~NASb-CR-182800-V01-1) HAPBIA-RICH HIXED b188- 221 97 

OXXDE CERAWZCS OF THE SYSTEM HfO2-Zr02-Ti02 

FOR HEATERS AND HEAT EXCXANGERS IN 

ELECTROTHERMAL THRUSTERS: THE EFFECTS OF Unclas 

TITANIA ON SELECTED ELECTRICAL ABD G3/27 0141230 
HAFNIA-RICH MIXED OXIDE CERAMICS OF THE SYSTEM 

Hf02-Zr02-Ti02 FOR HEATERS AND 

HEAT EXCHANGERS IN ELECTROTHERMAL THRUSTERS 

Final Report, Vol. I 

NAG 3-495 

The Effects of Titania on Selected Electrical and Mechanical 

Properties 

of Hafnia-Rich Mixes Oxides in the System 

Haf n ia-Zi rcon ia-Titan ia 

P. R. Staszak, G. P. Wirtz, M. Berg and S. D. Brown 

Department of Materials Science and Engineering 

204 Ceramics Building 

105 So. Goodwin Avenue 

University of Illinois at Urbana-Champaign 

Urbana, IL 61801 

May 1988 
THE EFFECTS OF TITANIA ON SELECTED 

ELECTRICAL AND MECHANICAL PROPERTIES 

OF HAFNIA-RICH MIXED OXIDES 

IN THE SYSTEM HAFNIA-ZIRCONIA-TITANIA 

Paul Russell Staszak, Ph.D. 

Department of Ceramic Engineering 

University of Illinois at Urbana-Champaign, 1988 

ABSTRACT 

A study of the effects of titania on selected properties of hafnia-rich mixed 

oxides in the system hafnia-zirconia-titania (HZT) was made in the region 

5 to 20 mol% titania. The studied properties included electrical conductivity, thermal 

expansion, and fracture strength and toughness. The electrical conductivity was 

measured using a two-point probe with 10 kHz ac current on bars over the 

temperature range 400 to 17OO0C and oxygen partial pressure range 10-3' to 21 kPa 

(10-33 to 0.21 atm). The thermal expansion was measured using a dilatometer from 

room temperature to 17OOOC. Fracture strengths were determined using a four-point 

bend fixture on nonindented bars at room temperature and on triple Vickers indented 

bars at room temperature, 1000°C, and 14OOOC. Fracture toughness (Klc) values 

were also determined from the indented bars. 

The effects of titania on the properties were studied for the reduced state as well 

as the oxidized state of the sintered mixed oxides. X-ray analyses showed that the 

materials were not always single phase. The oxidized compositions went' from being 

monoclinic solid solutiops at low titania additions to having three phases (two 

monoclinic and a titanate phase) at high additions of titania. The reduced 

compositions showed an increasing cubic phase presence mixed with the monoclinic 

phase as titania was added. The grains of all the samples were nominally 1 pm. 

Microcracking regions varying from 16 pm to 3000 pm were also observed in the 
microstructures. The wide variability in phases and microcracking regions accounted 

for the majority of the observed property trends. 

The electrical conductivity increased with temperature to approximately 

10-1 mhos/cm at 170OOC for all compositions. For low temperatures at high oxygen 

partial pressures, higher concentrations of titania lowered the conductivity suggesting 

that the titanium ions go into interstitial positions thereby reducing oxygen ion 

conductivity. For all temperatures at low oxygen partial pressures, the materials 

became highly anion deficient leading to higher conductivities suggesting that Ti4+ is 

being converted to Tis+ and Ti*+ thereby enhancing electronic conductivity. 

The thermal expansion coefficient decreased with increasing titania 

(6.2 x 10-6/OC at 5 mol% to 5.0 x 10-V°C at 20 mol% titania) as did the monoclinic to 

tetragonal transformation temperature. A hysteresis between the heating and cooling 

curves in the 15 and 20 mol% compositions suggested extensive microcracking. The 

low thermal expansion coefficient of the second phase titanate partially accounted for 

the decrease in thermal expansion in these higher titania compositions. 

The fracture strength of the oxidized bars tended to decrease with the addition 

of titania owing to the presence of the second phase titanate. The fracture strengths of 

the reduced bars exhibited a minimum corresponding to a two-phase region of 

monoclinic and cubic phases. When the second phases were suppressed, the titania 

tended to increase the fracture strength slightly in both the oxidized and reduced 

states. The fracture toughness followed similar trends. 
TABLE OF CONTENTS 

1 . 

2 . 

3 . 

4 . 

5 . 

6 . 

INlRODUCTlON ...................................................................................................................... 1 

LITERATURE REVIEW ........................................................................................................... 3 

2.1 Structure and Systems of Hafnia, Zirconia. and Titania ................................... 3 

2.2 Property Studies of Hafnia. Zirconia. and Titania. ............................................ 15 

EXPERIMENTAL PROCEDURES ...................................................................................... 52 

3.1 Compositions and Batching ................................................................................. 52 

3.2 Sample Preparation ............................................................................................... 53 

3.3 Property Measurements ........................................................................................ 55 

3.4 Analyses ................................................................................................................... 60 

RESULTS AND DISCUSSION ........................................................................................... 62 

4.1 Powder Analyses .................................................................................................... 62 

4.2 Densities. Length Changes. and Mass Changes ............................................. 62 

4.3 Microstructure Analyses ........................................................................................ 71 

4.4 Phase Analyses ...................................................................................................... 80 

4.5 Electrical Conductrvity ...... : .................................................................................... 85 

4.6 Thermal Expansion ................................................................................................ 96 

4.7 Fracture Strength and Toughness ..................................................................... 104 

.. 

CONCLUSIONS ................................................................................................................. 126 

SUGGESTIONS FOR FUTURE WORK .......................................................................... 128 

APPENDICES .......................................................................................................................... 129 

REFERENCES ......................................................................................................................... 132 

VITA ............................................................................................................................................ 141 

c 
1 

1. INTRODUCTION 

Hafnia-rich mixed oxides in the system hafnia-zirconia-titania (HZT) possess 

many properties which are ideal for a number of applications including heat 

exchangers, electrothermal thrusters, and fuel cells. They have high melting 

temperatures (>22OO0C), moderate (< 4 x lO-S/OC) to zero thermal expansion 

coefficients, and high thermal shock resistance.1-5 Owing to their similarity to hafnia 

and zirconia (they are isomorphous), they possess the potential to exhibit many of the 

property benefits of their parent oxides. Among these benefits are good hightemperature electrical conductivity and strength. 

The high electrical conductivity of zirconia (>10-2 mhos/cm at temperatures 

above 800OC) and its remarkable strengths (>300 MPa) have led to many studies of 

the electrical and mechanical properties of materials with the fluorite structure. 

Included among these studies have been investigations of the properties of hafnia, 

zirconia, and their binary system. In addition, the effects on the properties of various 

additives to these materials have been examined in an effort to improve the material 

characteristics. Included among the list of additives that have been investigated is 

titania. 

From studies of the zirconia-hafnia system, the material properties have been 

shown to follow the rule of mixtures (Le., the properties of the mixtures vary 

proportionally to the amounts of the pure phases present). Such a correlation would 

be expected since hafnia and zirconia form a complete solid solution with one another. 

From studies of the zirconia-titania and the hafnia-titania systems, solid solutions have 

been shown to result with additions of several mole percent titania. Beyond the 

solubility limit, a titanate phase forms. The effect of titania in these systems has been 

to increase the electrical conductivity, lower the thermal expansion, and increase the 

strength. In the regions of the phase diagram where the titanate phase is present, the 

strength of the material has been observed to decrease with titania additions. 

b 
2 

Although the HZT system has been investigated by several authorsl-5, a 

thorough study of the entire system has never been made. Solid solutions have been 

shown to exist in the hafnia-rich region of the ternary system: however, the amount of 

titania as well as the processing parameters govern whether the material remains a 

solid solution or becomes a mixture of a titanate phase in a monoclinic solid solution. 

