CHARACTERIZATION OF 10MOL% Sc 2 O 3 -1MOL% CeO 2 -ZrO 2 CERAMICS AS ELECTROLYTE MATERIAL FOR LOWER TEMPERATURE SOLID OXIDE FUEL CELLS by Devendra Ray A thesis submitted to the graduate faculty in partial fulfillment of the requirements for the degree of MASTER OF SCIENCE Department: Mechanical & Chemical Engineering Major: Mechanical Engineering Major Professor: Dr. Jag Sankar North Carolina A&T State University Greensboro, North Carolina 2007
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CHARACTERIZATION OF 10MOL% Sc2O3-1MOL% CeO2-ZrO2 CERAMICS
Material properties such as hardness, fracture toughness, densities and coefficients of thermal expansion are very important parameters for stability and reliability of solid oxide fuel cell (SOFC) operation. Ceria doped Sc2O3-CeO2- ZrO2 ceramics are also very promising materials as electrolyte in solid oxide fuel cells (SOFC) due to their high oxygen conductivity in the 800-1000°C temperature range. In this experimental work, sintering behavior of two commercial powders with the nominal composition 10mol% Sc2O3-1mol% CeO2-ZrO2 which were produced by Praxair surface technologies, USA and DKKK, Japan was studied. Ceramics were made by sintering uniaxially pressed pallets at temperature range 1100-1600°C in air. Hardness and fracture toughness of ceramics sintered at different temperatures were studied by microindentation method. The density measurement was done using Archimedes’ principle. Porosity level in ceramics was estimated using actual and theoretical densities calculated using lattice parameter a=5.09 Ǻ of FCC unit cell obtained from XRD data and actual composition of powder. Hardness of most dense ceramics reaches 15GPa and fracture toughness is in the range of 1.8-2.5MPa×m 0.5. These results are in good agreement. Hardness of ceramics depends on indentation depth due to indentation size effect (ISE). Indentation size effect parameters were analyzed using Power law model. CTEs were studied. These results can be further used for the optimal design of SOFC layered structures as well as for determination of their reliability and durability under operational conditions.
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CHARACTERIZATION OF 10MOL% Sc2O3-1MOL% CeO2-ZrO2 CERAMICS AS ELECTROLYTE MATERIAL FOR LOWER TEMPERATURE
SOLID OXIDE FUEL CELLS
by
Devendra Ray
A thesis submitted to the graduate faculty in partial fulfillment of the requirements for the degree of
MASTER OF SCIENCE
Department: Mechanical & Chemical Engineering Major: Mechanical Engineering Major Professor: Dr. Jag Sankar
North Carolina A&T State University Greensboro, North Carolina
2007
iii
DEDICATION
First and foremost, I would like to thank my Lord Shiva for giving me the
strength and passion to arrive at this point and for putting the right people in my life at
the right time to help guide me. I dedicate this thesis to my parents, Ramkailash Ray,
father, and late mother, Ramshakhi Devi Ray, for their love, encouragement and the
many sacrifices they made that have enabled me to arrive at this point of my life. I would
also like to dedicate this thesis to my wife, Sima Ray. Finally, I would like to thank all
my family and friends who have encouraged and supported me to reach this milestone in
my life.
iv
BIOGRAPHICAL SKETCH
Devendra Ray was born in Nepal and immigrated to the United States of America
(USA) on September 7, 2004. In 2002, he graduated from Kathmandu University, Nepal
with an undergraduate degree in Mechanical Engineering. He is a candidate for the
Master of Science degree in Mechanical Engineering at North Carolina Agricultural and
Technical State University. He has performed research on solid oxide fuel cells (SOFC)
electrolyte materials.
v
ACKNOWLEDGEMENTS
I would like to begin by thanking my major advisor, Dr. Jag Sankar, and senior
research scientist/co-advisor, Dr. Sergey Yarmolenko, for their consistent guidance and
support over the past two years. I significantly benefited from their knowledge and
valuable experience in the field of material science and its characterization. I would also
like to thank all my friends and the staff of the Center for Advanced Materials and Smart
Structures (CAMSS). I would also like to express my thanks to Dr. Z. Xu and Dr. M. T.
Saad for their suggestions. I am very grateful to my family for their support, sacrifice,
and encouragement in assisting me to achieve this goal.
vi
TABLE OF CONTENTS
LIST OF FIGURES ........................................................................................................... ix
LIST OF TABLES............................................................................................................ xii
ABSTRACT..................................................................................................................... xiii
2.1 Schematic representation of fuel cell ...........................................................................4
2.2 Phase diagram of Sc2O3-ZrO2 system .......................................................................11
3.1 Illustration of changes in the surface morphology after polishing on diamond paper with grits 3µm, 1 µm, and 0.5µm (from left to right) ......................................21 3.2 A schematic representation of the indentation process showing various parameters involved in the analysis ...........................................................................26
3.4 Typical images collected at 500x on optical microscope for hardness study ............33
3.5 (a) Schematic figure having crack and indent length; (b) Image collected at 1000x on optical microscope (Indentation load is 200g)..................................................35
3.6 Typical SEM images for grain size study ..................................................................38
3.7 Typical AFM working principle ................................................................................39
3.8 (a) Typical AFM images for grain size study; (b) Grain size and cavitations measurements............................................................................................................42 3.9 Temperature dependent shifting of XRD peaks positions ........................................45
3.10 Constant, linear and quadratic fitting with temperature vs. log (LT) (P1600- powder) .....................................................................................................................47 3.11 Instrumentally set temperature vs. residuals (P1600-powder)..................................48
3.12 Polynomial fitting of calibrated temperature vs. instrumentally set temperature ...............................................................................................................49
4.1 SEM micrographs of powders (a-Praxair, USA) and (b-DKKK, Japan)...................53
x
4.2 XRD pattern at different temperatures.......................................................................54
4.3 XRD pattern of as-received Praxair powder..............................................................55
4.4 XRD pattern of as-received DKKK powder..............................................................56
4.5 Fracture surfaces of samples P1600-(a) and J1600-(b)..............................................59 4.6 Grain structure of samples produced from Praxair powder sintered at temperatures 1500ºC (left) and 1600ºC (right) ..........................................................60 4.7 Grain structure of samples produced from DKKK powder sintered at temperatures 1400ºC (left), 1500ºC (middle) and 1600ºC (right) ...................................................60 4.8 SEM images of polished and low angle ion-milled surface of sample P1600 (a) and thermally etched surfaces of samples P1600 (b) and J1600(c).................................61 4.9 The dependence of density on sintering temperatures...............................................62
4.10 Porosity level vs. sintering temperature....................................................................62
4.11 SEM images of DKKK samples from left to right sinter correspondingly 1200-1600°C.............................................................................................................64 4.12 SEM images of Praxair samples from left to right sinter correspondingly 1500-1600°C.............................................................................................................64 4.13 Grain size vs. sintering temperatures ........................................................................65 4.14 SEM images collected after 1hr ion milling at 30 deg (left - J1600 right - P1600,
right - P1600) ............................................................................................................66 4.15 SEM images collected after ion milling (left-J1600 milled at 9-deg for 18hrs, right - P1600 milled at 5 deg for 20hrs)....................................................................67 4.16 SEM images of P1600 collected after ion milled at 5 deg for 20hrs (right - high porous area, left - less porous) ..........................................................................67 4.17 Phase content in the sample DKKK-1200 at different temperatures from XRD (Retvield multiphase analysis) ........................................................................69 4.18 Changes in XRD pattern of powdered sample J1600 heated from 300 to 500°C with step 10°C...........................................................................................................70
xi
4.19 Phase transition in DKKK ceramic sintered at 1600°C............................................70
4.20 Instrumentally set temperature vs. residuals (Praxair powder).................................74
4.21 Constant, linear and quadratic fitting with temperature vs. log (LT) (Praxair powder) .....................................................................................................................74
4.22 Instrumentally set temperature vs. residual(DKKK powder) ...................................75
4.23 Constant, linear and quadratic fitting with temperature vs. log (LT) (DKKK powder) .......................................................................................................75 4.24 Instrumentally set temperature vs. residuals (J1500 Powder) ..................................76
4.25 Constant, linear and quadratic fitting with temperature vs. log (LT) (J1500 Powder) .....................................................................................................................76 4.26 Instrumentally set temperature vs. residuals (J1500 ceramic) ..................................77
4.27 Constant, linear and quadratic fitting with temperature vs. log (LT) (J1500 ceramic).....................................................................................................................77 4.28 Hardness vs. sintering temperature plot for etched and non-etched sample.............78
4.29 Variation in fracture toughness value with temperature ...........................................79 4.30 (a) Effect of sintering temperature on microhardness and fracture toughness at the load 1000g; (b) Micrograph of typical Vickers indents and cracks at different loads (thermally etched J1600) ..................................................................81 4.31 ISE analysis using power law (a) and Nix-Gao (b) models......................................86
4.4 Density and porosity of sintered samples ..................................................................58
4.5 Grain size and in-grain cavities at different sintering temperatures ..........................65
4.6 Coefficients of thermal expansion for Praxair and DKKK samples in ceramic and powder forms ......................................................................................................73 4.7 Hardness (H) and fracture toughness (K1C) values at different sintering temperatures and loads for non-etched samples .......................................................82 4.8 Hardness (H) and fracture toughness (K1C) values at different sintering temperatures and loads for thermally etched samples ...............................................83 4.9 Comparison of hardness (H) and fracture toughness (K1C) of non-etched and etched samples. .................................................................................................84 4.10 ISE parameters for Praxair and DKKK-based samples ............................................86
Ray, Devendra. CHARACTERIZATION OF 10MOL% Sc2O3-1MOL% CeO2-ZrO2 CERAMICS AS ELECTROLYTE MATERIAL FOR LOWER TEMPERATURE SOLID OXIDE FUEL CELLS (Major Professor: Dr. Jag Sankar), North Carolina Agricultural and Technical State University.
