, the
100
mixtures were homogenized by shaking polypropylene vials in a SPEX mill for 30
mins. The addition of binders and plasticizers such as those employed for tape casting
(Section 5.1.1) was trialled but did not improve the ease in which the material could
be handled. After drying the powder overnight at 160°C it was crushed in a mortar
and pestle before compaction. Additional graphite was used as a die lubricant/pin
release agent because the green compacts were weakened significantly by the
inclusion of graphite and had a tendency to stick to the pin, leading to fracture during
unloading. This was the primary reason why compositions above 17 vol% graphite
were not studied. A fracture surface of a specimen that failed in such a manner was
examined and Figure 6.9 shows the larger graphite particle completely covered by
~100 nm clusters of 55 wt%NiO-SDC.
Figure 6.9 Electron micrograph from a fractured pellet of 55 wt%NiO-SDC with 20 vol% graphite
Dilatometry measurements plotted in Figure 6.10 show that the inclusion of graphite
displaced the temperature of the maximum sintering rate from 1200°C to 1350°C.
Thermogravimetry showed that the exothermic combustion of graphite occurring at
101
777°C, coinciding with the onset of the first stage of shrinkage. The decomposition of
graphite occurred without introducing significant flaws to the microstructure, the
specimens were reasonably dense but qualitatively more fragile than those without
graphite added.
-30
-25
-20
-15
-10
-5
0
-0.12
-0.1
-0.08
-0.06
-0.04
-0.02
0
500 700 900 1100 1300 1500
Shr
inka
ge [%
DL/
Lo]
Der
ivat
ive
(T) S
hrin
kage
[d/d
T(D
L/Lo
)]
Temp [deg C]
Figure 6.10 Dilatometry curve for 55 wt%NiO-SDC with 10 vol% graphite included.
The green pellets were sintered at 1350°C for 4 hours, then reduced in flowing H2-Ar
at 600°C for 1 hr and cooled to room temperature under Ar. Large voids were visible
to the naked eye. Electron microscopy examination of polished sections (Figure 6.11)
shows the elongated cavities left by the graphite and the access that these give to
internal anode regions.
102
Figure 6.11 Electron micrographs of cermet formed from 55 wt%NiO-SDC with 15 vol% graphite;
Secondary image(right), backscattered image(left)
The densities of the reduced specimens were 55%, regardless of the original graphite
content, in comparison to 75% density for the cermet without graphite. The areal
porosity was calculated by image analysis of polished specimens, derived from images
typical of those shown in Figure 6.12 for 9.1 vol%(left) and 16.8 vol%(right) graphite.
The graphite particles appear to agglomerate into planar clusters around 16.8 vol%
resulting in several hundred micrometre pores in Figure 6.12(right).
Figure 6.12 Scanning electron micrographs of reduced and polished 50vol%Ni-SDC from
9.1vol%(left) and 16.8vol%(right) included graphite
The pressure dependence of the permeability of the reduced cermets is shown in
Figure 6.13, with the exception of 50 vol%Ni-SDC which was found to be
103
impermeable. At 9.1 vol% graphite the cermet has the same permeability as 13.1
vol% graphite and the doubling that occurs at 16.8 vol% graphite was caused by the
formation of a percolated network.
0.0E+00
2.0E-10
4.0E-10
6.0E-10
8.0E-10
1.0E-09
1.2E-09
1.4E-09
1.6E-09
1.8E-09
0.0E+00 5.0E+04 1.0E+05 1.5E+05 2.0E+05 2.5E+05 3.0E+05 3.5E+05
p[Pa]
K[m
ol m
/m2
Pa
s]
10vol% graphite
15vol% graphite
20vol% graphite
Figure 6.13 Permeability vs pressure difference for reduced cermets with added graphite
Particles with high aspect ratios such as graphite are able to percolate at lower
thresholds due to the fact that the large excluded volume associated with flake-like
particles gives an increased probability of inter-particle contact. The porosity induced
by the inclusion of 16.8 vol% graphite gave a permeability around 700 times larger
than the 2.5*10-12 mol.m/(m2Pas) reported for a porous aluminosilicate [110]. When
calculated according to the standard AS-1774.7 the permeability of ca. 10-11 m2 was
100 times lower than the value reported for a coarse Ni-YSZ cermet [111]. The
conductivity of the 16.8 vol% graphite cermet was measured at 600°C and it had
decreased by 60% to 475 S/cm. For a 500 μm thick anode this yields a resistance of
104
105 μΩ.cm2 which is negligible when compared to the 150 mΩ.cm2 resistance
allowed for the electrolyte.
The triple phase boundary (TPB) length and the contiguity of anode, electrolyte and
pore phases are important parameters for the characterisation of a fuel cell anode, yet
there was no consistent approach to its determination found in the literature. Lee et. al.
[112] apply a 3D extension of 2D quantitative microscopy relations in Equation 17
and Equation 18 to determine the contiguity (Cαα) of three phases α, β and χ (i.e. in
Ni, SDC and Pore) with one another.
αγαβαα
αααα
LNLNLNLN
C++
=2
2
Equation 17
βγββαγαβαα
αβαβ LNLNLNLNLN
LNC
++++=
22
2
Equation 18
In these equations, NLαβ is the number of contact points between the respective phases
per unit length. A backscattered micrograph, such as the one shown in Figure 6.14,
was used for the determination of this microstructural data. Fortunately the contrast
between Ni and SDC and the pores allows these phases to be easily distinguished, this
is not the case for Ni and YSZ.
105
Figure 6.14 Backscattered electron micrograph of 50vol%Ni-SDC with 13.2vol% graphite added.
White regions=SDC, Grey=Ni, black=pores.