The amount of titania also affects whether the system will duplicate the properties of its 

parent oxides. The exact effect of titania in the HZT system and the effect of titania on 

the electrical and mechanical properties of the system are two areas this research 

examines. 
3 

2. LITERATURE REVIEW 

This review will be divided into two sections. The first will examine the structure 

and systems of hafnia, zirconia, and titania. The second will survey selected electrical 

and mechanical properties in the systems; namely, electrical conductivity, thermal 

expansion, and fracture strength and toughness. From the review, certain insights and 

knowledge of the effects that titania has on the structure and properties in the HZT 

system will be gained. 

2.1 Structure and Systems of Hafnia, Zirconia, and Titania 

2.1 .l Structure 

Three crystalline structures are encountered in the systems of hafnia, zirconia, 

and titania. They are the fluorite, the rutile, and the titanate. The more important of the 

three in the hafnia-rich region of the HZT system are the fluorite and the titanate 

structures. 

2.1 . 1 .1 Fluorite 

Both hafnia and zirconia have distorted fluorite crystal structures whereas titania 

has the rutile structure. From considerations of ionic radii and Pauling's rules, 

however, all three oxides should have the rutile structure.6 Most AX2 compounds with 

high cation to anion radius ratios (approaching 1.0) have the fluorite structure, while 

those with intermediate ratios (0.43 to 0.73) have the rutile structure and those with low 

ratios (< 0.4), one of the silica structures.7 The difference in the structures lies in the 

coordination of anions to cations. For the fluorite structure, there is eightfold 

coordination. For the rutile structure the coordination is sixfold, and for the silica 

structures, fourfold. Cation to anion radius ratios for typical fluorite structures are 

CeO2 0.68, U02 0.70, Tho2 0.75, CaF2 0.84, HgF2 0.88, SrF2 0.94, etc..7 Table 1 

shows the ionic radii for hafnia, zirconia, and titania along with the cation to anion 

+ 
4 

I 

I radius ratios, coordination numbers, and crystal structures. Even though ionic radius 

ratios suggest the opposite, the degree of covalency that exists in Hf02 and Zr02 

apparently causes the fluorite structure to be more stable than the rutile.* Covalency 

in the structure was suspected to exist due to the sevenfold coordination of anions 

I about cations in the low temperature monoclinic phase. Although the degree of 

covalency in the monoclinic phase of the ZrO2 fluorite structure was not calculated, the 

I 

I 

covalency of the high temperature forms was found to be 28 and 22%, respectively, for 

the tetragonal and cubic phases? The monoclinic form should have a similar degree. 

Table 1. Ionic Radii, Radius Ratios, Coordination Numbers, and Crystal Structures 

for Hf02,ZrO2, and Ti02. 

~ 

Ion Ionic Radii (Ref. 9) M4+/@- Crystal 

8CN 7CN 6CN 8CN 7CN 6CN Structure 

HP+ .83 .76 -71 .58 .54 51 Fluorite 

Z1Q+ .84 .78 .72 .59 .56 .51 Fluorite 

Ti4+ .74 - .605 .52 .43 Rutile 

Ti3+ - - .810 

Ti2+ - - 1 .oo 

02- 1.42 - 1.40 

As mentioned, three polymorphs of fluorite exist for HfO2 and ZrO2. The low 

temperature phase is monoclinic. The higher temperature phases are tetragonal and 

cubic. Figure 1 shows the three polymorphic phases of the fluorite structure. 

The monoclinic phase is stable at all temperatures below 1650OC for HfO2 and 

1170OC for ZrO2. The boordination of anions about cations is sevenfold having a 

range of bond lengths and angles. Triangularly coordinated and distorted tetrahedrally 

coordinated O-M4+ layers with different thicknesses result.8 The lattice sites of 

monoclinic ZrO2 are shown in Figure 2. The movement off normal fluorite positions 

can be seen. HfO2 (being isomorphous with ZrO2) has a similar lattice configuration. 
5 

Figure 1. Polymorphic phases of the fluorite structure (a) cubic, (b) tetragonal, and 

(c) monoclinic. Solid circles: Z$+. Open circles: 02- (Ref. 7). 

The tetragonal phase is stable at temperatures between 1650°C and 2700°C 

for HfO2 and 1170°C and 2370°C for ZrO2. The monoclinic to tetragonal 

transformation has a thermal hysteresis of approximately 200°C between heating and 

cooling. The coordination of anions to cations is eightfold with two sets of four 

oxygens at slightly different bond lengths8 Figure 3 shows the tetragonal lattice sites. 

The cubic phase is stable above 2700°C for HfO2 and 2370°C for Zr02 to their 

respective melting temperatures, 2900°C and 2680°C. Eightfold coordination exists 

having equidistant cation*anion bond lengths. The lattice sites of the cubic fluorite 

structure are shown in Figure 4. The slight distortion from the ideal fluorite structure is 

illustrated. 

The lattice parameters of the three polymorphs of HfO2 and Zr02 are given in 

Table 2. Also included are the cell volumes and theoretical densities, 
6 

2- 

Figure 2. Lattice positions of the monoclinic cell of ZQ. Solid circles: Z++. Open 

circles: 02-. Broken lines become tetragonal cell after transformation 

(Ref. 10). 
7 

Figure 3. Lattice positions of the tetragonal cell of ZrO2. Solid circles: Z++. Open 

circles: 02- (Ref. 10). 

s. 

0 Cations 0 Oxygen 

Figure 4. Distortion of cubic lattice sites of ZrOp from ideal fluorite crystal structure 

(Ref. 8). 
I 

I 

I 8 

Table 2. Summary of Lattice Parameters for Hf02, ZQ, Ti02, HfTiO,, and ZrTiO,. 

~ ~ 

Phase a b C 0 v p Measurement Ref. 

1 (A) (A) (A) (degrees) (As) (g/cm3) Temperature I 

I Hf02 m 

t 

1 C 

m 

t 

Ti02 t 

HfTi04 o 

ZrTiO, o 

I 

I z02 

C 

5.116 5.172 5.295 

5.140 5.25 

5.145 5.208 5.311 

5.166 5.298 

5.070 

4.5933 2.9592 

4.750 5.043 5.573 

4.806 5.032 5.447 

- 

99.18" 138.3 

138.7 

99.23" 140.5 

141.4 

130.3 

62.43 

133.5 

131.7 

- 

10.1 1 

10.08 

5.826 

5.790 

6.283 

4.250 

7.223 

5.1 23 

- 

Frr 

1865°C 

RT 

1400°C 

Frr 

RT 

FIT 

RT 

- 

- 

11 

12 

11 

13 

14 

15 

16 

17 

- 

2.1.1.2 Titanate 

In the study of the hafnia-rich portion of the HZT system, the fluorite structure 

plays a main role since it tends to be the dominating phase structure. The rutile 

structure, on the other hand, is not encountered. Therefore, an elaboration of the 

structure is not warranted in this review. A structure that is encountered, however, is 

that of the titanate phase. Beyond the solubility limit of titania in monoclinic solid 

solutions of hafnia and zirconia, hafnium titanate and zirconium titanate form as 

second phases. The titanate phases will presently be discussed for completeness of 

the structure review. 