Material properties such as hardness, fracture toughness, densities and
coefficients of thermal expansion are very important parameters for stability and
simpler and less expensive than lithium ion batteries [2].
2.5 Basic Components of SOFC
There are different types of electrolytes that are commercially available; however,
more research and development is needed to characterize and explore new and promising
electrolytes for an intermediate temperature (IT) SOFC. The basic components of SOFC
are:
1. Cathode
2. Electrolyte
3. Anode
4. Interconnect
This thesis focuses on electrolyte material.
8
2.5.1 Electrolyte
The most effective electrolytes employed in fuel cells have ionic conductivities of
around 0.10 Ώ-1cm-1. The search for better electrolytes has led to the development of
three major candidate material classes for fuel cells: aqueous, polymer, and ceramic
electrolytes. Regardless of the class, however, any fuel cell electrolyte must meet the
following typical properties to be a better electrolyte:
• High ionic conductivity
• Low electronic conductivity
• High stability (in both oxidizing and reducing environments)
• Low fuel crossover
• Reasonable mechanical strength (if solid)
• Ease of manufacturability.
Except for the high conductivity requirement, the electrolyte stability requirement
is often the hardest to fulfill. It is difficult to find an electrolyte that is stable in both the
highly reducing environment of the anode and the highly oxidizing environment of the
cathode [1, 52].
2.5.2 Electrolyte for Solid Oxide Fuel Cells
The electrolyte is the central part of an SOFC. Within the electrolyte the oxygen
ions (O2-), which are reduced on the air electrode side (cathode), are transported and
react with, for example, hydrogen to form water on the fuel electrode side (anode).
Conversely, electrons (e-) are formed and while moving in the opposite direction, are
available for an outside current use. Nowadays, the most frequently used electrolyte
9
material is zirconia (ZrO2). Zirconia is at ambient conditions a poor ionic conductor. If
ZrO2 is heated up to temperatures above 2000°C, it becomes ionic conducting due to a
phase transformation from tetragonal to cubic structure. By adding stabilizing scandia
(ScSZ) the cubic structure is stabilized even at ambient conditions. Because of this
stabilization zirconia becomes a reasonable ionic conductor at SOFC operating
temperatures (750-1000°C) and can be used as electrolyte material. Ionic conduction
proceeds through oxygen vacancies due to insertion of a di- or trivalent element (Ca2+,
Y3+, Sc3+) instead of the tetravalent Zr4+. The lack of positive charge is balanced by free
oxygen lattice sites. Through these free sites oxygen can move through the cubic
structure.
Besides good ionic conductivity, gas tightness of the electrolyte is the most
important characteristic the material should have. If the gas tightness of the electrolyte is
insufficient, a reaction between oxygen (cathode side) and hydrogen (anode side) may
occur. Gas tightness is ensured by sintering the electrolyte at temperatures of
approximately 1400°C. Lower temperatures lead to inadequate gas tightness, but pure
yittria-stablized zirconia (YSZ) cannot be sintered to high densities below 1400°C. In
contrast to the high sintering temperature, which is necessary for the electrolyte, the
cathode tolerates only maximum temperatures of approximately 1200°C. If sintered at
higher values, the amount of triple phase boundaries reduces drastically due to enhanced
sintering. These two facts are the reason for one area of major research and development
on electrolyte materials, namely, is the reduction of the sintering temperature of the
electrolyte (goal ≤ 1300°C). Decreasing sintering temperatures can be reached by a)
10
using a nanosized starting material or b) the use of sintering additives. Additional
research and development is focused on the coating technologies for SOFC electrolytes
[4].
2.6 Introduction of Zirconia
Zirconia has become one of the most industrially important ceramic materials of
the present time. The traditional applications of ZrO2 and ZrO2-containing materials are
foundry sands and flours, refractory ceramic, and abrasives. Because of its high oxygen
ion conduction and high refractive index, it is also used in a wide range of newer
applications which include fuel cells, catalysts, oxygen sensors, and jewelry. Zirconias
have also been utilized in many mechanical applications. Along with high strength and
toughness, zirconia also possesses good hardness, wear resistance, and thermal shock
resistance. These properties have led to the use of zirconia-based components in a
number of engineering applications such as automobile engine parts, wire drawing dies
and cutting tools. The low thermal conductivity together with relatively high coefficiency
of thermal expansion makes zirconia a suitable material for thermal barrier coating on
metal components [6].
2.6.1 Crystallographic Structure
At atmospheric pressure, pure zirconia (ZrO2) has three crystalline polymorphs
with monoclinic, tetragonal and cubic structures. The monoclinic from is stable up to
1170°C when it transforms to the tetragonal modification which remains stable up to
2370°C; from 237°C to the melting point (2680°C), the cubic form is stable. The reverse
11
transformations take place on cooling. The phase transformations are martensitic in
nature. The monoclinic and tetragonal conversion is accompanied by a contraction in
volume of approximately 5% which can cause mechanical failure evident by cracking in
the ceramic. The phase transformation can be avoided and the high temperature cubic
phase can be stabilized at low temperature by substituting low-valency cations for the
zirconium (see Figure 2.2).
Figure 2.2: Phase diagram of Sc2O3-ZrO2 system
Cubic ZrO2 has the fluorite structure with the O2- ions arranged in simple cubic
packing and half the interstices in this lattice occupied by Zr4+ ions. The substitution of
12
lower-valency cations leads to O2- ion vacancies. The vacancies that stabilized the
structure also lead to high mobility in the oxygen sublattice and to behavior as a high-ion
conductor.
The elements that stabilized the cubic fluorite structure in zirconia include
scandium, lanthanides, yttrium, magnesium, calcium, manganese and indium. Scandium
has been found to give a material with a higher conductivity, which is particularly
valuable at a lower temperature. Scandia stabilized zirconia (ScSZ) is one of the solutions
to the problem of stabilizing the cubic form of zirconia [6].