There were no differences in the bulk microstructure as the graphite content increased.
The only change was the increased areal pore fraction, the ultra-fine composite of Ni,
SDC and pores is seemingly unaffected by the inclusion of graphite, which affects
only the bulk porosity. The contiguities and interfacial areas were calculated and
compared to the results of Lee et. al. for 50 vol%Ni-YSZ in Table 6.2. The interfacial
area was defined as 2NL.
106
Table 6.2 Contiguity data for Ni-SDC(E)-Pore cermet
Parameter
(E=Electrolyte)
50vol%
Ni-YSZ [112]
50vol% Ni-SDC
Ni grain size [μm] 4.6 0.20
E grain size [μm] 2.0 0.16
Ni-Ni contiguity 0.25 0.25
Ni-E contiguity 0.175 0.52
E-E contiguity 0.17 0.16
Ni-pore contiguity 0.08 0.18
E-pore contiguity 0.06 0.22
Ni-Ni grain boundary area
[μm2/μm3]
0.25 0.86
E-E grain boundary area
[μm2/μm3]
0.11 0.62
Ni-pore interfacial area
[μm2/μm3]
0.26 0.58
Ni-E interfacial area
[μm2/μm3]
0.3 4.5
E-pore interfacial area
[μm2/μm3]
0.15 0.81
The contiguity between Ni-Ni and E-E are identical for both cermets. Because the
relative particle sizes of Ni and electrolyte within each data set, the formation of a
percolating network of Ni within the electrolyte phase dictates that the homogeneous
contiguity will be similar for both cermets. A significant difference can be seen for
the Ni-E and Ni-pore contiguity, and on average the Ni-SDC cermet is 200% larger
than Ni-YSZ. The 1410% increase in the Ni-E interfacial area is also associated with
a finer grain structure, but unfortunately this is accompanied by a mere doubling in the
Ni-pore interfacial area. It is visibly evident that the interfacial area between Ni and
pore is much larger and clearly this quantitative approach can not account for the
107
observed bimodal porosity distribution. Furthermore, the electrochemical reaction
zone is reported to extend, at most, 20 μm from the electrolyte surface into the anode
so the significance of these equations is questionable and another approach was
required.
If the dense electrolyte layer exists as if superimposed over the microstructures and
pores, shown in Figure 6.12 and Figure 6.14, then for a given areal fraction of porosity,
the cermet’s TPB length can be determined by the TPB length per pore. Since all
cermets were exposed to the same temperatures during sintering and reduction, there
were no significant differences in the grain sizes of Ni or SDC and the only factor that
will change the TPB between the three porous cermets fabricated will be their pore
fraction as determined by image analysis. The pore selected for analysis is shown in
Figure 6.15, which has been enlarged from the upper-left region of Figure 6.14.
Figure 6.15 Backscattered electron micrograph of 50vol%Ni-SDC showing the pore under
consideration
From the mean size of the graphite particles, 68 μm at a 10:1 aspect ratio and
cylindrical geometry, the pores can be modelled as 66*6.8 μm rectangles. If the
108
electrochemical reaction zone is at most 20 μm in depth, then the electrode area
exposed within such an idealised pore is 300 μm2. Microstructural analysis gave an
interfacial area per unit volume, equivalent to the interfacial boundary length per unit
area of 4.5 μm/μm2, so the Ni-SDC boundary length within each pore was 13.5μm.
The theoretical number of ideal rectangular pores present on the cermet-electrolyte
interface can be determined from the pore area of 462 μm2 and the measured areal
pore fraction for cermets originating from the 9.1, 13.1 and 16.2 vol% graphite. This
process is described graphically in Figure 6.16.
Figure 6.16 Microstructure model used to determine TPB length
These calculations yield TPB lengths of 95, 166 and 211 m/cm2 respectively. These
values are comparable to the 45-61 m/cm2 reported by Bouwmeester [113] for a
sputtered Ni anode on YSZ electrolyte. For the highest porosity obtainable in the
cermets fabricated here, synthesized from 55wt%NiO-SDC with 16.2 vol% of
109
graphite, a 250% increase over the literature value was obtained by virtue of an ultra-
fine microstructure.
Anode supported fuel cells must traditionally sacrifice a considerable amount in triple-
phase boundary length to allow for sufficient structural integrity. Whilst it is essential
for open porosity to exist, it can be difficult for this porosity to generate an
appreciable triple phase boundary length. There are very few reports of actual or
estimated triple phase boundary lengths in the literature, especially for anode-
supported fuel cells. In this instance the use of a nanoscale NiO-SDC composite
powder gave an ultra-fine cermet microstructure. The cermets inherent microstructure
was not sufficient to allow a high triple phase boundary length for it had negligble gas
permeability. Increasing the porosity with fugitive graphite particles allowed gas to
permeate throughout the structure and gave a TPB length that was several times larger
than reported in the literature.
110
7 Fabrication and Performance of a Single Chamber
SOFC
The properties of the materials described in previous chapters provided loose limits on
the processing conditions necessary to fabricate the fuel cell. Firstly a ~500 μm thick
anode consisting of 55 wt%NiO-SDC must be compacted. The exact thickness is
unimportant provided the specimen had sufficient strength to be manipulated during
subsequent processing steps.
Secondly, based on its measured conductivity reported in Section 6.1, an electrolyte
layer of SDC with a maximum sintered thickness of 13 μm was required. Either
airbrushing or spin coating were suitable methods for depositing such a layer. In its
green, unsintered form the layer must be continuous and as defect free as possible.
This has implications for the surface of the anode; it must be clean and as flat as
possible – preferably as pressed from the die. The sintering required to achieve
appreciable density in the electrolyte must be balanced with the coarsening of the
anode microstructure, hence firing of the anode before electrolyte deposition as
described by some research is not to be recommended.