Figure 5 shows the titanate crystal structure. The unit cell is orthorhombic with 

sixfold coordination of anions about cations. Lattice parameters and unit cell volumes 

for HfTi04 and ZrTi04 are given in Table 2. The Hffi04 solid solution range is between 

35 and 55 mol% titania At room temperature and is stable to 1980°C.18 The ZrTi04 

solid solution range is 42 to 50 mol% titania and extends to 19OO0C.19 

2.1.2 Systems 

Several investigations of the three binary systems of the HZT ternary have been 
9 

I 

I #.SO I 

f I cQ724 o~-~~~ 

c I 

I 

l 

I 

I 

I i 

I I 

I iA.776 -0.224 I 

Figure 5. Lattice positions in the titanate crystal structure. Open circles: 02-. Solid 

circles: Zr4+ (Ref. 20). 

made. A thorough investigation of the ternary, however, has never been done. 

Nonetheless, a good deal of insight into the ternary can be gained from information on 

the solubility limits, two-phase regions, and phase transformations of the binaries. 

2.1.2.1 HfO2-ZrO2 1 

Hafnia and zirconia form a complete solid solution with one another.12s21-22 The I 

zirconium ions substitute homogeneously and directly for the hafnium ions in the 

fluorite structure? As zirconia is added, the lattice parameters increase smoothly,22 

the densities decrease,21 and the monoclinic to tetragonal, tetragonal to cubic, and 

cubic to liquid transformation temperatures Figure 6 is the binary 

Hf02-Zr02 phase equilibrium diagram constructed by Ruh et. al.? The triangles in the 

diagram represent the experimental liquid to cubic transformation temperature. (The 
10 

800 ' 

HfO*(WtX) 

to 80 sp KH) IO 20 ?o 40 50 60 

3200 

3000 Liquid 

22200 

a frtrogonat 

Figure 6. Phase equilibrium diagram for the Hf0,-ZrO, system (Ref. 22). 

solidus and liquidus temperatures as well as the tetragonal to cubic inversions are 

estimated.) The open circles represent the monoclinic to tetragonal heating inversions 

while the X's represent the tetragonal to monoclinic cooling inversions. 

2.1.2.2 HfOrTiO2 

The solubility limit of TiO, in HfO, is 10 to 13 mol% at 160OOC.183 Beyond this 

limit, HfTi04 in a monoclinic solid solution occurs up to the HtTiO, solid solution region 
11 

which begins at 30 mol%.ls The most recent phase equilibrium diagram (by Coutures 

and Couturesls) for the Hf02-Ti02 system is shown in Figure 7 along with the older 

Ruh et. al.24 diagram. (The results of Coutures and Coutures agree with the findings 

of Schevchenko et. al..25) Coutures and Coutures explain the discrepancy in the 

diagrams to be the qualitative rather than quantitative determination of the HfTi04 

melting temperature in the Ruh et. al. work. 

Figure 7. Phase equilibrium diagrams for the Hf02-Ti02 system (a) (Ref. 18) 

and (b) (Ref. 24). 
12 

The phase boundaries of the equilibrium diagram of Coutures and Coutures 

were determined by examining both the transformation temperatures and the changes 

in lattice spacings of the titanate and monoclinic phases. Compositional limits of the 

phases were determined knowing: (1) The interplanar spacings of the monoclinic 

HfO, lattice decrease with TiO, additions in the solid solution region? (2) Beyond the 

solubility limit (in the two-phase region), the spacings become constant with increasing 

Ti02 as the titanate forms. (3) The titanate lattice parameters are constant in the twophase regions and decrease with Ti02 in the titanate solid solution region? Using 

lattice parameter measurements and phase transformation temperatures, Coutures 

and Coutures constructed the Hf02-Ti02 phase diagram accordingly. The phase 

diagram shows that the the monoclinic to tetragonal and tetragonal to cubic inversion 

temperatures decrease with KO2 concentration and become constant in the two-phase 

region. 

Concluding on the Hf02-Ti02 system, both Ruh et. al. and Coutures and 

Coutures observed that under certain thermal treatments, a second monoclinic phase 

formed in addition to the normal monoclinic and orthorhombic phases in the 

0 to 50 mol% Ti02 range. For samples annealed at 1600OC and above, traces of the 

second monoclinic phase appeared at 15 mol% TO, and became the only monoclinic 

phase by 30 mol%. For samples annealed at 1300°C, the second monoclinic phase 

was not observed. Coutures and Coutures suggest that the lower volume second 

monoclinic phase has a higher Ti02 concentration and that it decomposes to 

monoclinic HfO, and orthorhombic HfTi04. 

b 

2.1.2.3 zro2-Ti0, 

The solubility limit of KOe in ZrO, is about 10 to 15 mol%?6 Beyond this limit, 

ZrTi04 in a monoclinic solid solution results. Figure 8 shows the Zr02-Ti02 phase 

diagrams of Brown and Duwez,19 Coughanour et. al.17 and Noguchi and Mizuno.26 
13 

0 

a 

a 

(b) (a 

Figure 8. Phase equilibrium diagrams for Zr02-Ti02 (a) (Ref. 19), (b) (Ref. 26), 

and (c) (Ref. 17). 
14 

Many anomalies exist between the diagrams, probably as a result of the 

advancements in zirconia studies between 1954 and 1968. A comparison with the 

HfO,-TiO2 phase diagrams suggests that the phase equilibrium in the Zr02-Ti02 

system is best depicted by Noguchi and Mizuno. The larger Noguchi and Mizuno TiO, 

solubility limit is probably more reliable than the lower value predicted by Brown and 

Duwez since the latter was based on less than accurate lattice parameter calculations. 

The tetragonal to cubic transformation in the Noguchi and Mizuno diagram probably 

needs to be redrawn such that the transformation temperature increases with 

increasing TO2 (similar to the Hf02-Ti02 system). In addition, the diagrams show that 

a decrease in the monoclinic to tetragonal transformation temperature with increasing 

titania occurs and a ZrTi04 solid solution range exists. Unlike in the Hf02-Ti02 studies, 

none of the studies of the Zr02-Ti02 system showed a second monoclinic phase. 

2.1.2.4 Hf02-ZQ-Ti02 

The HZT ternary has been investigated by a number of authors;l-s however, the 

overwhelming task of producing a complete phase diagram has never been 

attempted. The primary difficulty in attempting such a task from literary information 

arises from the failure of most authors to specify the phases and compositions that they 

studied. Simpsonl investigated the range of compositions 60 to 100 mol% Hf02, 

0 to 13 mol% ZrO, and 0 to 64 mol% Ti02, generally referring to all compositions as 

hafnium titanates. (No phase analyses were performed.) Skaggss in summarizing the 

HZT studies of Holcombe et. al.27 (compositions in the range 16 to 47 mol% Hf02, 

16 to 47 mol% Zr02 and37 mol% Ti02), also recognized the failure of the authors to 

state the exact phases of the materials they were examining. Monoclinic solid 

solutions of 60 to 85 mol% Hf02, 10 to 30 mol% ZQ, and 10 to 30 mol% Ti02 were 

reported by Brown et. al.,4 however, only a select few compositions were specified 

exactly. Complicating the situation further, Schroeders concluded that compositions 
15 

with 32 mol% or greater Ti02 formed titanates in hafnia-rich mixed oxides which is 

contrary to the findings of Ruh et. al.24 and Coutures and Coutures.18 Basically, the 

degree of solubility of Ti02, the two-phase regions, and the effect of Ti02 on the lattice 

parameters and transformation temperatures in the HZT system have never been 

reported conclusively. 

2.2 Property Studies of Hafnia, Zirconia, and Titania 

In the last section, a basic understanding of the structures and systems of 

hafnia, zirconia, and titania was given. This section will describe selected electrical 

and mechanical properties, then cite typical property values and compositional effects 

in the various hafnia, zirconia, and titania systems. 

2.2.1 Electrical Conductivity 

The electrical conductivity is a measure of the ability of a material to transport a 

charge. For typical ionically bonded ceramics, point defects are primarily responsible 

for charge transport. Ionic solids contain point defects at all temperatures above OOK. 