2.6.2 Application of Scandia Stabilized Zirconia (ScSZ)
Fully stabilized zirconia based ceramic materials have several areas of application
such as an electrolyte in solid oxide fuel cells (SOFC), oxygen sensors and structural
applications due to their high oxygen conductivity in the 800-1000oC temperature range,
electrical and mechanical properties. Scandia stabilized zirconia (ScSZ) ceramics have
the highest oxygen conductivity among zirconia based materials [7] and, with the trend in
reducing the operation temperature of SOFCs to an intermediate temperature (IT) range
of 600-800ºC, ScSZ ceramics are very attractive materials for IT SOFCs and catalytic
membrane reactors. Important requirements for these applications are reproducibility of
the electrical and mechanical properties, high density and a low level of internal flaws.
However, most commercial ZrO2-based powders typically need high sintering
temperatures above 1400ºC in order to form dense ceramics. High sintering temperatures
create a variety of problems such as temperature induced degradation and shape change
of materials, interface reactions between zirconia-based ceramics and other components,
13
excessive grain growth as well as energy production costs. Lower processing
temperatures can limit grain growth to nanocrystalline size which has beneficial effect on
the improvement of electrical properties of solid electrolytes [8] due to specific grain
boundary conductivity. Recent studies of nanocrystalline scandia stabilized zirconia thin
films prepared by sol-gel using a polymer precursor solution showed that the electrical
transport of this material can also be enhanced [9]. Therefore the reduction of the
sintering temperature towards an SOFC�s operating temperature can improve mechanical
and electrical properties and make possible co-firing zirconia-based electrolytes with
other SOFC components and significantly lower production costs. Properties such as
hardness, fracture toughness and coefficients of thermal expansion (CTEs) are very
important parameters for stability and reliability of SOFC operation.
Extensive research has shown that a phase transition from cubic to rhombohedral
structure of 11mol%Sc2O3-89%ZrO2 occurs when the temperature decreases below
600°C [10, 11]. This phase transformation prevents the direct use of this ceramic in
SOFC applications. The aging process of scandia doped zirconia is also accompanied by
the formation of a less conductive rhombohedral phase resulting in degradation of ionic
conductivity of Sc2O3-ZrO2 electrolytes [12, 13]. Usually the cubic phase can be
stabilized over a wide temperature range by substitution of 1 mol% of scandia by another
stabilizer such as Yb2O3 [13], Y2O3 [10], Bi2O3 [14], Al2O3 [15], and so on. The addition
of a low amount of other oxides usually does not greatly deteriorate the sintering activity
and other properties of the resulting powder. It was reported recently that replacement of
1mol% Sc2O3 by CeO2 stabilized cubic phase led to only 50ºC increase in sintering
14
temperature [14]. Also it was found recently that scandia-ceria doped zirconia ceramics
have a very high ionic conductivity (0.151 S/cm at 800oC) [16].
In this experimental work the study of sintering behavior of two commercially
available powders with a nominal composition of 10mol%Sc2O3-1mol%CeO2-ZrO2
(ScCeSZ) that can be used to fabricate dense and stable ceramic electrolytes was carried
out. The crystal structure, the morphology and phase transitions of powders and ceramics
were also investigated. Mechanical properties, such as hardness and fracture toughness,
of the Sc2O3-CeO2-ZrO2 ceramics were reported. These results can be further used for the
optimal design of SOFC layered structures as well as for determination of their reliability
and durability under operational conditions.
2.6.3 Different Electrolyte Materials
Table 2.2 briefly describes all types of commercially available SOFC electrolyte
materials currently available in different forms for developing solid oxide fuel (SOFC)
[18].
Table 2.2. Commercially available types of electrolyte materials
No Name Form Description 1 Scandia
Stabilized Zirconia (10 mole %) Nanopowder
Nano Powder Scandia Stabilized Zirconia (Sc2O3)0.10(ZrO2)0.90 Nanoscale powder for use as sintering aid, catalyst support, and as a component for mixed conducting anodes and cathodes to enhance catalytic activity Surface area: >100 m2/g Particle size: softly agglomerated
15
Table 2.2. (Continued) No Name Form Description 2 Gadolinium
Doped Ceria (10% Gd) Nanopowder
Nanopowder Gadolinium Doped Ceria Gd0.10 Ce0.90 O2- Nanoscale powder for use as sintering aid, catalyst support, and as a component for mixed conducting anodes and cathodes to enhance catalytic activity. Surface area: >100 m2/g Primary crystallite size: 5-10 nanometers Secondary particle size: softly aggolomerated
Suspension Gadolinium Doped Ceria (10% Gd) aqueous suspension specially formulated suspension for colloidal deposition and spray coating methods. Suspension is stable for an extended period of time. Surface area: >20-40 m2/g (estimated) Particle size: 50-80 nm Solids loading: 10-15 volume %
4 Gadolinium Doped Ceria (10% Gd) Ceramic Powder
Ceramic Grade Powder
Gadolinium Doped Ceria Gd0.10 Ce0.90 O2- Powder suitable for tape casting, ink manufacture, pellet pressing and other non-aqueous manufacturing processes. Surface area: >5-8 m2/g Particle Size: d50-0.5 microns
5 Gadolinium Doped Ceria (20% Gd) Nanopowder
Nanopowder Gadolinium Doped Ceria Gd0.20 Ce0.80 O2- Nanoscale powder for use as sintering aid, catalyst support, and as a component for mixed conducting anodes and cathodes to enhance catalytic activity. Surface area:>100m2/g Primary crystallite size: 5-10 nanometers Secondary particle size: softly agglomerated
6 Gadolinium Doped Ceria (20% Gd) Ceramic Powder
Ceramic Grade Powder
Gadolinium Doped Ceria Gd0.20 Ce0.80 O2- Powder suitable for tape casting, ink manufacture, pellet pressing and other non-aqueous manufacturing processes. Surface area: >5-8 m2/g Particle Size: d50-0.5 microns
16
Table 2.2. (Continued) No Name Form Description 7 Samarium
Doped Ceria (15% Sm) Nanopowder
Nanopowder Samarium Doped Ceria Sm0.15 Ce0.85 O2- Nanoscale powder for use as sintering aid, catalyst support, and as a component for mixed conducting anodes and cathodes to enhance catalytic activity. Surface area: >100 m2/g Particle size: softly agglomerated
8 Samarium Doped Ceria (15% Sm) Ceramic Powder
Ceramic Grade Powder
Samarium Doped Ceria Sm0.15 Ce0.85 O2- Powder suitable for tape casting, ink manufacture, pellet pressing and other non- aqueous manufacturing processes. Surface area: >5-8 m2/g Particle Size: d50-0.5 microns
Ytrria Stabilized Zirconia (Y2O3)0.03 (ZrO2)0.97 Partially stabilized zirconia powder suitable for pellet pressing, injection molding and aqueous manufacturing processes. Can be calcined for use in non-aqueous applications. Surface area: >13-19 m2/g Particle Size: d50-0.5 microns
10 YSZ 8 mole% Engineered Coating
Suspension Yttria Stabilized Zirconia (8% Yttria) Aqueous suspension for colloidal deposition, dip coating and spray coating methods. Has engineered particle sizes to match shrinkage to typical anode materials. Can be engineered to match the shrinkage of cathode supports. Suspension is stable for an extended period of time. Surface area: Engineered Particle Size: Engineered Solid loading: 10-15 volume %
11 Yttria-Stabilized Zirconia 8 mole% Nanopowder
Nanopowder Ytrria Stabilized Zirconia (Y2O3)0.08 (ZrO2)0.92 Nanoscale powder for use as sintering aid, catalyst support, and as a component for mixed conducting anodes and cathodes to enhance catalytic activity. Surface area: >100 m2/g Particle size: softly agglomerated
17
Table 2.2. (Continued) No Name Form Description 12 Yttria-
Stabilized Zirconia 8 mole% Aqueous Suspens
Suspension
Yttria Stabilized Zirconia (8% Yttria) Aqueous Suspension specially formulated suspension for colloidal deposition and spray coating methods. Suspension is stable for an extended period of time. Surface area: 20-40 m2/g (estimated) Particle Size: 50-80 nm Solid loading: 10-15 volume %
The critical parameters used in the calculation of hardness and elastic modulus are
the peak load (Pmax), the maximum indentation depth (hmax), the residual or final
indentation depth after unloading (hf) and the initial unloading stiffness (S = dP/dh). S is
also known as the elastic contact stiffness, or simply contact stiffness. Hardness, the
mean pressure a material can withstand, is calculated using Equation 1,
APH = (1)
28
where P is the load and A is the projected contact area at that load. The hardness
determined by dividing the applied load by the projected contact area under the load
should not be confused with the traditional definition of hardness, the load divided by the
projected area of contact of the residual hardness impression. These two definitions yield
similar values for hardness when the deformation process is mainly dominated by a
plastic region and a fully plastic permanent hardness impression is formed. When the
contact is predominantly elastic, they exhibit different values for hardness since the
residual contact area is very small for a purely elastic contact, which results in infinite
hardness based on the traditional hardness definition [30]. This phenomenon plays a
crucial role while indenting materials with very sharp indenters at small indentation
depths. In these conditions, traditional definition of hardness yields a higher value than
that calculated using Equation 2.