Thirdly, the 10 wt%SDC-SCC cathode was to be deposited on to the electrolyte with a
thickness around 20 μm. The application methods investigated were painting, spin-
coating and airbrushing. To avoid short circuiting the cell the cathode must not touch
the anode and this was achieved by application over a circular 8 mm mask. The
evolution of a suitable processing regime to optimise the fuel cell performance,
operating in single chamber mode, is discussed in the following sections.
111
7.1 Electrolyte Fabrication by Spin Coating
Spin-coating is a well established technique in the electronics industry, allowing high
reproducibility for a given speed and coating composition. It has been shown to work
for nanoparticle suspensions of SnO2 gas sensors [114, 115]and CeO2 [73]. When
coating SDC onto the anode precursor substrate a practical issue arises. Excess slurry
is spun off the surface and some becomes entrained in the vacuum that attaches the
specimen to the rotor and results in electrolyte material sticking to the lower surface.
The extra handling required to clean the opposing face risks damaging the unfired
layer. Measures such as sticking the anode to a glass slide with tape prevent contact
with the slurry means that the specimen must be detached, again there is a risk of
damage.
Suspensions of 3-5 wt% MCP-SDC in ethanol, from centrifuged residues, dispersed
with 3.5 wt% Disperbyk-111 were spun onto as-pressed 55 wt%NiO-SDC compacts.
The nature of the spin-coating process tends to leave a circle of increased thickness
around the rim of the specimen, which upon drying will crack and render the region
unusable. Similarly, when more viscous suspensions are applied, such as those
containing binders or plasticizers, the central region retains more slurry and this will
also crack. Sintering this bilayer at 1200°C for 1 hour did not achieve appreciable
densification for the electrolyte, it appeared powdery and was readily scraped or
blown off, evidently the result of a lower green density. Further sintering was
therefore necessary, but the drastic difference in microstructure when sintered for 1 hr
at 1300°C vs 1400°C is shown in Figure 7.1.
112
Figure 7.1 55wt%NiO-SDC//SDC anode substrate spin-coated, sintered at 1300°C (left) vs 1400°C
(right). Arrows denote the anode//electrolyte interface
More problematic than the increase in electrolyte grain size from 250 nm to 613 nm, is
the increased density of the anode layer and the absence of open porosity. Figure 7.1
(left) has ca. 250 nm pores visible throughout the anode section whereas Figure 7.1
(right) has no visible pores as such. Gas permeation measurements performed on a
reduced anode pellet revealed that it was impermeable; what is needed therefore is the
porosity induced as described previously in Section 6.2.2.1. As a compromise and for
all following fuel cell anode/electrolyte bilayers, sintering was performed at 1350°C
for 4 hours unless otherwise specified. Under these conditions, spin-coating of three
SDC layers, with oven drying between coats gave a sintered electrolyte of 3-5 μm
thickness.
Examination of the SDC surface, as coated and after firing, highlighted the
weaknesses of this technique. Surface imperfections on the substrate imprinted from
the pin during compaction caused the green film to crack during air-drying, as shown
in Figure 7.2. In this micrograph the suspension consisted of AMR-SDC (Section
113
4.1.2.1.1) and the agglomerates present on the surface may also have acted as crack
nucleation sites.
Figure 7.2 Micrograph of an oven-dried spin coated layer of AMR-SDC showing agglomerates and
surface cracking
The presence of atmospheric moisture appeared to accelerate this process as the
unfired layer would visibly rehydrate once removed from the oven. The cracking was
present to varying degrees in all specimens but by applying more than 2 coats it was
believed that complete coverage of the anode by the electrolyte would be achieved.
The cross-section of a cell fabricated using spin coating (Figure 7.3) shows that fuel
cells with thin electrolytes were obtained using these methods. However, since such
cells in operation yielded no significant OCV’s, it was assumed that the cathode was
in electrical contact with the anode via electrolyte imperfections. For the deposition
of such thin films on an uneven substrate airbrushing will be shown to be a superior
method.
114
Figure 7.3 Fuel cell microstructure showing spin-coated SDC electrolyte and SDC-SCC cathode
7.1.1 In-Situ Reduction of Cermet Precursors
Cermet formation was monitored by measuring the conductivity as the temperature
within the atmosphere (6% CH4-6% O2-88% Ar, “Fuel Cell Gas” or FCG) of the
reaction vessel increases. The reduction was complete before the specimen reached
800°C, as shown by the steep increase in conductivity in Figure 7.4. The final
conductivity at 600°C was 460 S/cm and this did not change when the atmosphere
was changed from FCG, to Ar or 10%H2-Ar. Complete reduction of NiO had
occurred and this was verified by its absence in the XRD pattern. In order to initialise
fuel cell specimens for operation, exposure to FCG at 800°C was necessary.
115
Figure 7.4 Conductivity and temperature vs annealing time for the reduction of 55wt%NiO-SDC
ceramic under 200mL/min FCG
7.1.2 Cell construction using Pt mesh current collector
The choice of a suitable current collecting material for a single chamber SOFC was
not obvious from the literature. Hibino [48] suggests Pt for the anode and Au for the
cathode and presumably Ag will perform equally well. Some researchers use Pt
exclusively, which is questionable since Pt is a combustion catalyst, hence a reaction
with the methane would be expected. Initial experiments were performed using Pt
mesh and Ag wires, later Ag wires were used exclusively. Finally a Ag-SCC porous
paint was developed to address cell cathodic performance issues.