The presence of ionic defects gives rise to ionic conductivity, while that of electronic 

defects results in electronic conductivity. The total electrical conductivity is given by 

the equation: 

m=Cni (zi e) (1) I 

where ni, 3, and are the concentration, valency, and mobility, respectively, of the ith 

charge carrying species and e is the electronic charge? Ionic and electronic defects 

in ionic solids are presented using KrOger-Vink notation which specifies the kind, 

location, and effective charge of a defect.29 Some of the typical defects for a M02 

fluorite ionic crystal are given in Table 3 along with a basic description. Which of the 

point defects is the main mobile species depends on temperature, oxygen partial 

pressure, and dopant concentration. Figure 9 shows an example of the kinds of 
16 

pressure-temperature dependences that can occur. Dopants also alter the total 

conductivity by causing anion vacancies to be formed (thereby increasing ionic 

conductivity) or by acting as donors or acceptors (thereby increasing electronic 

conductivity). 

Table 3. 

Symbol Description Effective 

Charge 

00 Oxygen in Oxygen Site Neutral 

O;( Interstitially Incorporated Oxygen Ion Negative 

VOW Doubly-Ionized Oxygen Vacancy Positive 

VM"" Tetravalent Cation Vacancy Negative 

FM' Substitutionally Incorporated Trivalent Cation Negative 

Interstitially Incorporated Tetravalent Cation Positive 

Electron in Valence Band Neutral 

Flev 

Typical Point Defects in M02 Ionic Crystals. 

e' Electron in Conduction Band 

h" Hole in Valence Band 

Negative 

Positive 

LOG up 

. LOG 

Figure 9. Schematic diagram of the Px,, T dependence of partial and total 

conductivities. A: OP >> eon, 0,. B-B: OP = On > Wan. C: On >> eon, O~ 

D: eon >> OP + On (Ref. 30). 
17 

2.2.1.1 Systems 

An examination of several fluorite systems will elucidate some of the finer points 

of defect chemistry and electrical conductivity. 

Pure Hf02 and Zr02 in both the monoclinic and tetragonal form have been 

studied by several investigators.l2#31*7 From these studies, and those of Choudhary 

et. al.,28 the following summary of the defect nature of Hf02 and Zr02 can be made. 

The defect formation reactions can be written as: 

% 

F2 

K3 

00 -L 0~+VoW (2) 

o0 * fo2+voH+2e* (3) 

bo2* Or+2h’ (4) 

where the mechanism described in Equation (2) is a thermally generated Frenkel 

defect. The mass-action law for each becomes, respectively: 

K1 = [Oi’l V0-l (5) 

K2 = n2 PQ [voo3 (6 1 

KS = PQ-~ [Or] (7) 

1 

1 

where [ ] indicates defect concentrations and n and p are the electron and hole 

concentrations, respectively. Electroneutrality requires that n + 2[0(’] = p + 2[V,’q. The 

defect concentrations as a function of partial pressure then become: 

In Region I (low Po,) where the prevailing electroneutrality condition is n =2D/o”l, 

1 1 

n = (2K2)S PO,'^ (8) 

[VO”] = (K2/4)5 PQ’~ (9) 

(1 0) 

(11) 

= [Or] = K1Z (12) 

(1 3) 

(14) 

1 1 

[Or] = (4Kl3/K2)5 PO,^ L 

and p = (K*K33/4K13)i PO) 

In Region II (medium PO,) where [Otl 6 Po”), 

1 

1 1 

n = (K22/K1)s Po,“ 

and p = (K32/Kl)S PO,’ 

1 1 
18 

And in Region 111 (high PO,) where p = 2[0?], 

I 1 

p = (2K3)5 Po2Z (15) 

[Or] = (K3/4)5 PO,^ (1 6) 

I I 

= (4K13/K3)% Po;) (1 7) 

1 

and n = (K3K23/4K& PO;^ (1 8) 

The Kroger-Vink diagram which shows the expected defect concentration 

dependence of Hf0, and ZrO, on the oxygen partial pressure is shown in Figure 10. 

The diagram was constructed by Douglass and Wagner38 using equations similar to 

Equations (8)-(18). 

"/ [VnOOI \ ~ 

LOG OXYGEN PARTIAL PRESSURE 

Figure 10. Kroger-Vink diagram showing pressure dependence of defect 

concentration for HfOe and ZrO, (Ref. 38). I 

The partial ionic conductivity (ob") and partial electronic conductivity (a,,) in 

terms of the defect concentrations are given by: 

a, = 2F(pi[Oi"] + J.LJV;~)/VM (1 9) 

and oel = F(ven + p+,p)/V~ (20) 

where F is the Faraday constant; VM is the molar volume of Za2; and pi, p,,, pet and ph 

are the mobilities of interstitial oxygen ions, anion vacancies, and excess electrons 
19 

and holes, respectively.36 In general, the mobility of an electron is 100 to 1000 times 

greater than that of interstitial anions. Also, in M02 fluorites, the mobility of anion 

vacancies is much greater than that of interstitial anions.36 

The oxygen pressure dependence of the ionic and electronic conductivities are 

then: 

In Region I (low pressures), 

1 1 1 

Om, = (~F/VM){R(~K~~/K~)’ P0,C’ + h(K2/4)’ PO;’} (21) 

1 1 

and Gel E (Fk/V~)(2K2)5 PO,'^ (22) 

when pen phP 

In Region II (medium pressures), 

00, = (~F/VM)& + h)Kli (23) 

and Gel P (F&,/VM)(K~~/K~)~ PQ’ (24a) 

Gel E (F~/VM)(K~*/K~)’ PO,“ (24b) 

1 1 

when phP y, pen 

1 1 

when pen y, phP 

And in Region 111 (high pressures), 

A1 

OD, = (2F/V~){&(K3/4)3 PO,’ + h(4K13/K3)f Po[&} (25) 

(26) 

1L and Gel P (Fk/V~)(2K3)5 Pq‘ 

when PhP y, Pen 

The total electrical conductivity is the sum of the ionic and electronic 

conductivities. Figure 1 1 shows the experimentally determined pressure and 

temperature dependence of the total electrical conductivity of pure monoclinic Zr02 

from the work of Kumar e€. al.? At low pressures, the conductivity is shown to have a 

limiting PO;& dependence (n-type). The conductivity at high pressures is shown to 

have a PO,^ dependence (p-type). An apparent discrepancy exists in Region 111 

between the expected PO,^ dependence and the observed PO,$ dependence. A 

1 

1 

similar discrepancy was observed in monoclinic thoria by Choudhary et. al..28 
20 

Figure 11. Temperature and pressure dependence on electrical conductivity of pure 

monoclinic ZrOe (Ref. 36). 

Choudhary et. al. and Kumar et. al. examined the differences in the theoretical 

and experimental high pressure conductivities. Both concluded that the discrepancy 

occurred as a result of impurities which modified the controlling defect mechanisms at 

high pressures. In an effort to characterize the effects, Kumar et. al. chose to examine 

the electronic and ionic contributions of the conductivity. Coupling the experimentally 

determined transference numbers with the measured total conductivities, they were 

able to separate the ionic and electronic components from the total electrical 

conductivity. They showed that the electrical conductivity of zirconia was ionically 

controlled below approximately 7OOOC and electronically controlled above 700°C for 

10-22 S PO, S 1 atm. A schematic diagram of the expected pressure dependence on 

the electronic and ionic conductivities for the electronically controlled situation 

(modelled after Kumar et. al.) is shown in Figure 12. Two possibilities are presented. 
I 21 

Kumar et. al. found that their data was not accurate enough to distinguish between the 

two. However, based on Arrhenius plots of the ionic conductivity, Kumar et. al. 

concluded that the data for ZrO, was consistent with an anti-Frenkel disorder in which 

singly ionized oxygen interstitials predominated at high pressures. Singly ionized 

1 

I 

I 

oxygen interstitials occur when impunties are present and lead to a PO,’ dependence 

in the high pressure region in agreement with the experimental data. Nonetheless, in 

order to fully verify the fluorite defect model presented in Equations (2)-(26), 

conductivity data for monoclinic hafnia, zirconia or thoria without impurities should be 

collected and plotted to show the proper PO, dependence. 