Elastic modulus of the test material, E, is calculated using the equation,
i
i
r EEE
22 111 νν −+−= (2)
where Er is the reduced elastic modulus, E and ν are the elastic modulus and Poisson�s
ratio of the test material, and Ei and νi are the elastic modulus and Poisson�s ratio of the
indenter, respectively, and for a diamond indenter these values are 1141 GPa and 0.07
respectively. Reduced modulus is used to account for the elastic displacements that occur
in both the indenter and the sample. Er is related to the contact stiffness by the equation,
A
SEr βπ
2= (3)
29
where S is the contact stiffness, A is the projected contact area and β is a constant that
depends on the geometry of the indenter. This equation is derived from elastic contact
theory [41-42] and holds good for any indenter that can be described as a body of
revolution of smooth function [43]. Though the equation formally applies only to circular
contacts, it has been shown that it works well for different geometries provided a
different value of β is used [44-45]. For indenters with square cross sections β = 1.012;
for triangular cross sections like the Berkovich and cube-corner indenters, β = 1.034. It
has been shown recently that yet another correction factor has to be added to the equation
[42, 46-49].
The principal difference between conventional hardness testing techniques and
the nanoindentation technique is the manner in which the contact area is determined.
Instead of imaging of the indent impressions, the contact area is estimated from the
indentation load-displacement data. Oliver and Pharr [22] developed a method by fitting
the unloading portion of the load-displacement data to the power-law relation
mfhhAP )( −= (4)
where A and m are constants determined by a least squares fitting procedure; hf is the
final displacement after complete unloading determined from curve fit. The initial
unloading slope is found by differentiating the equation and evaluating the derivative at
the peak load and displacement. Contact stiffness is obtained by differentiating Equation
(4) and evaluating the value at the maximum indentation depth.
The contact depth, hc, is estimated using, the expression,
30
SPhhc ε−= (5)
where ε is a constant that depends on the indenter geometry. For spherical indenters ε =
0.75, for conical indenters ε = 0.72 and for Berkovich indenters ε = 0.75. Equation (5)
does not account for pile-up formed during indentation because it is assumed that the
contact is perfectly elastic and only sink-in occurs continuously [30]. The contact area is
calculated as a function of the contact depth, hc, and is written as:
A (hc) = 24.5 hc2 (6)
for a perfect Berkovich indenter. Taking into consideration the tip blunting effects, the
modified area function is written as,
128/18
64/17
32/16
16/15
8/14
4/13
2/12
11
25.24)( cccccccccc hChChChChChChChChhA ++++++++=
(7)
where C1 through C8 are constants. The area function is determined by making a series of
indentations at various depths in a calibration sample of well-known elastic properties.
The basic assumption in this process is that the elastic modulus is independent of the
indentation depth. It is also imperative that there is no pile-up. Fused silica is the widely
accepted calibration sample. The machine compliance is also calculated during the
calibration procedure [22]. Correcting for machine compliance, the load-displacement
data are reduced and used to obtain the contact stiffness (S) and the contact depths (hc).
For fused silica material, E = 72 GPa; ν = 0.17 and for diamond indenter, E = 1141 GPa;
ν = 0.07. The contact area, determined using Equation (6), is plotted against contact
depth, hc, and is fitted according to a polynomial order. A weighted fitting procedure
31
ensures that data from all depths are given equal importance. The area function
coefficients determined from the fit are used in further calculations.
3.4.2.2 Microindentation
The resistance to indentation or deformation of material is known as its hardness.
It is one of the most frequently measured properties of a ceramic. There are many
hardness scales and methods of measurements to characterize resistance to deformation,
densification, and fracture. In fact, many ceramic specifications list minimum hardness
requirements. For example, a new ASTM zirconia�s specification for surgical implants, F
1873-98, stipulates that Vickers hardness (HV) shall be no less than 11.8 GPa (1200 kg/
mm2) at a load of 9.8 N (1 kg). Although measuring and interpreting ceramic hardness
should be routine, pitfalls, controversies, and surprises abound.
The Vickers hardness test was used for hardness and fracture toughness
measurements in this research work. In most engineering and characterization
applications, approximately 60% of worldwide published ceramic hardness values are
Vickers, with loads typically in the range of a few newtons to 9.8 N (1 kg) and occasional
data for soft or high-toughness ceramics as high as 98 N (10 kg). Knoop hardness is
more frequently used for glass, glass ceramics, and ceramic white wares. About 35% are
Knoop with loads from as low as 0.98 N (100 g) to 19.6 N (2 kg). Rockwell hardness are
about 5% [21]. For research purposes, Vickers, Knoop, and Berkovich (triangular
pyramid) indenters are customary; Rockwell and Brinell indenters are rarely suitable for
ceramics research. Measurements were carried out on conventional microhardness
machines with Vickers diamond indenters.
32
A LECO Microhardness testing Machine Model: M-400-H1/H2/H3 was used for
indentations. The indentions loads were 100, 200, 300, 500, 1000g. At each load 20
indents were created. The following procedures were undertaken for indentations:
• A highly polished flat sample was placed on stage at a right location of
microhardness testing machine.
• The sample was focused and adjusted to create 20 indents at each load from 100,
200, 300, 500, 1000g.
• A loaded indentor was lowered against a polished flat surface of the sample for a
specific load time of 30 seconds and retracted automatically afterwards.
• During loading, the indentor penetrated into the ceramic and, on retracting, left a
permanent pyramidal indentation.
• A high quality impression image was collected on an Optical Microscope
adjusting the indent on the center of the lens to avoid possible errors of focusing
on Zeiss Axiovert 10 at magnification 500x (1000, 500g) and 1000x (100, 200,
300g).
• Diagonal lengths of the impression were measured by using Image Pro-Plus
software for Vickers hardness.
• Crack lengths were measured for fracture toughness analysis.
A typical impression is shown in Figure 3.4.