An anode/electrolyte bilayer was fabricated; the anode was 55 wt%NiO-SDC with 20
vol%Ni that had been pre-fired at 1000°C and polished. After sintering at 1350°C for
116
4 hours a slurry of 10%SDC-SCC (MicroCoating) in ethyl carbitol was painted over
the electrolyte as evenly as possible. As this paint dried, the cracks were touched up
and any spillages over the edge onto the anode were scraped off. Firing for 2 hrs at
900°C formed a comparatively dense layer and no regions of electrolyte were visible
through any of the cracks. The fuel cell was clamped between two 8 mm circles of Pt-
mesh, insulated from the clamp by 40% porous Al2O3 bricks and was attached to leads
inside the flow cell as described in Section 3.4.7.2. After annealing at 800°C for 30
mins under 260 mL/min of FCG, the OCV was notably higher than its initial value,
and electrochemical activity was confirmed by the change in this voltage with
temperature as shown in Figure 7.5.
050
100150200250
300350400450500
400 500 600 700 800T[oC]
OC
V[m
V]
Figure 7.5 OCV as a function of furnace temperature for the fabricated fuel cell
The power characteristic for selected temperatures is shown in Figure 7.6 and the
maximum power output was 161 μA/cm2 at 625°C. A low value was anticipated due
to the dilute gas mixture employed here, in contrast to that of Hibino [10] who
obtained 150 mW/cm2 at 700°C using 52 mL/min of 17% CH4 and O2. Aside from
117
the difference in gas composition, the power output was so poor that no further
analysis was performed.
Figure 7.6 Power vs current drawn for fuel cell 7.1.2 at selected temperatures between 550 and 750°C
in 260 mL/min FCG
The three orders of magnitude difference in power output and an OCV two thirds of
the literature value suggest that there is much to be optimised in this system, starting
with the current collector material. A fine black deposit, presumably catalytically
deposited carbon, was observed on the Pt mesh and wires inside the cell after testing.
It would be preferable to have a less catalytically active material on the anode side
since the methane reacting on the current collector will reduce the amount that is able
to penetrate the anode and react at the TPB. For the cathode, Pt has a high activity for
oxygen reduction but poor oxygen diffusivity. Silver electrodes appear to be the most
sensible choice and are suitable given the low operating temperatures. Silver has a
118
high oxygen diffusivity, high conductivity and low catalytic activity and is therefore a
more sensible, and cheaper, current collector than gold which is impermeable to
oxygen [116].
7.1.3 Cell construction using Ag wire current collector
A cell was fabricated by the same method as described in Section 7.1.2, where SDC
deposited by spin coating onto pellets of porous NiO-SDC, fired previously at 1000°C.
In the final assembly loops of ø0.5 mm Ag wire were placed between the electrode
surface and the porous alumina separators. This gave a much improved power output
of 483 μW/cm2 at 650°C shown in Figure 7.7.
0
0.1
0.2
0.3
0.4
0.5
0.6
0 1000 2000 3000 4000 5000I [uA]
P[m
W/c
m2]
600650700750800
Figure 7.7 Power vs current drawn for fuel cell 7.1.3 over selected temperatures between 550 and
800°C in 260 mL/min FCG
119
The dependence of the terminal voltage and maximum power on operating
temperature in Figure 7.8 shows a significant drop in both OCV and maximum power
at 750°C, and a recovery by 800°C.
160
180
200
220
240
260
280
300
320
0
0.1
0.2
0.3
0.4
0.5
550 600 650 700 750 800 850
OC
V[m
V]
Pm
ax[m
W/c
m2]
Temp [deg.C]
Figure 7.8 OCV and Maximum power output for fuel cell 7.1.3 as a function of temperature
Galvanic current interruption (CGI) was used to separate the internal ohmic losses
from polarisation losses. Because a considerably dilute gas mixture was used, the
reduced rate of reaction will generate smaller currents, hence the resistances obtained
can only be compared internally to other cells operating from FCG. It is not possible
to simply correct for the difference in gas concentration because the differing reaction
rates on either surface are also likely to be functions of concentration. There was
sufficient stability to separate the polarisation overpotential from the cell potential for
a range of current densities at 600 and 650°C, as shown in Figure 7.9. The
120
measurements were unstable in the upper temperature ranges for this experiment so
the shape of the curves above 700°C in Figure 7.8 remain anomalous.
Figure 7.9 Polarisation Loss as a function of current drawn from the fuel cell
The decrease in polarisation resistance from 778 to 271 Ω.cm2 with increasing
temperature is associated with the 556% increase in cell power output, despite a 29%
loss in open cell voltage. The internal ohmic resistance was constant at 510±45 Ωcm2.
When broken down by the components thicknesses and known conductivities, shown
in Table 7.1, the internal resistance was 490 Ωcm2 larger than calculations suggest.
Post test inspection of the cell revealed that the cathode had changed colour from
black to orange/brown over most of its surface, so its real contribution to the cell’s
resistance was probably much larger than predicted.
121
Table 7.1 Summary of cell component resistances based on GCI measurements
Component Thickness,
μm
Resistivity at
600°C, S cm-1
Resistance in
cell, Ω.cm2
Anode 675 460 0.15
Electrolyte 1.8 0.009 20
Cathode 3.1 400 7.75E-4
TOTAL - - 20.15
XRD analysis showed that the cubic (SmSr)0.5CoO3 phase was no longer detectable
amongst an ambiguous mixture of peaks and only CoO (and SDC) could be positively
identified. Although other researchers employing SCC in single-chamber fuel cells
have not reported the same effect, the fine microstructure of the SCC composite may
have lead to its rapid internal decomposition during fuel cell operation. This was
rectified by the use of coarse SCC in subsequent results and discussions. Furthermore,
there is evidence in the literature that the thickness of LSM cathode in YSZ fuel cells
should be greater than 25 μm, as excessive cathodic polarisation was reported [117],
presumably the SCC cathode is too also thin in this instance.