I 

LOG OXYGEN PARTIAL PRESSURE 

I Figure 12. Schematic diagram of the expected pressure dependence on the 

electronic and ionic conductivities for Zr02 at high temperatures 

(Ref. 36). 
22 

Numerous studies of doped HfO, and ZrO, have been made. Some of the more 

recent review articles are by Choudhary et. a1.,28 Dell and Hooper,39 and Subbarao.8 

A summary of the main aspects follows. 

Several types of dopants can be added which may alter the conduction 

mechanisms of the fluorite structure. Mixed oxides having cations with essentially 

invariant valence (e.g., ZrS+, Hf4+, Th4+, S$+, Y3+, La3+, Mg2+, and Ca2+), generally tend 

to exhibit purely ionic conductivity (independent of oxygen partial pressure) over a 

wide range of oxygen pressures. Those cations which reduce to lower valences 

(e.g., Ce4++Ce3+, Fe3++Fe2+, Eu~++Eu~+), or those which oxidize to higher valences 

(e.g. Ni2++Ni3+, Co2++Co3+, and U4++U6+), will tend to introduce electron or hole 

conduction, respectively, via hopping mechanisms to the otherwise ionic conductor. 

When aliovalent oxides with stable valences are substituted as impurities into 

the fluorite structure, they usually tend to stabilize the cubic phase and produce 

vacancies in the oxygen sublattice if their valence is less than four.28 The electrical 

conductivity increases due to the increased concentration of oxygen vacancies. The 

temperature dependence exhibited by the electrical conductivity approximates 

Arrhenius behavior.28 A limit to the enhancement of conductivity with dopant is 

reached when the oxygen vacancy mobility decreases as a result of lattice distortions 

and the formation of (FM"-V~'") complexes or vacancy cl~sters.*8~4o41 Figure 13 shows 

the maxima in conductivity with impurity concentration for some typical Zr02-based 

systems. 

For aliovalent impurities added to the fluorite structure as in Y203-stabilized 

Tho,, the electroneutra1if)r requires: 

p + 2[VJ = n + 2[0/'] + [yTh'] (27) 

where [Y,,'] is the concentration of Y3+ ions on Th4+ sites. Following an analysis 

similar to that used for pure ZrO,, the Kroger-Vink diagram in Figure 14 for a doped 
23 

mol X M2O3, MO 

Figure 13. Maxima in conductivity for various doped ZQ-based materials at 

T= 1080°K (Ref. 41). 

I I ZONE1 I ZONE II ! 7ONE 111 

LOG OXYGEN PARTIAL PRESSURE 

Figure 14. Kroger-Vink diagram showing pressure dependence of doped fluorite 

structure (Ref. 28). 
24 

and is constant with temperature.28 On the other hand, 

fluorite structure was constructed. Also, Figure 15 shows the expected oxygen partial 

pressure dependence on conductivity. The effects of aliovalent impurities is primarily 

seen in the high pressure region. For the condition that [YTh’] = p, the mass-action 

laws lead to the pressure dependence seen in Zone 111 of Figure 15. In this region, the 

hole conductivity becomes independent of oxygen partial pressure. A fourth zone 

which is not shown in Figures 14 and 15 occurs for very high oxygen partial pressures. 

The condition p = 2[0(’] exists giving concentration and conductivity dependencies 

similar to those observed in the high pressure regions of the pure state. 

I I I 1 I n-T y pe I Mixed I lonic I Mixed I p-Type 

semiconductor I conductor Iconductor I iconductor I semiconductor 

I I I I ‘ QP 

Qion 

I 

I 

I 

I 

I / \ I 

I I 

I I 

LOG OXYGEN PARTIAL PRESSURE 

Figure 15. Schematic diagram of the expected pressure dependence on the 

electrical conductivity of a doped fluorite structure (Ref. 28). 
25 

(30) 

-AG [VJ = [or] = (~~0); exptd) 

at which point the material shows intrinsic behavior.28 Figure 16 demonstrates this 

point. Intrinsic behavior is rarely observed in doped oxide electrolytes, however.28 

1/T 

I Figure 16. Schematic diagram demonstrating the effect of the temperature on defect 

concentration in M02-F203 at constant oxygen partial pressure (Ref. 28). 

When multivalent cations are added to the fluorite structure, an increase in 

electronic conductivity results. A possible defect ionization reaction for a dopant 

incorporated substitutionally into the fluorite structure would be: 

(31 1 *+ FM+e +FM* 

where FM and FM' represent substitutionally incorporated tetravalent and trivalent 

cations on cation sites, respectively. The electronic conductivity is increased via 

electrons hopping between dopant cations.4 

If a multivalent cation is incorporated interstitially into the fluorite structure, the 

reaction would be: 

F02 3 FY"+ 46' + O2 (9) (32) 

where F is a tetravalent cation.4 The excess electrons could then be incorporated 

either by FT or Vooo following the equations: 
26 

FI"" + e'+ FI" (33) 

or Vom+ 2e'+V0 (34) 

The latter situation is the case for TiO, doped Y203-stabilized ZrO, as reported 

by Worrell and Weppner.44 In their study, they found that titanium is incorporated 

interstitially into the fluorite structure: however, its effective charge is unknown. The 

result is a mixed ion-electron conductor. The electronic conductivity is high at low and 

high pressures. Figure 17 shows the temperature dependence of the electrical 

conductivity at various oxygen partial pressures for two titania dopant concentrations. 

ft'C1- 

1oo0950 900 850 800 750 700 - -1.01 I1 I 1 I 6 I c 

k 

b 

z -1.5 - 

u 

c 

Y 

0 

- 

-20 - 

I' 

-2c 

PO, = 0.21 atm (air) 

(88mlo 210, -12mlo Y,O,) 

STRICKLER L CULSON 

8 8.5 9 9.5 10 105 4' 

- 

* 

x loL CK"1 

(a) 

f['Cl - 

1m950900 850 800 750 700 I1 I I I I 

Figure 17. The dependence of electrical conductivity on temperature for various 

oxygen partial pressures for (a) 1 mol% and (b) 10 mol% Ti02 in 

Y203-stabilized ZrO, (Ref. 44). 
27 

2.2.1.2 Measurements 

High temperature electrical conductivity measurements are generally made 

using two- or four-probe methods.39 One of the primary concerns in measuring with 

probes is electrode polarization (Le., the build-up of charge at an electrode). The 

design of the four-probe method eliminates these effects.45 Polarization in the twoprobe method is eliminated by using ac current (500 kHz to 5 MHz).~~ The sample 

configuration is best when the ratio of probe separation to sample thickness is large. 

This assures that current paths are straight. 

Resistance measurements are taken for a sample of known cross sectional area 

(A) and probe separation (I). When a known current (I) with known voltage (V) is 

passed through the sample, the resistance (R) becomes: 

(35) 

V R=- I 

The relation between measured resistance and the resistivity (p) is: 

(36) 

A p=Rr 

And the conductivity is then just the inverse of the resistivity. 

When measurements are made as a function of oxygen partial pressure, two 

factors need to be considered. First, the movement of gas away from the electrode 

should be rapid; second, the gas mixture which establishes the PO, should not react 

with the sample (altering the defect chemistry). In address to the second of these 

factors, most conductivity studies use either H2/H20 or CO/CO2 gas mixtures to 

establish low oxygen partial pressures. McClaine and Coppel35, however, reported 

observations of interactions of H2/H20 and CO/CO2 gas mixtures with pure tetragonal 

Zr02. (Carniglia et. al.46 later noticed similar behavior in their studies of Tho2 and 

Zr02.) McClaine and Coppel found that increased electronic conductivities occurred 

when H2/H20 gas mixtures were used to establish oxygen partial pressures as 
28 

opposed to when CO/CO2 gas mixtures were used. They attributed this increased 

conductivity to the presence of hydrogen entering the matrix as interstitial positive ions. 