33
Figure 3.4: Typical images collected at 500x on optical microscope for hardness study
The Vickers hardness is calculated by using the following formula
= 28544.1
DFVHN GPa
where, F = Load in Kg, D = average diagonal length in µm, VH = kg/mm2 (1 kgf/mm2 =
9.81 MPa.)
Fracture toughness, K1C, is a measure of a ceramic part�s resistance to fracture
starting from a pre-existing crack. It is one of the most important properties of material
for virtually all design applications. If the material has a large value of fracture
toughness, it will undergo ductile fracture. Brittle fracture is a characteristic of materials
with a low fracture toughness value. There are four different types of fracture toughness,
34
KC, K1C, K11C, and K111C. KC is the fracture toughness of the sample, which is in a state
plane stress. K1C, K11C, and K111C are the fracture toughness of material under three
different modes of fracture, mode I, mode II, mode III, respectively.
Ceramic materials have a much lower K1C value than the metals. The low K1C
value shows that ceramic materials are very susceptible to cracks and undergo brittle
fracture, whereas the metals undergo ductile fracture. Thus for this study, the fracture
toughness, K1C is considered.
For fracture toughness calculation Young�s modulus was obtained from
nanoindentation data 230GPa and 210GPa for DKKK, Japan and Praxair samples
respectively.
The equations used for fracture toughness calculation is:
2/3
2/1
1 016.0C
PHEK C
=
where, E = Young�s modulus (230/210GPa), H = Vickers hardness (GPa), F =
Indentation load (mN), C = Crack length (µm). Figure 3.5 shows diagonals and cracks
taken for measurements. Young�s modulus was obtained by nanoindentation technique.
The Young�s modulus value for the DKKK based ceramic sample sintered at 1600°C was
230GPa. The same value for the Praxair based ceramic sintered at 1600°C was 210GPa.
Both values are in good agreement with the literature review of this experimental work.
35
(a)
.
(b)
Figure 3.5: (a) Schematic figure having crack and indent length; (b) Image collected at 1000x on optical microscope (Indentation load is 200g)
D
C
36
3.5 Microstructure Characterization
Microstructure observations of the Scandia Stabilized Zirconia (ScSZ) were
performed on the thermally etched samples of DKKK, Japan and Praxair materials using
a scanning electron microscope (SEM) Model: Hitachi S3000N and atomic force
microscope (AFM) respectively. A brief description of both microscopes and the
technique adopted for the collection of images follows.
Scanning Electron Microscope
The SEM uses electrons rather than light to form an image. There are many
advantages to using the SEM instead of a light microscope. The SEM has a large depth of
field, which allows a large amount of the sample to be in focus at one time. The SEM
also produces images of high resolution, which means that closely spaced features can be
examined at a high magnification. Preparation of the samples is relatively easy, since
most SEMs only require the sample to be conductive. The combination of higher
magnification, larger depth of focus, greater resolution, and ease of sample observation
makes the SEM one of the most heavily used instruments in materials research today.
The scanning electron microscope (SEM) has a very fine �probe� of electrons
with energies of up to 40 keV focused at the surface of the specimen in the microscope
and scanned across it in a �raster� or pattern of parallel lines. A number of phenomena
occur at the surface under electron impact: most important for scanning microscopy is the
emission of secondary electrons with energies of a few tens of eV and re-emission or
reflection of the high-energy backscattered electrons from the primary beam.
37
Backscattered Electron Imaging
In this particular research backscattered electron imaging was used to collect
images for microstructure characterization. The intensity of emission of both secondary
and backscattered electrons is very sensitive to the angle at which the electron beam
strikes the surface, that is, to topographical features on the specimen. When the electron
beam strikes the sample some of the electrons will interact with the nucleus of the atom.
The negatively-charged electron will be attracted to the positive nucleus but if the angle
is just right instead of being captured by the "gravitational pull" of the nucleus, it will
circle the nucleus and come back out of the sample without slowing down. These
electrons are called backscattered electrons because they come back out of the sample.
Because they are moving so fast, they travel in straight lines. In order to form an image
with BSE (backscattered electrons), a detector is placed in their path. When they hit the
detector a signal is produced which is used to form the image.
Different elements have different size of nuclei. As the size of the atom nucleus
increases, the number of BSE increases. Thus, BSE can be used to derive an image that
shows the different elements present in a sample. A typical image collected at SEM is
shown in Figure 3.6.
The magnification of this microscope is the ratio between the dimensions of the final
image display and the field scanned on the specimen. Usually, a magnification range of
SEM is between 10 to 200000X and the resolution (resolving power) is between 4 to 10
nm (40 - 100 Angstroms).
38
Figure 3.6: Typical SEM images for grain size study
There are several types of SEMs designed for specific purposes ranging from
routine morphological and microstructural studies of materials, to high-speed
compositional analyses.
Atomic Force Microscope
The AFM is also called the scanning force microscope (SFM) and was invented in
1986 by Binnig, Quate and Gerber. The AFM utilizes a sharp probe moving over the
surface of a sample in a raster scan. The probe is a tip on the end of a cantilever which
bends in response to the force between the tip and the sample. Figure 3.7 illustrates the
working of the AFM. As the cantilever flexes, the light from the laser is reflected onto the
39
split photo-diode. By measuring the difference signal (A-B), changes in the bending of
the cantilever can be measured.
Figure 3.7: AFM working principle
The interaction force between the tip and the sample can be found as the
cantilever obeys Hooke's Law for small displacements. The movement of the tip or
sample is performed by an extremely precise positioning device made from piezo-electric
ceramics, most often in the form of a tube scanner. The scanner is capable of sub-
angstrom resolution in x-, y- and z-directions. The z-axis is conventionally perpendicular
to the sample.
40
The AFM operates by measuring attractive or repulsive forces between a tip and
the sample. In its repulsive contact mode, the instrument lightly touches a tip at the end of
a cantilever to the sample. As a raster-scan drags the tip over the sample, some sort of
detection apparatus measures the vertical deflection of the cantilever, which indicates the
local sample height. Thus, in contact mode, the AFM measures hard-sphere repulsion
forces between the tip and the sample.
In non-contact mode, the tip does not touch the sample and the AFM derives
topographic images from measurements of attractive forces. AFMs can achieve a
resolution of 10 pm, and unlike electron microscopes, can image samples in air and under
liquids.
Feedback operation
The AFM can be operated in two principal modes:
• with feedback control
• without feedback control.
Once the electronic feedback is switched on, then the positioning piezo which is
moving the sample or tip up and down can respond to any changes in force which are
detected, and alter the tip-sample separation to restore the force to a predetermined value.
This mode of operation is known as constant force mode, and usually enables a fairly
faithful topographical image to be obtained.
When the feedback electronics are switched off, then the microscope is said to be
operating in constant height or deflection mode. This is particularly useful for imaging
41
very flat samples at high resolution. The following procedures were considered to obtain
high quality images for grain size study of all samples:
• The controller and the main computer were turned on and the SPM cockpit
software was launched.
• The close-contact EZ-Mode was selected and the corresponding configuration file
was loaded.
• The cantilever close-contact mode operation was installed in the AFM scanner
head.
• The cantilever tip was aligned so that the laser light was projected from the back
of the cantilever onto the photodetector.
• The AFM scanner was moved down and the surface was focused to obtain the
high quality image.
• The tip approach was carried out and a specific area on the sample was selected to
obtain better quality images with distinct grain boundaries.
• The probe was retracted from the sample.
• The above procedure was repeated with remaining samples.
• Grain size measurements were performed by the intersect method of Image Pro-
Plus software using Scanning Electron Microscope (SEM) and Atomic Force
Microscope (AFM) images (see Figure 3.8).
42
(a)
(b)
Figure 3.8: (a) Typical AFM images for grain size study; (b) Grain size and cavitations measurements
43
The measurements were performed, collecting data by moving the red lines from
top to bottom at 10 locations in the same images. Almost 20 grain intersects per line were
collected (see Figure 3.8 (b)). For each sample, the data were collected from 4/5 images
which gave 800-1000 measurements. Cavities density was collected using the threshold
method as shown in Figure 8.3 (b).