7.1.4 In-Situ Reduction of Cathode
In order to assess the behaviour of MicroCoatings’ SDC-SCC, a sintered specimen
was sectioned and prepared as 55 wt%NiO-SDC as described in 7.1.1. The electrical
conductivity of the ceramic was measured as a function of temperature and as shown
in Figure 7.10, it appears to be metallic in nature with conductivities between 300-500
S/cm decreasing with increasing temperature. The conductivity is significantly higher
than reported for pure (SmSr)0.5CoO3 by Kang [101], which was at most 1.0 S/cm
from 773-1273°C.
122
100
150
200
250
300
350
400
450
500
550
300 400 500 600 700 800 900
Temperature [deg. C]
Tot
al C
ondu
ctiv
ity [S
/cm
]
Figure 7.10 Conductivity vs temperature for sintered 10 wt%SDC-SCC in flowing air
The conductivity was monitored as the conditions required to reduce NiO in the FCG
atmosphere were reached, that is as the atmosphere in 6%CH4-O2 increased in
temperature from 600-800°C. During this process the conductivity decreased by four
orders of magnitude as the temperature reached 800°C, shown in Figure 7.11.
123
600
650
700
750
800
850
0.01
0.1
1
10
100
1000
0 50 100 150 200
Tem
p [o
C]
Con
duct
'y [S
/cm
]
Time [mins]
FCG -> air
air -> Ar
Ar -> H2/Ar
Figure 7.11 Temperature and conductivity vs. annealing time for 10%SDC-SCC ceramic. Gas
composition is FCG(6%CH4, O2) and is changed as marked.
When the temperature returned to the 600°C baseline for fuel cell operation, the
conductivity increased to 16 S/cm. This recovery is a significant decrease from the
starting value of 400 S/cm and it did not improve upon exposure to air. A similar
decrease in conductivity as the atmosphere changed to hydrogen showed that this
material decomposes in the atmosphere that was intended for single-chamber fuel cell
operation. Microstructural examination did not reveal any useful findings.
To eliminate the possibility that the SDC-SCC, as received from MicroCoating, was
too fine it was annealed at 1200°C before application. This was compared with the
material prepared by mixing 10 wt% of MCP-SDC with SCC as synthesized by the
124
literature method [101]. An elemental assay (ICP) of the MicroCoating material
(Table 7.2) shows that the molar ratio of Sm to Sr was correct at 1.042. However both
the Sr:Co and Sm:Co ratios are 2.3 and 2.23 respectively, significantly larger than
dictated by stoichiometry.
Table 7.2 ICP Elemental assay for 10%Sm0.2Ce0.8O1.9-x-(SmSr)0.5CoO3
Element Concentration,
wt%
Ce 5.31
Co 22.3
Sm 27
Sr 14.3
The influence that this excess cobalt oxide has on the stability of the crystal structure
under reducing conditions is unknown. Given the observation of CoO and Co3O4
phases in the poorly resolved XRD pattern of decomposed cathodes, it became
apparent that traditionally fabricated SCC would make a more consistent cathode.
7.2 Electrolyte Fabrication by Airbrush
Airbrushing or spray-painting is a ubiquitous technique ideally suited for the
application of nanoparticles. Layers were fabricated by spraying 20 mL of 0.3 wt%
suspensions of SDC in ethanol onto green NiO-SDC pellets held at 80-100°C on a
hotplate. After sintering at 1350°C for 4 hours the cross section and surface
microstructures shown in Figure 7.12 give an indication of the uniformity and absence
of cracks.
125
Figure 7.12 Electron micrographs of airbrushed SDC on NiO-SDC substrates, sintered at 1350°C for
4hrs. The cross section shows the slightly porous SDC layer on the base(left); a typical SDC surface
region(right)
The 5 μm electrolyte was highly uniform along this cross-section, the only
imperfections being micron sized spherical regions, visible in Figure 7.12 (right)
which were not believed to impact negatively on the electrolyte’s mechanical
properties.
7.2.1 Cell construction using Ag wire current collector
An electrolyte/anode precursor bilayer was prepared using the airbrush-hotplate
techniques above, sintering for 4 hours at 1350°C. A suspension of calcined SCC
(MicroCoating) was mixed with 10 wt% SDC in ethanol and diluted to 3 wt% solids.
This was sprayed onto the electrolyte over an ø8mm mask and fired at 900°C for 2
hours. The cell, attached with Ag wires, produced a power characteristic as shown in
Figure 7.13 and had a maximum power output of 2.02 mW/cm2 at 650°C, a fourfold
increase from that described previously. This was ascribed to the increased size of the
cathode precursor starting material forming a thicker cathode layer.
126
0
0.5
1
1.5
2
2.5
0 1000 2000 3000 4000 5000 6000 7000I[uA]
P[m
W/c
m2]
600 650
850 800
450 500
550
Figure 7.13 Power vs current drawn for the fuel cell at selected temperatures in 30 mL/min FCG
The OCV and maximum power output again did not share a common temperature and,
as shown in Figure 7.14, the maximum power was obtained at 650°C which had the
lowest OCV. This shows that the electrode reaction kinetics were more important
than cell thermodynamics in determining the power output of a single chamber SOFC.