(A similar effect occurs in ZnO.) They found that the use of CO/CO2 gas mixtures had 

much less of an effect on the electrical properties. (Both a high concentration and a 

highly active species is required for the gas reactants to dominate the defect chemistry. 

Apparently, neither condition prevails in the electrical conductivity of Zr02 electrolytes 

in CO/CO2 gas mixtures.) From their studies, therefore, McClaine and Coppel 

recommended the use of CO/C02 gas mixtures for low pressure conductivity studies 

of pure Zr02. Nonetheless, further examination of the use of CO/CO2 gas mixtures, as 

well as H2/H20 gas mixtures, is warranted. 

2.2.2 Thermal Expansion 

Thermal expansion is the linear or volumetric change that accompanies an 

increase in temperature. The linear thermal expansion coefficient (a) is defined as the 

ratio of the change in length per "C to the length at 0°C and can be expressed as: I 

where Lo and LT are the lengths at 0°C and some temperature T, respectively. 

Most materials when heated tend to enlarge or expand in response to the 

increase in temperature. Such materials are said to have positive thermal expansion 

coefficients. Typically, thermal expansion coefficients of ceramic materials are of the 

order of 4.0 to 14 x lO-S/OC. For applications where rapid cycling between high and 

low temperatures is necessary, materials with low or zero coefficients are best. (These 

materials tend to exhibit high thermal shock resistance.) Skaggss points out that 

I 

I 

I 

common factors in low thermal expansion materials are: (1) The symmetry is never 

cubic, infrequently tetragonal, possibly monoclinic, and usually orthorhombic. 

(2) Strong axial expansions and contractions are present (e.g., the axial thermal 
29 

expansion of orthorhombic Hffi04 is very strongly negative along the b-axis, while it is 

moderately positive along the a-axis). (3) The material is usually in ternary 

combinations. (4) A large open spaced lattice is usually involved. (5) A hysteresis in 

the thermal cycle usually occurs. 

2.2.2.1 Systems 

A number of ceramic materials with low thermal expansion coefficients have 

been cited in the literature.1-5 Many compositions in the hafnia, zirconia, and titania 

systems are among these low coefficient materials. An examination of several 

materials will elucidate the mechanisms involved in low thermal expansion. 

The thermal expansion of Hf02 and Zr02 has been studied by several 

a~thors.l%l8,22~47-54 Table 4 summarizes the axial and linear thermal expansions of 

the monoclinic and tetragonal phases found in these investigations. Table 4 also 

gives the calculated linear thermal expansion coefficient (acalc.) of the monoclinic 

phases, determined from the x-ray data of Ruh et. al.,54 as well as the volumetric 

thermal expansion coefficients of the unit cells (Vunfi cell). For ZrO,, the measured 

thermal expansion coefficient compares very well with the calculated value, whereas, 

a difference is observed for Hf02. This difference is associated with microcracking 

which Dole et. al.52 have shown indeed occurs in Hf02. Microcracking leads to open 

cracks at low temperatures that close as temperature increases? The effect is that the 

linear thermal expansion coefficient of the bulk material appears lower than the true 

lattice thermal expansion coefficient. In thermal expansion measurements, 

microcracking is observe8 as a hysteresis in the thermal cycle. Figure 18 shows the 

type of hysteresis associated with microcracking found in thermal expansion data. 

This hysteresis, and thus microcracking, is eliminated in fine grained (<2 pm) 

materials.52 In addition to the hysteresis associated with microcracking, a hysteresis in 

the thermal cycle associated with the monoclinic to tetragonal phase transformation 
30 

Table 4. Axial and Linear Thermal Expansion Coefficients of Hf02 and Zr02. 

Thermal Expansion Coefficient (1 0-VOC)' 

Parameter z*2, zro2tet "f02rnO"O HfO2tet 

a 9.2 1 1.649 6.118 6.118 

b 1.4 0.018 

C 10.9 16.149 9.618 5.618 

0 -4.5 -3.6 

Vunit cell 22.9 37.3749 20.6 37.218 

ameas. 7.2 10.6349 4.4-6.5 10.047 

%ale. 7.6 6.9 

'All values from Ref. 54 unless othetwise noted. 

SOLID SYMBOLS REPRESENT COOLING DATA 

400 800 I200 

TEMPERATWE ('C) 

Figure 18. Linear thermal expansion coefficient of microcracked (>2 pm grains) and 

crack-free (<2 pm grain size) Hf02 (Ref. 52). 
31 

occurs. This temperature hysteresis can be seen in the phase diagram in Figure 6 for 

the system Hf02-Zr02. Figure 19 shows a typical hysteresis in the thermal expansion 

associated with the monoclinic to tetragonal phase transformation. The volume 

change of the transformation is 2.7%51 and 3.0%49 for Hf02 and Zr02, respectively. It 

is this volume change which is used in transformation toughening and which is 

deleterious to pure unstabilized Zr02 and HfO,. 

The thermal expansion coefficients for a wide range of compositions in the 

Hf02-Zr02 system were determined by Stansfield.21 In general, the coefficient 

increased as Zr02 was added to Hf02 while the transformation temperature 

decreased . 

An extensive study of the effect of various additives (e.g., CaO, MgO, etc.) to 

Zr02 on the thermal expansion was made by Curtis.55 Among other findings were the 

following: (1) Small additions of CaO and MgO (metastable region) reduce the 

inversion hysteresis and leave the thermal expansion coefficient relatively unchanged. 

(2) Large additions of CaO and MgO (cubic stabilized region) eliminate the inversion 

hysteresis but raise the thermal expansion coefficient. 

The effect of adding Ti02 to HfO, and Zr02 on the thermal expansion has been 

studied by several investigators.1.5*laD19,*4.s6-5a Analogous to the pure oxides, the 

thermal expansion behaviors of the two systems are quite different. For additions of 

TiO, to Zr02, Brown and DuwezlQ showed that the thermal expansion coefficient 

decreased with increasing TiOp. Figure 20 shows the thermal expansion curves from 

their study. The thermal expansion coefficient is ca. 7.5 x 10-6/OC for both pure Zr02 

and the 10 mol% Ti02 cbmposition. For the 20, 30, and 40 mol% compositions, a 

gradual decrease from 6.7 to 5.0 x 10-6/OC occurs. Phase analyses show that for 

compositions greater than 15 mol% TiO,, ZrTi04 begins to appear in the monoclinic 

solid solution. The increase in the concentration of the titanate phase slightly reduces 
32 

J 0 028 \ 

2 0024 

- 

C $ 0020 c 

L" 

., 0016 0, 

C 

0 

c" 0012 

- 

0.008 c 

u 

f 0004 

0 1000 2000 3000 4000 

c 

'C 

0 

0 

5 

a 

C 

0 

n 

C 

0 

Q 

.- 

w. 

W 

6 

0 

C .- 

A 

* 200 400 600 800 y)o I200 

Temperature ,t 

(b) 

Figure 19. Linear thermal expansion demonstrating transformation temperature 

hysteresis of (a) Ha2 (Ref. 47) and (b) ZrO, (Ref. 49). 
0.0 

0.6 

0.4 

0.2 

0 

IIII1(Lj 

-0.5 0 400 m 

2 

Figure 20. Linear thermal expansion curves of Zr02 containing various mole 

percentages of TO2 (Ref. 19). 

the expansion and leads to a decrease in 

tetragonal phase transformation hysteresis. 