3.6 XRD Measurements
The X-ray diffraction experiment is mostly performed to characterize
crystallographic structure and chemical composition of the materials. In the present work,
the XRD experiment was used for phase identification of ceramic sintered at different
temperature, and calculation of coefficient of thermal expansion by obtaining lattice
parameters.
The experiments were performed using the AXS-Bruker D8 Discover
diffractometer with Cu Kα X-ray radiation source and Eurlean cradle. Parallel beam
optics was used to minimize errors associated with sample displacement and surface
roughness. The Anton Paar H-900 high temperature stage mounted on the XYZ-
positioning stage was used for measurements at temperatures up to 900ºC. The zero
point of detector was <0.004º, and sample height was controlled with accuracy of 0.005
mm. Ceramic samples were polished down to 0.5 mm thick tablet form with final grit
size 0.05 µm to avoid stress induced effects on XRD data.
Powder samples for high temperature measurements were deposited on Si(100)
substrate in the slurry form using colloidal silver paste. Silver matrix has been used for
44
temperature calibration. Diffraction patterns were collected at 25ºC and then between 100
and 800ºC in 100ºC intervals for CTE measurements and 10°C for phase transition
analysis. The heating rate was 10 ºC/min and the waiting time was 30 minutes before
XRD scans at every temperature for temperature stabilization. High temperature XRD
scans were performed in the interval 2θ of 28-66º with integration time 5s/point and step
0.01º. The waiting time for temperature stabilization before the XRD scan at each
temperature was 30 minutes.
3.6.1 Coefficient of Thermal Expansion (CTE)
When heat is given to a body, it normally expands. The coefficient of thermal
expansion is the relative change in a given dimension when the body is heated. Thermal
expansion of the crystalline phase means a change in the cell lattice parameters with the
temperature variation. Increase of the lattice parameter in the cubic phase with
temperature increase leads to shifting of diffraction peaks towards lower angles. An
example of this behavior is shown in Figure 3.9 for the P1600 powdered sample. To
determine coefficient of thermal expansion (CTE), cell parameters at every temperature
were obtained by Retvield refinement of at least 5 peaks in XRD data using Full-Proof
Software. The CTE is defined by the following formula:
dTdL
LCTE 1=
where, L is original length, dL is change in length and dT is temperature difference. This
demonstrates the linear expansion of solid material. The cubical or volume expansivities
45
of solids are three times the linear expansivity. The unit of linear CTE is /°C or
mm/mm/°C.
Figure 3.9: Temperature dependent shifting of XRD peaks positions
The expansivities of the majority of solid materials increase with increasing
temperature and can be fitted in three different ways:
1) Model-I:CTE- constant
obCTE =
By substitution and integration,
46
.*ln constTbL oT +=
2) Model-II: CTE-varying linearly
TbbCTE o 1+=
By substitution and integration,
.2
*ln 21 constTbTbL oT ++=
3) Model-III: CTE - varying qudratically
221 TbTbbCTE o ++=
By substitution and integration,
.32
*ln 3221 constTbTbTbL oT +++=
For all three models, coefficients bo, b1, b2 were obtained through curve fitting log
(LT) by linear, quadratic, and cubic functions respectively. An example of this fitting is
presented in Figure 3.10. Analysis of fitting residuals can help to select the appropriate
model for CTE (see Figure 3.11). By substituting value of coefficients and given
temperature, the CTE was calculated. In the case of the P1600 powdered sample
dependence Ln (LT) is not linear and the best fit is achieved by linear and quadratic
models. In cases like this, selection between linear and quadratic models is difficult
because residuals are similar. Despite the constant and the linear models, the quadratic
model is the better fitting (less residuals) model. The accuracy of coefficients of fitting
was used for selection of the appropriate model. In the case of the P1600 powdered
47
sample (see Figures 3.9, 3.10, 3.11) for the linear model b0 = (7.14±0.39)×10-6 and
b1=(4.61±0.53)×10-9, while for the quadratic model b0= (5.83±2.33)×10-6,
b1=(8.47±6.75)×10-9 and b2=-(2.68±4.67)×10-12. It is evident that statistical accuracy of
coefficients for the linear model is significantly better than for the quadratic model.
Therefore, the best model was decided based on analysis of residuals, overall standard
deviations and standard errors of parameters. To compare our data with the known value
of CTE of cubic zirconia (10.5×10-6 deg-1), we have used the results of a constant model.
300 400 500 600 700 800 900 1000 1100
1.628
1.629
1.630
1.631
1.632
1.633
1.634
1.635
1.636 Constant Linear Qudratic
Ln(L
T)
Temperature (K)
Figure 3.10: Constant, linear and quadratic fitting with temperature vs. log (LT) (P1600-powder)
48
300 400 500 600 700 800 900 1000 1100
-0.00015
-0.00010
-0.00005
0.00000
0.00005
0.00010
0.00015
Del
ta
Temperature, K
Constant Linear Quadratic
Residuals
Figure 3.11: Instrumentally set temperature vs. residuals (P1600-powder)
3.6.1.1 Temperature Calibration
It is known that the coefficient of thermal expansion is temperature dependent
parameters. The actual value of temperature is essential for the accurate calculation of
coefficient of thermal expansion. Based on literature reviews, the CTE at known
temperature of silver was taken as reference value for temperature calibration. By back
substitution and polynomial fitting using Microsoft excel and the Origin program, the
coefficient of polynomial equation was obtained for known temperature, and for each set
temperature the actual temperature was calculated as shown in Figure 3.12 and Table 3.3.
49
Figure 3.12: Polynomial fitting of calibrated temperature vs. instrumentally set temperature Table 3.2. Calculation of actual temperature from set temperature
The results show that the porosity level in ceramics sintered from Praxair and
DKKK powders has significant sinteribility difference. It was found that the DKKK
powder was significantly more active in the sintering process and required ~200ºC lower
sintering temperature to produce ceramics of the same level of porosity as in samples
sintered from Praxair powder (Figure 4.1). Amounts of porosity had been calculated
using theoretical densities 5.49 and 5.74 g/cm3 of ScCeZrO2 ceramics Praxair and DKKK
powder respectively which were estimated using XRD data and the actual chemical
composition of the powders. Density data showed that for ceramics produced from
59
DKKK powder, density increases rapidly from 1100 to 1400ºC and at higher
temperatures become constant (Table 4.4). Similarly, in the same temperature range,
grain size in DKKK ceramics increases significantly from 200 nm to 3 µm and stabilizes
at higher temperatures. Even at 1200°C samples produced from DKKK powder are dense
enough for polishing, grain size measurements and microhardness testing while samples
sintered from Praxair powder at 1400°C are still porous for microindentation. Analysis of
grain size with increase of sintering temperature allowed for the conclusion that DKKK
powder produces dense ceramics at temperatures higher than 1400ºC while Praxair
powder is still not fully sintered at 1600ºC. This conclusion was supported by analysis of
fracture surfaces. Figure 4.5 shows typical fracture surfaces of the samples sintered at
1600ºC from Praxair powders and DKKK powder. It is evident that DKKK powder
created almost
Figure 4.5: Fracture surfaces of samples P1600-(a) and J1600-(b)
60
fully dense ceramic while the sample produced from the Praxair powder had noticeable
amounts of porosity. It should be noticed that thermal etching effects on polished
surfaces of samples produced from Praxair powders and DKKK are different (see Figures
4.6-4.8).