127
400
450
500
550
600
0
0.5
1
1.5
2
2.5
400 450 500 550 600 650 700
OC
V [m
V]
Pm
ax [m
W/c
m2]
T [deg.C]
Figure 7.14 OCV and Maximum power output for fuel cell 7.2.2 as a function of temperature
Galvanic current interruption measurements allowed the separation of internal ohmic
(Rint) and polarisation (Rpol) losses and hence resistances. As discussed previously,
the magnitude of these resistances is high because of the dilute fuel mixture. This
issue is overcome in the following discussion by using undiluted gas. The
temperature dependence of these resistances can be modelled with the Arrhenius
equation, and Figure 7.15 shows that the experimental data was consistent with this
approach. The activation energies for the ohmic and interfacial resistances are 32 and
92 J/mol respectively, far lower than the 70-170 kJ/mol reported in the literature for a
Ni-YSZ//YSZ fuel cell [118].
128
1
10
100
1000
1.05 1.1 1.15 1.2 1.25 1.3 1.35 1.4
Rint [Ohm.cm2]Rpol [Ohm.cm2]
Rin
t [O
hm.c
m2]
1000/T [1/K]
Figure 7.15 Temperature dependence of ohmic (Rint) and interfacial (Rpol) resistances for fuel cell
from Section 7.2.2
Decomposition of the cathode was suspected due to a brown discolouration which
appeared on the black cathode material after testing. Analysis of the XRD pattern
showed that the original SCC perovskite had persisted, but there was evidence for
CoO and a complex mixed oxide Sr2Co2O5.82 phase, presumably a decomposition
product. Since the (SmSr)0.5CoO3 supplied by MicroCoating was known to have an
excess of Co, residual CoO and other decomposition products suggests that the
material was not suitable for single-chamber solid oxide fuel cell cathodes.
Furthermore, any such cathode in a mixed-gas environment must have strict
stoichiometric controls to obtain large currents.
129
There was no visual change to the cathode, or other defects that would suggest its
failure from post-testing microstructural examination. The 34 μm cathode, 4.7 μm
electrolyte and 34.1 μm cathode are shown in a fracture cross-section in Figure 7.16
supported on the 873 μm anode.
Figure 7.16 Cross section of fuel cell after testing with the anode support on the left side
7.2.2 Cell construction using Ag wire current collector and Ag
composite paste on Cathode
To further improve electrical contact between the cathode and the Ag wire current
collector, a porous Ag paste was formulated. Compliance with the cathode was
achieved by incorporating 30 vol% of 10 wt%SDC-SCC (MCT treated at 1200°C)
relative to Ag (ProSciTech). For gas diffusion 20 vol% final porosity was achieved
by the inclusion of graphite. By mass the paste was 5.92 wt% graphite, 77 wt% Ag,
130
16 wt% SCC-SDC and it was carried in minimum amount of dowanol (OnGlaze).
When dried at 160°C it exhibited electrical continuity and as shown in Figure 7.17
(left), appeared to have considerable porosity. To achieve graphite decomposition it
was fired at 800°C, the fired pieces were found to have maintained their shape,
conductivity and were highly porous as shown in Figure 7.17 (right). Although
elemental analyses were not performed, it was assumed that the continuous, low-
contrast phase was Ag and the rougher, faceted agglomerates were the SDC-SCC
cathode material. The testing of the material is described in the following section.
Figure 7.17 Dried 160°C (left) and fired 800°C (right) Ag-based cathode current collector paste
7.3 Single Chamber Fuel Cell Operation in Undiluted Gas
Mixtures
The testing of fuel cell specimens in dilute gas mixtures has confirmed three facts.
Firstly the anode functions as an electrical conductor and allows the facile diffusion of
gases due to a combination of macro-pores and nanoscale porosity. Secondly, a thin
layer of electrolyte was successfully sintered to this anode and the anode can be
reduced whilst maintaining the electrolyte’s integrity and ultra-fine grain structure.
Thirdly, a cathode layer was sintered onto the electrolyte and the fuel cell displayed
131
evidence of catalytic activity by the formation of an open cell voltage and the supply
of a current over varying loads. The gas atmospheres to this point have been as
concentrated as it was possible to obtain from suppliers. The mixture of 6 vol% CH4,
O2 corresponds to a methane:air ratio of 1:5 so 50 mL/min of CH4 and 250 mL/min of
air were introduced to a fuel cell formed by the previously proven steps, namely:
Table 7.3 Anode supported NiO-SDC//SDC//SCC-SDC solid oxide fuel cell fabrication regime
1. Compaction of 55wt%NiO-SDC mechanochemically milled nanopowder
composite mixed with 17% graphite pore former.
2. Application of 15mL 5wt% SDC in ethanol via aerosol to as-pressed pellet
held on a hot-plate at ca. 100°C
3. Sintering of anode/electrolyte at 1350°C for 4 hours
4. Application of 10%SDC-SCC (prepared by calcination of oxides) suspension
in ethanol via aerosol over an 8mm mask; pellet held at 120°C on hot-plate
5. Sintering of anode/electrolyte/cathode at 900°C for 2 hrs
6. Cell secured between Pt mesh in clamp, leads attached to flow cell terminal
The cells thus fabricated were found to have component thicknesses similar to the cell
shown in Figure 7.16. The specific dimensions were an anode thickness between 550-
700 μm, an electrolyte of 5-7 μm and a cathode of 25-40 μm. The fuel cell was tested
according to the method described previously, but it was found that the introduction of
air to methane at 400°C resulted in anode fracture and cell failure. The fracture was
presumably caused by thermal gradients induced by the rapid change in atmosphere
allowing catalytic combustion reactions to occur. Microstructural examination shown
in Figure 7.18 shows a network of cracks within the anode. The originally flat disc
had been considerably deformed due to these cracks and it is possible that the high
132
catalytic activity allowed localised methane combustion or decomposition reactions
that were detrimental to the anode’s operation.