In contrast, for additions of Ti02 to t 

the magnitude of the monoclinic to 

f02, large reductions in the thermal 

expansion have been observed. Low, zero, and even negative thermal expansion 

coefficients have been reported for compositions in the Hf02-Ti02 system.115,18,24.56-s8 

Figure 21 shows the linear thermal expansion of a number of Hf02-Ti02 compositions. 

For compositions between 30 and 40 mol% Ti02, very low coefficients (4 x 10-6/oC) 

are observed, while for compositions greater than 40 and less than 30 mol%, 

moderate coefficients (-4 x lO-S/OC) are found. The compositional dependence of the 

thermal expansion is summarized in Figure 22. An examination of the Hf02-Ti02 
34 

I 

Figure 21. Linear thermal expansion for several Hf02-Ti02 compositions (in mol%) 

(Ref. 56). 

L4 - 

e ‘I I831 

0 

0 

I I 1 I I I n* a n Y Y Y n 88 

4.1 

Tii WLC u 

Figure 22. Linear thermal expansion vs. Hf02-Ti02 composition at 750°C. Open 

squares (Ref. 56). Open circles (Ref. 24). 
35 

phase diagram shows that the compositions with the lowest coefficients lie within the 

two-phase (HfTiO, plus monoclinic HfO,) region. Three conditions which are present 

in the Hf0,-TiO, system but are not present in the Zr02-Ti0,system have been 

associated with the low thermal expansion coefficients. First, the b-axial thermal 

expansion is negative for HfTi04 in contrast to the positive expansion which occurs for 

ZrTi04 (see Figure 23).24 Second, a hysteresis in the thermal expansion data is 

I observed for HfTiO,, but is not present in the ZrTi04 data.58 Third, microcracking 

occurs in the Hf02-Ti02 system but not in the ZrO2-TiO2system.58 Each of these 

conditions leads to low thermal expansion in the Hf02-TiO, system. Similar to the 

ZrOe-Ti02 system, the titanate phase decreases the magnitude of the monoclinic to 

tetragonal phase transformation hystereses for compositions in the Hf02-Ti02 system. 

Figure 23. Axial thermal expansion of orthorhombic HfTi04 (solid circles), 

Hfl.26Ti0.7404 (open circles) and ZrTiO, (open squares) (Ref. 57). 
36 

In the HZT system, many reports of compositions with low thermal expansion 

coefficients have been made. In almost every study, the low coefficient materials were 

within the titanate plus monoclinic two-phase region. Exceptions were found in the 

work of Schroeder3 and Brown et. aL.4 Both investigations report compositions with 

low thermal expansion coefficients in the monoclinic solid solution region of the phase 

diagram. For compositions where the titanate phase is present, the low thermal 

expansion coefficients are attributed to microcracking which results from the negative 

b-axial thermal expansion of the titanate phase. Schroeder showed that the b-axis of 

the monoclinic phase in the composition that he studied also exhibited negative 

thermal expansion. He attributed the low linear thermal expansion coefficient to the 

negative b-axial thermal expansion, but never addressed the possibility of 

microcracking. He presented an openness model to explain the reduced expansion 

tendencies. In the openness model, titanium ions, which are smaller than either 

hafnium or zirconium ions, assume cation positions on the lattice which is oversized 

for the titanium ion. The result is an increased openness in the lattice which can better 

accommodate thermal agitation, thereby decreasing the rate of thermal expansion.59 

The openness model, however, states that once the titanate phase begins to form, an 

increase in thermal expansion occurs which is contrary to the reported observations. 

Table 5 summarizes the thermal expansion behavior of various compositions in 

the HZT system. For compositions which are well within the known titanate plus 

monoclinic solid solution region, the coefficients decrease with additions of Zr02 up to 

a limit (ca. 13%) and then increase. In addition, from the work of Brown et. al.,4 a 

range of compositions and coefficients were reported. The range of coefficients were 

-1 to 5 x lO-S/OC for compositions 60 to 85 mol% Hf02, 10 to 30 mol% ZrOe and 

10 to 30 mol% Ti02. It is suspected that the lowest coefficient corresponded to a 

Ir 

composition with approximately 60 mol% Hf02-l 0 mol% zr02-30 mol% TO2. One final 

comment regarding the effect of Ti02 on the expansion behavior of HZT compositions, 
37 

the monoclinic to tetragonal transformation range is widened and the volume change 

is lowered when the concentration of titania is increased.3 

Table 5. Thermal Expansion Coefficients for Various Compositions in the HZT 

System. 

Composition (mol%) Thermal Expansion Coefficient Ref. 

63 0 37 0.5 1 

61 3 36 -0.9 1 

59 5 36 -1 .o 1 

58 8 35 -1.5 1 

55 10 35 -1.8 1 

52 13 35 -4.4 1 

47 16 37 2.2 5 

32 32 37 5.9 5 

16 47 37 6.7 5 

60 20 20 3.1 4 

H Z T (1 O-VOC) 

~ 

2.2.2.2 Measu re men ts 

Measurements of thermal expansion can be made on both the bulk material 

(linear thermal expansion) and the lattice or unit cell (axial thermal expansion). 

For linear thermal expansion measurements, a dilatometer is used to monitor 

the change in length of a bulk sample when heat is applied. A factor to contend with in 

using a dilatometer is the expansion of the unit itself during the measurement. 

Usually, a standard with a known thermal expansion is run to calibrate the apparatus. 

Once the expansion of the apparatus is known, the thermal expansion of a test sample 

can be determined by adding the apparatus expansion to the measured expansion. 

The final equation for determining the expansion of a material using a single rod 

dilatometer is then: 

&le = (&tandad - &and. meas.) + &le meas. (38) 
38 

where AI- is the change in length. As long as the conditions are identical for the 

calibration run and the sample run, the final expansion of the material can accurately 

be determined. To assure identical conditions, a dual rod dilatometer can be used 

which measures the standard and the test sample simultaneously. 

For axial thermal expansion measurements, hot-stage x-ray units are used. The 

lattice parameters are determined at various temperatures and the parameter change 

versus temperature is plotted to give the axial thermal expansion. For isotropic 

materials, the axial thermal expansion is equal to the linear thermal expansion. For 

anisotropic orthonormal polycrystalline materials: 

a= Q (019 + at, + ac) (39) 

-L 

where G is the linear thermal expansion coefficient and aa, ab, and & are the axial 

thermal expansion coefficients. For non-orthorhombic systems, (e.g., monoclinic), the 

conversion between axial and linear thermal expansion coefficients requires matrix 

transformations. 

2.2.3 Fracture Strength and Toughness 

The fracture behavior of a structure depends on many factors besides the 

inherent strength of the material. Factors such as flaws, microstructure, and 

environment strongly influence the fracture mechanics. This section will discuss the 

strength and toughness of materials and address the factors which influence the 

fracture behavior. 

The theoretical cohesive strength of a solid is governed by the strengths of the 

bonds between constitu&t atoms. Fracture occurs when the separation between 

atoms along a plane reaches a critical distance whereby the attractive forces no longer 

can bind the atoms. This critical distance and the energy required to create new 

surfaces are used in determining the theoretical strength of a solid. For covalently 
39 

bonded solids, the attractive forces can be approximated by the Morse potential.60 

The cohesive strength is then: 

0, = 1/yE/4a (40) 

where y is the surface energy, E is the elastic modulus, and a is the separation 

between atoms. For an elastic material, the theoretical strength is of the order of 

2 x 106 psi (1 4 GPa). 

In general, measured strengths are orders of magnitude lower than theoretical 

strengths. This difference is attributed to flaws which are present in all materials. 