Figure 4.6: Grain structure of samples produced from Praxair powder sintered at temperatures 1500ºC (left) and 1600ºC (right)
Figure 4.7: Grain structure of samples produced from DKKK powder sintered at temperatures 1400ºC (left), 1500ºC (middle) and 1600ºC (right)
61
Figure 4.8: SEM images of polished and low angle ion-milled surface of sample P1600 (a) and thermally etched surfaces of samples P1600 (b) and J1600 (c)
Praxair-based samples unusually exhibit significant increase of in-grain cavities
due to thermal etching with sintering temperature while the in-grain cavitations level in
DKKK-based samples remain constant. SEM analysis reveals that all Praxair-based
samples have significant non-uniformity. Even the densest sample, P1600, has highly
dense and very porous areas with well distinguished borders between them (see Figure
4.8(a)). The difference in overall density between samples P1400, P1500 and P1600
(Table 4.4) results predominantly from different ratios between dense and porous areas in
the samples. Grain structures in Figure 4.8(b) demonstrate a higher density area of
sample P1600 while average porosity of the sample is 4.4%. Contrary to samples
(Praxair), DKKK-based samples have high uniformity, and the density of in-grain
cavities remains constant for samples J1400, J1500 and J1600 (Table 4.4). One can
conclude that DKKK powder produces dense ceramics at temperatures higher than
1400ºC while ceramics produced from Praxair powder are still not fully sintered at
1600ºC (see Figures 4.9-4.10).
62
1100 1200 1300 1400 1500 1600
2.0
2.5
3.0
3.5
4.0
4.5
5.0
5.5
6.0
Praxair Japan
Den
sity
(g/c
cm)
Temperature (C)
Figure 4.9: The dependence of density on sintering temperatures
1100 1200 1300 1400 1500 16000
5
10
15
20
25
30
35
40
45
Porous
Poro
sity
Lev
el (%
)
S intering Temperature (0C)
Praxair (I) DKKK (II)
Figure 4.10: Porosity level vs. sintering temperature
63
The difference in sintering activity of Praxair and DKKK powders is a result of
differences in powder morphology, chemical composition and impurities. Praxair powder
consists of relatively large 1-2 µm pre-sintered agglomerates of grains while DKKK
powder has a significantly higher surface area due to relatively uniform nanograins (see
Figure 4.1) which favors lower sintering temperature for DKKK powder. Also Praxair
and DKKK powders have a slightly different chemical composition that could have an
additional effect on the sintering process.
4.2.2. Microstructure and Grain Size
The microstructure properties were characterized by collecting images at SEM
and AFM of sintered ceramics. The high quality images were obtained with distinct grain
boundary and grain size measurement was performed as described in Chapter 4 using
Image Pro-Plus. It is observed from the experimental data that the sintering temperature
has a significant effect on the grain size of the ceramic and its microstructure is shown in
Figures 4.11, 4.12, and 4.13. Results of grain size and cavities are listed in Table 4.5.
Figure 4.30: (a) Effect of sintering temperature on microhardness and fracture toughness at the load 1000g and; (b) Micrograph of typical Vickers indents and cracks at different loads (thermally etched J1600)
Tables 4.7-4.8 contain the hardness and fracture toughness values for each load from
1000g-100g. The weighted value for hardness and fracture toughness was also calculated.
The results obtained for DKKK based and Praxair-based ceramics are listed in Table 4.7
The following conclusions can be made based on the results of the research work
outlined:
• Phase and chemical composition, morphology, and grain size of the 10 mol%
Sc2O3 � 1 mol% CeO2 - ZrO2 ceramic powders produced by Praxair Specialty
Ceramics, USA and by DKKK, Japan has been characterized.
• The sinterability of the powders has been studied by sintering at 6 different
temperatures (1100ºC -1600ºC).
• The sintering rate of the ceramic powder manufactured by DKKK is much higher
than the powder manufactured by Praxair. The DKKK powder was able to reach
its lowest porosity level at lower temperatures and shorter sintering times than the
Praxair powder.
• DKKK-based ceramics have phase instability in the region 300-500ºC.
Transformation from c- to b-phase at these temperatures is slow process and it is
required to use high heating and cooling rates to avoid formation of significant
amounts of b-phase.
• The ionic conductivity of selected samples is reported to be ~0.088 S/cm for
Praxair Ceramics and ~0.156 S/cm for DKKK ceramics at 800ºC.
90
• Coefficients of thermal expansion in the range RT-800ºC of both types of
ceramics are in the range 10.5 � 10.6×10-6 grad-1 and very close to CTE of cubic
zirconia. CTE of DKKK samples have less linear dependency on temperature than
Praxair samples.
• Young�s modulus and hardness of DKKK samples are higher than Praxair
samples.
• High temperature fracture toughness measurements confirm that these ceramics
are very brittle, especially at elevated temperatures. The highest K1C was 2.5
MPa×m.5 at 300ºC for DKKK ceramics.
5.2 Future Work
The following are recommendations for future work:
• Study of mechanical and microstructure properties of J1300 samples by changing
the dwell time and heating and cooling rate.
• Study of mechanical and microstructure properties of rhombohedral phases of
samples J1300-J1600
91
REFERENCES
[1] Rayne O�Hayre, Suk-Won Cha, Whitney Colella, Fritz B. Prinz, Fuel Cell Fundamentals, Chapter 1, Published by John Wiley & Sons, New York, (2006). [2] María Mercedes González Cuenca, Novel anode materials For Solid Oxide Fuel Cells, Chapter1, Twente University Press, P.O. Box 217, 7500 AE Enschede, the Netherlands, www.tup.utwente.nl. [3] Sharon Thomas and Marcia Zalbowitz at Los Alamos National Laboratory in Los Alamos, New Mexico, Fuel Cells Green Powder http://www.scied.science.doe.gov/nmsb/hydrogen/Guide%20to%20Fuel%20Cells. pdf. [4] Materials Synthesis and Processing (IEF-1)
http://www.fz-juelich.de/ief/ief-1/index.php?index=70 (April 5, 2007). [6] A. J. Moulson and J.M. Herbert, Electroceramics, Materials, properties. Applications, chapter 4, Chapman & Hall, London, UK 1990. [7] R. Chiba, T. Ishii, F. Yoshimura, "Temperature dependence of ionic conductivity in (1-x)ZrO2-(x-y)Sc2O3-yYb2O3 electrolyte material", Solid State Ionics, 91(3,4), 1996, 249-256. [8] Z. Lei, Q. Zhu, "Low temperature processing of dense nanocrystalline scandia- doped zirconia (ScSZ) ceramics", Solid State Ionics, 176(37-38), 2005, 2791-2797 (2005) [9] I. Kosacki, H.U. Anderson, Y. Mizutani, K. Ukai, "Nonstoichiometry and electrical transport in Sc-doped zirconia.", Solid State Ionics, 152-153, 2002, 431- 438. [10] M. Yashima, M. Kakihana, M. Yoshimura, "Metastable-stable phase diagrams in the zirconia-containing systems utilized in solid-oxide fuel cell application", Solid State Ionics, 86-88 (Pt. 2), 1996, 1131-1149. [11] H. Fujimori, M. Yashima, M. Kakihana, M. Yoshimura, "The b-cubic phase transition of scandia-doped zirconia solid solution: Calorimetry, x-ray diffraction, and Raman scattering", Journal of Applied Physics, 91(10, Pt. 1), 2002, 6493- 6498.