Figure 7.18 Cross section of cell failed during exposure to CH4:Air atmosphere. Extensive crack
propogation in the lower anode region has caused the electrolyte and cathode (upper region) to also
fracture.
As a departure from the procedure set out above, the silver composite paste was
employed to achieve better contact between terminal lead and cathode. In one
instance the majority of the cathode surface was coated with the paste and a silver
wire was attached to this; in another instance the silver wire was made into a hoop and
this was attached in three locations with the paste to the cathode. In neither instance
was there a visible benefit to the use of Ag paste over Pt wire, evidenced by
comparable OCV and power generation characteristics. Furthermore the silver paste
did not prevent the degradation in cell performance previously associated with
133
cathode decomposition. This suggests that the use of silver, or gold, is not central to
the operation of a single chamber SOFC and platinum was preferred.
Cells fabricated according to the steps in Table 7.3 were heated under an atmosphere
of 50 mL/min CH4 with 250 mL/min air under direct control by monitoring the
resistance of the cell and its open circuit voltage. In doing so the rate of cell failure
due to rapid heating was significantly reduced. It was known from Section 7.1.1 that
of NiO would be reduced in an atmosphere of dilute methane:air as the temperature
approached 800°C. In the undiluted atmosphere the situation was similar, for the
anode material remained non-conductive and the cell was electrochemically inactive
when the annealing temperature was below 700°C. Around 750°C an OCV of 150
mV was generated and a small current could be drawn from the cell, indicating that
the NiO was reduced and the fuel cell was functioning. The output was recorded for
the temperature range between 550 and 650°C and as shown in Figure 7.19, the
maximum power generated was 5.5 mW/cm2 at 600°C.
134
0.0
1.0
2.0
3.0
4.0
5.0
6.0
0 2 4 6 8 10 12Current, mA
Pow
er, m
W/c
m2
600oC650oC550oC
Figure 7.19 Cell power generation from 50mL/min CH4, 250mL/min air
Despite the improvement in power generation by a factor of 2.5 compared to an
identically constructed fuel cell in dilute gas conditions, the result was far below those
reported by Hibino and Jasinski [53]. The cells of Hibino however were SDC
electrolyte supported, and in the more recent report from Jasinski was supported on a
Ni-YSZ anode and employed a LSCF cathode under a lean gas composition of 17
vol% CH4, so the cell processes are not directly comparable. This is illustrated by the
variation of OCV and maximum power with temperature, and as shown in Figure 7.20,
the maximum power was extracted at 600°C.
135
550
600
650
700
750
800
4
4.2
4.4
4.6
4.8
5
5.2
5.4
5.6
540 560 580 600 620 640 660
OC
V [m
V]
Pm
ax [m
W/c
m2]
Temperature [oC]
Figure 7.20 Variation of open-cell voltage and power output for SOFC in 50 mL/min CH4, 250
mL/min air
The OCV decreases with increasing temperature, here and in the dilute fuel case
(Figure 7.14) but when comparing power outputs there are significant differences.
The loss of power at higher temperatures has been ascribed to an increase in the
internal ohmic resistance, as measured by GCI and shown in Figure 7.21, increases
fractionally despite a concurrent decrease in polarisation resistance. Furthermore, an
ohmic resistance of this order is dominated by the electrolyte resistance, and since this
decreases with increasing temperature a 20% increase is significant. This
phenomenon is therefore related to the decomposition of a component, either chemical
136
or physical, both of which were observed in the cathodes and anodes tested in fuel rich
atmospheres. In this instance however, there was no visible degradation to the
cathode.
5
10
15
20
25
500 550 600 650 700
Rint [Ohm.cm2]
Rpol [Ohm.cm2]
R [O
hm.c
m2]
Temperature [oC]
Figure 7.21 Cell losses as a function of temperature as determined by GCI
The polarisation resistance of 4.1 Ω.cm2 is several times larger than the anode or
cathode reaction resistances reported by Hibino as 0.7 Ω.cm2 and 0.2 Ω.cm2
respectively [49]. Given that a proven cathode composition and fabrication method
was employed the cell losses must be associated with mass transport limitation of
methane through the anode so two possible causes are proposed. Firstly, the porosity
introduced by 16 vol% of graphite (Section 6.2.2.1) may be inadequate, despite the
137
fact that the cermet had open porosity and a comparable permeability to literature
examples. It was found difficult to increase the porosity beyond this limit due to
structural weakness of the anode oxide precursor. Secondly, although the NiO was
reduced to Ni below 800°C, it would appear that the triple-phase boundary area of 211
m/cm2 catalyses the complete combustion of CH4, rather than producing CO and H2
via partial oxidation or steam reformation as per Equations 4 and 5. In this single
chamber SOFC, the rate of CH4 oxidation and combustion appears to be greater than
the diffusion rate of CH4 to the electrochemical reaction zone close to the electrolyte
interface, thus forming a large anodic overpotential. The fact that on many occasions
the anode was observed to have suffered physical damage, presumably from methane
combustion, suggests that these types of fuel cells are not useful unless dilute gas
conditions are maintained. Hibino and others suggest that methane is not the ideal
fuel for a single chamber SOFC unless catalysts such as Ru, Pd or Pt are employed.
Hydrocarbons with longer chain lengths show a reduced tendency for carbon
deposition and yield higher cell potentials, but the use of nanostuctured Ni-SDC
anodes does not provide any advantage here and appears to be inherently
disadvantageous to cell operation. The vast range of Ni-SDC-pore geometries and
volume fractions that could be obtained using mechanochemical milling was beyond
the scope of this project, which has identified a number of fabricating and operating
parameters that will yield functioning fuel cells. It is hoped that this work will allow
others to further understand single-chamber SOFC’s and the application of
nanostructured materials to fuel cells in general.