92
[12] C. Haering, A. Roosen, H. Schichl, M. Schnoeller, "Degradation of the electrical conductivity in stabilized zirconia system. Part II: Scandia-stabilized zirconia", Solid State Ionics, 176(3-4), 2005, 261-268. [13] O. Yamamoto, Y. Arati, Y. Takeda, N. Imanishi, Y. Mizutani, M. Kawai, Y. Nakamura, "Electrical conductivity of stabilized zirconia with ytterbia and scandia", Solid State Ionics, 79, 1995, 137-142. [14] M. Hirano, T. Oda, K. Ukai, Y. Mizutani, "Suppression of rhombohedral-phase appearance and low-temperature sintering of scandia-doped cubic-zirconia", Journal of the American Ceramic Society, 85(5), 1336-1338, 2002. [15] R. Chiba, F. Yoshimura, J. Yamaki, T. Ishii, T. Yonezawa, K. Endou, "Ionic conductivity and morphology in Sc2O3 and Al2O3 doped ZrO2 films prepared by the sol-gel method", Solid State Ionics, 104(3,4), 1997, 259-266. [16] I. Kosacki, "Nanoscaled oxide thin films for energy conversion.", NATO Science Series, II: Mathematics, Physics and Chemistry, 202 (Fuel Cell Technologies), 2005, 395-416. [17] J. Rodriguez-Carvajal, "Recent advances in magnetic structure determination by neutron powder diffraction", Physica B: Condensed Matter (Amsterdam, Netherlands), 192 (1-2), 1993, 55-69. [18] The A to Z of Materials (Azom.com) http://www.azom.com/details.asp?ArticleID=2975#_Electrolyte_Materials_For_S olid Oxi. ( April 7, 2007). [19] Mettler Toledo GmbH, Laboratory & weighing Technologies, Ch- 8606 Greifensee, Switzerland, Phone +41-1-944 22 11, Model: AX105 Delta Range(R). Density Determination of Solid. [20] http://www.answers.com/density&r=67, Answers.com ( April 9, 2007). [21] D. Tabor The Hardness of Metals, Oxford: London (1951). [22] W.C.Oliver and G.M.Pharr, An Improved Technique for Determining Hardness and Elastic Modulus Using Load and Displacement Sensing Indentation Experiments, J. Mater. Res. Vol 7 (No. 6), 1992, 1564-1583. [23] G.M.Pharr and W.C.Oliver, Measurement of Thin Film Mechanical Properties Using Nanoindentation, MRS Bull., Vol 17, 1992, 28-33.
93
[24] G.M.Pharr, Measurement of Mechanical Properties by ultra-low Load Indentation, Mater. Sci. Eng. A, Vol 253, 1998, 151-159. [25] W.D.Nix, Mechanical Properties of Thin Films, Metall. Trans. A, Vol 20, 1989, 2217-2245. [26] M.F.Doerner and W.D.Nix, A Method for Interpreting the Data from Depth- Sensing Indentation Instruments, J.Mater. Res., Vol 1, 1986, 601-609. [27] J.B.Pethica, R. Hutchings and W.C.Oliver, Hardness Measurements at Penetration Depths as Small as 20nm, Philos. Mag. A, Vol 48(No. 4), 1983, 593-606. [28] W.C.Oliver, Progress in the Development of a Mechanical Properties Microprobe, MRS Bull., Vol 11 (No. 5), 1986, 15-19. [29] Yeon-Gil Jung and Brian R Lawn, Evaluation of Elastic Modulus and Hardness of Thin Films by Nanoindentation, J. Mater. Res., Vol 19 (No. 10), 3076-3080 [30] J.L.Hay and G.M.Pharr, Instrumented Indentation Testing, ASM Handbook, Vol 8, Mechanical Testing and Evaluation (2000). [31] J.W.Harding and I.N.Sneddon, Proc. Cambridge Philo. Soc., Vol 41, 1945, 12. [32] J.S.Field and M.V.Swain, A Simple Predictive Model for Spherical Indentation, J.Mater. Res., Vol 8 (No. 2), 1993, 297-306. [33] J.S.Field and M.V.Swain, Determining the Mechanical Properties of Small Volumes of Material from Submicron Spherical Indentations, J. Mater. Res., Vol 10 (No. 1), 1995, 101-112. [34] M.V.Swain, Mechanical Property Characterization of Small Volumes of Brittle Materials with Spherical Tipped Indenters, Mater. Sci. Eng. A, Vol 253, 1998, 160-166. [35] J.L.Loubet, B.N.Lucas and W.C.Oliver, Some Measurements of Viscoelastic Properties with the Help of Nanoindentation, NIST Special Publication 896: International Workshop on Instrumented Indentation, 1995, 31-34. [36] B.N.Lucas, C.T.Rosenmayer and W.C.Oliver, Mechanical Characterization of Sub-Micron Polytetraflouroethylene (PTFE) Thin Films, in Thin Films-Stresses and Mechanical Properties VII, MRS Symposium Proc., Vol 505, Materials Research Society, 1998, 97-102.
94
[37] B.N.Lucas, W.C.Oliver, J.L.Loubet, and G.M.Pharr, Understanding Time Dependent Deformation During Indentation Testing, in Thin Films-Stresses and Mechanical Properties VI, MRS Symposium Proc., Vol 436, Materials Research Society, 1997, 233-238. [38] G.M.Pharr, D.S.Harding and W.C.Oliver, Measurement of Fracture Toughness in Thin Films and Small Volumes Using Nanoindentation Methods, Mechanical Properties and Deformation Behavior of Materials Having Ultra-Fine Microstructures, Kluwer Academic Publishers, 1993, 449-461. [39] D.S.Harding, W.C.Oliver and G.M.Pharr, Cracking During Nanoindentation and Its Use in the Measurement of Fracture Toughness, in Thin Films-Stresses and Mechanical Properties V, MRS Symposium Proc., Vol 356, Materials Research Society, 1995, 663-668. [40] D.Tabor, M.J.Adams, S.K.Biswas and B.J.Briscoe, Solid-Solid Interactions, Imperial College Press (1996). [41] I.N.Sneddon, The Relation Between Load and Penetration in the Axisymmetric Boussinesq Problem for a Punch of Arbitrary Profile, Intl. J. Eng. Sci., Vol 3, 1965, 47-56. [42] A.Bolshakov and G.M.Pharr, Influences of Pile-Up on the Measurements of Mechanical Properties by Load and Depth Sensing Indentation Techniques, J. Mater. Res., Vol 13, 1998, 1049-1058. [43] G.M.Pharr, W.C.Oliver and F.R.Brotzen, On the Generality of the Relationship among Contact Stiffness, Contact Area, and Elastic Modulus, J. Mater. Res., Vol 7 (No. 3), 1992, 613-617. [44] R.B.King, Elastic Analysis of Some Punch Problems for a Layered Medium, Int., J. Solids Struct., Vol 23, 1987, 1657-1664. [45] B.C.Hendrix, The Use of Shape Correction Factors for Elastic Indentation Measurements, J. Mater. Res., Vol 10 (No. 2), 1995, 255-257. [46] A.Bolshakov and G.M.Pharr, Inaccuracies in Sneddon�s Solution for Elastic Indentation by a Rigid Cone and Their Implications for Nanoindentation Data Analysis, Thin Films-Stresses and Mechanical Properties VI, MRS Symposium Proc., Vol 436, Materials Research Society, 1997, 189-194. [47] J.C.Hay, A. Bolshakov, and G.M.Pharr, A Critical Examination of the Fundamental Relations in the Analysis of Nanoindentation Data, J. Mater. Res., Vol 14 (No. 6), 1999, 2296-2305.
95
[48] J.C.Hay, A.Bolshakov and G.M.Pharr, Applicability of Sneddon Relationships to the Real Case of a Rigid Cone Penetrating an Infinite Half Space, in Fundamentals of Nanoindentation and Nanotribology, MRS Symposium Proc., Vol 522, Materials Research Society, 1998, 263-368. [49] J.C.Hay and G.M.Pharr, Experimental Investigations of the Sneddon Solution and an Improved Solution for the Analysis of Nanoindentation Data, in Fundamentals of Nanoindentation and Nanotribology, MRS Symposium Proc., Vol 522, Materials Research Society, 1998, 39-44. [50] I .Manika, J. Maniks, "Size effects in micro- and nanoscale indentation", Acta
Materialia, 54(8), 2006, 2049-2056 [51] W.D. Nix, H. Gao, "Indentation size effects in crystalline materials: a law for
strain gradient plasticity", Journal of the Mechanics and Physics of Solids, 46(3), 1998, 411-425
[52] MRS Bulletin, Publication of the materials research society, Volume 30, Number- 8, ISSN: 0883-7694 CODEN: MRSBEA, (2005).