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8 Conclusions
8.1 Summary
This study has endeavoured to apply a unique synthesis processing technique,
mechanochemical processing, to the fabrication of all components of a solid oxide
fuel cell using numerous processing techniques. The objective was to form thin
electrolytes and high internal surface area electrodes for the operation of the fuel cell
at intermediate temperatures from a fuel mixture of methane and air.
Mechanochemical synthesis of the electrolyte, samarium-doped ceria (SDC), yielded
particles of high-purity with dimensions below 20 nm. When densified, the ionic
conductivity was unremarkable and the grains had grown to several microns in size.
Compositing SDC with the anode precursor oxide, NiO, using mechanochemical
milling formed a nanocomposite oxide and the two phases were completely separate
with a grain size of around 15 nm. The densification of this composite proceeded
below 1350°C and the size of individual grains remained well below 100 nm. The
electrical conductivity of these Ni-SDC composites was 1068 S/cm at 600°C and
image analysis suggested a triple-phase boundary area of 211 m/cm2 and this was
consistent with the ultra-fine scale of the cermet’s microstructure.
The fabrication of an anode supported cell required the deposition of SDC electrolyte
by spray-painting, sincemethods such as spin-coating were found to be impractical.
Additionally, it was not possible to synthesize the cathode material, samarium-doped
strontium cobaltite cathode, using mechanochemical techniques so conventional and
commercial sources were employed.
139
The fuel cells thus fabricated were able to operate in dilute, commercially supplied
mixtures of methane and air, producing a specific power density of 483 μW/cm2 at
650°C. Operation in 25 vol% CH4 in air was hindered by the very same ultra-fine
anode microstructure that was so painstakingly fabricated. Anodic polarisation due to
mass transport limitation was attributed to the overly large polarisation resistance of
4.1 Ω.cm2. It was proposed that the high triple-phase boundary area, in conjunction
with insufficient porosity, caused the catalytic combustion of methane to occur in the
mixed fuel atmosphere.
This dissertation contributes much to the understanding of nanoscale synthesis
techniques and their applicability to alternative power generation solutions. These are
both timely topics, for the direct generation of electricity from natural gas would solve
many problems associated with the production and distribution of energy faced by
both developing and developed nations today.
8.2 Suggestions for Further Work
The methodologies and outcomes of this project can provide the basis for many
avenues of further work. Firstly, although it does not appear that nanostructured
anodes are of revolutionary significance to solid oxide fuel cells, there are a number of
other technological applications where it remains untested. Remaining focussed on
SOFC technology, there are numerous permutations in the anode structure, the exact
cathode composition and the hydrocarbon fuel that could define many similar studies.
The stability of the cathode material appeared to be an issue and given more time this
would have been investigated further. Impedance spectroscopy would have provided
140
a more detailed understanding of internal resistances and Further specific examples
are provided.
i. It was evident that the scale of the anode microstructure might not be as
important as the relative scale of catalyst and electrolyte. Large, elongated
particles of catalyst phase with a fine dispersion of electrolyte over them
would be ideal, but there are no methods to synthesize anisotropic particles of
(doped) NiO. Whether mechanochemically synthesized electrolyte particles
would be employed, or that flaky or rod-like particles can be obtained using
some combination of techniques are largely unknown.
ii. In single chamber SOFC it has been shown that higher chain length
hydrocarbons, in particular propane and butane, give higher performance than
methane. Furthermore, the use of Cu-Ni anodes has been shown to limit
graphite deposition [119]. This thesis has provided a framework for the
fabrication of Ni-based anodes and the utility of other more reactive
hydrocarbons suggests that this would be a worthwhile pursuit.
iii. Doped-NiO oxide composites with YSZ can be employed as thermoelectric
materials and as cathodes for molten carbonate fuel cells [108, 120]. The
electrical properties of nanoscale composites were briefly visited in this project
and NiO doping occurred inadvertently. A more thorough experimental
investigation of potential dopants and composite compositions could generate
interesting materials and contribute significantly to this field.
141
iv. Fabrication of small tubular electrolytes using electrophoretic deposition.
Suspensions of mechanochemically synthesized nanopowders may be suitable
for deposition onto graphite cores as demonstrated by Sarkar and Rho [121]
for a micro-tubular SOFC. These tubes should not be limited only to SOFC
research as they would find application in oxygen pumps/sensors and ion-
selective electrodes. The consolidation of a ceramic from the green state
would be challenging as evidenced by the many processes investigated in this
dissertation.
142
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10 Publications and Presentations during PhD
Canditature
1. “Mechanochemical synthesis and characterisation of nanoparticulate
samarium-doped cerium oxide” poster presented at NANO2002, The 6th
International Conference of Nanostructured Materials in Orlando, Florida.
Subsequently published in Scripta Materialia 48(2003) 85-90.
2. “Slip Characterisation and Tape Casting of Cerium Oxide Nanopowder” oral
presentation at AUSTCERAM 2002, the Exhibition of the Australasian
Ceramic Society and subsequently published in The Journal of the
Australasian Ceramic Society 38(2002) 139-44.
3. “Evaluation of Mechanochemically Synthesized NiO/SDC Composite
Nanopowders for the Development of Nanostructured Cermet Anodes”, oral
presentation at SOFC VIII Paris, France 2003. Published in the 203rd Meeting
of the Electrochemical Society, Proceedings Volume 7, 752-61.
4. “Mechanochemically Synthesized Nanopowders for High Performance Solid
Oxide Fuel Cell Materials”, Poster presentation at the2003 Sir Mark Oliphant
Conference “Scaling Down to a Nano-Materials World” Melbourne, Australia.
150