Flaws reduce the strength of a structure by introducing stress concentration effects, 

residual stress effects, or a combination of both. 

lnglis61 was the first to show that flaws in the form of fine cracks act as stress 

concentrators. When a stress is applied perpendicular to the fine cracks within a 

structure, the stress is magnified at the crack tips. The resultant concentration of 

stresses causes the structure to fracture at applied stresses which are much lower than 

expected from considerations of theoretical strengths. Griffith62 used the stress 

analysis of lnglis to show that the fracture strength of a flawed structure is: 

01 = d2yE/~c (41 1 

where c is the length of the critical flaw. Griffith assumed that in brittle fracture, the 

material is entirely elastic and that at failure, the strain energy around the crack equals 

the energy to create the new surfaces. Irwin63 extended the ideas presented in the 

Griffith theory to include the plastic nature of materials and to take into account energy 

lost to mechanisms other than the creation of surface area. Irwin showed that the 

fracture strength of a math is: 

of = -\IG~E/TCC (42) 

where G, is the elastic energy release rate. G, becomes equal to 2y for a purely 

elastic material having only surface energy dissipation at fracture. 
40 

The term G, is related to the fracture toughness (K1c). The fracture toughness, 

also called the critical stress intensity factor, is the resistance of a material to crack 

propagation. A structure will catastrophically fail when the stress intensity factor 

reaches Klc. The applied stress is then: 

KlC 

YdE Of = - (43) 

where Y is a geometrical constant and c is the critical flaw size. 

As Equation (43) points out, the fracture strength of a structure is related to the 

fracture toughness of the material and to the critical flaw size. The fracture toughness 

is considered by many to be a material property that is dependent upon microstructural 

effects (e.g., second phases, grain size, porosity, grain boundaries, microcracks, 

agglomerates, etc.) and environmental effects (e.g., temperature, atmosphere, stress 

distribution, etc.). Consequently, the fracture strength is influenced by these effects 

also. Flaws, which are the principal source of fracture, however, dominate the fracture 

behavior? 

Detailed reviews of the various types of flaws and their effect on the fracture 

behavior of ceramics have been made by Evans64 and by Singh.65 As Evans points 

out, three categories of failure can be distinguished: (1) failure from preexistent flaws, 

(2) failure from the coalescence or joining of microcracks, and (3) failure from the 

formation and coalescence of cavities.64 A brief description of the various failure 1 

I mechanisms and flaw types in ceramic materials follows as well as an examination of 

their relationship to strength. I 

Preexistent flaws are those that are present in a structure prior to the application 

of a stress. They can be separated into two types, internal and surface. Internal flaws 

(e.g., inclusions, voids, pores, etc.) are generally introduced during fabrication. 

Surface flaws, which tend to dominate the fracture behavior, occur in the as-fired 

condition or may be introduced during polishing. Inclusion-initiated failure usually 

I 
41 

involves a combination of stress concentration and residual stress influences, while 

void-induced failures are related to stress concentration effects and surface crack 

failures are a consequence of residual stresses.64-65 

Inclusions are a major source of premature fracture. In general, they introduce 

residual stresses as a result of a mismatch in the thermal expansion coefficients of the 

matrix and the inclusion. . The residual stress can be expressed as: 

where or and at are the radial and tangential stresses, respectively; i and m refer to 

inclusion and matrix; AT is the difference between room temperature and the maximum 

temperature below which stresses are not relaxed during cooling; R is the radius of the 

inclusion; and r is the radial distance from the inclusion.65 Depending on the thermal 

expansion coefficients, elastic moduli, and toughnesses of the matrix and inclusion, a 

variety of fracture mechanisms can occur. (For example, when an inclusion has a 

thermal expansion coefficient that is much lower than that of the matrix, residual tensile 

stresses occur in the matrix. These stresses when large enough produce radial cracks 

that severely degrade the strength of the structure.) A schematic diagram of the 

various inclusion-initiated fracture mechanisms is presented by Evans.64 Some 

fracture modes tend to be more detrimental to the strength than others. Not all 

inclusions when added to a material, however, degrade the strength. Transformation 

toughened materials and glass-ceramics have second phase inclusions that actually 

enhance the strength.66 The strengthening effects of inclusions, however, will not be 

addressed here. b 

Voids or large pores are a frequent source of failure in sintered ceramics. They 

usually serve as stress concentrators by acting in conjunction with other flaw types 

(e.g., extending surface cracks or microcracks).M The reduction in strength depends 
42 

on the size, shape and distribution of the porosity. The strength generally decreases 

with increasing pore size. The relation between the strength and the amount of 

porosity is empirically given by: 

op = a, e-BP (45) 

where op and a. are the strengths with and without pores, B is a constant that depends 

on the distribution and morphology of the pores, and P is the fractional pore 

content .65#67-70 

Surface flaws tend to strongly influence fracture behavior. A large number of 

as-fired surface flaws are introduced in virtually every ceramic fabrication process. In 

many instances, it is necessary to grind or polish the surface to eliminate these flaws. 

The finishing process, however, tends to introduce surface cracks which may also act 

as failure-initiating sites. These surface cracks originate from the residual stresses that 

accompany the localized deformation in the vicinity of the grinding particles. A 

correlation between the grit size of the polishing media and the crack size exists. 

Likewise, a correlation exists between the grit size and the fracture strength. The 

strengths tend to be higher for finer grit polishes.64-65 

Preexistent flaws, whether internal or at the surface, generally influence fracture 

behavior at low and intermediate temperatures (Le., in temperature regions where 

fracture is elastic). Strength in this region is either constant or increases slightly with 

temperature. At high temperatures where plastic deformation occurs, the influence of 

preexistent flaws in reduced.= The stress concentration and the residual stress effects 

introduced by flaws are lessened when plastic deformation occurs. 

#, 

Failure by microcrack coalescence also occurs in the low and intermediate 

temperature range. Microcracks (which form in grain boundaries as a consequence of 

thermal expansion anisotropy) generally act to decrease the strength. The initial 

formation of microcracks does not lead to failure. Fracture in microcracked structures 

occurs when microcracks coalesce or join to form a crack of critical size. Both the 
43 

orientation and the concentration of microcracks strongly influence the fracture 

behavior. For example, dilutely microcracked materials may experience no 

degradation in strength, and highly microcracked materials with the proper relative 

crack orientation, may enhance strength by stopping crack propagation (microcrack 

to ug hen i ng ) .71-7* 

The fracture process at high temperatures usually involves nonlinear 

deformation. Generally, the strength decreases as the temperature is elevated. 

Failure occurs as a result of the formation and coalescence of cavities. Cavity sites 

initiate from the corners of grains, from grain boundaries, from local deviations in the 

chemical composition of the grain boundary phase or from thick amorphous regions 

within the grain boundary. Cavities grow by diffusion and/or viscous deformation. 

Coalescence or the joining of cavities progresses until fracture occurs. The 

degradation of the strength is dependent upon the diffusive transport of material from 

the growing cavity and the nature of the grain boundary phase which forms the 

cavi t y.65 

2.2.3.1 Systems 

Several of the major points of fracture mechanics will be illustrated by 

examining the effects of various additives to hafnia and zirconia. 

A large number of studies on Zr02-based materials have produced 

considerable amounts of data on fracture strength and toughness. Since the data for 

Hf02-based materials is limited, emphasis here will be upon Zr02-based compositions. 

The range of fracture strengths is from 17073 to 1500 MPa74 depending on the 

composition, purity, phases present, and fabrication technique. Toughnesses range 

from 1.875 to 12 MPa-mf.76 

-G 

The strongest, toughest materials are tetragonal Zr02 polycrystals which are 

predominantly single phase? For grain sizes less than 0.1 to 0.3 pm, the tetragonal 
44 

phase is metastable and will transform to the monoclinic phase when a stress is 

applied, thereby transformation toughening the material.77 Strength and toughness 

values for ZrO, with optimum metastabilizing concentrations of MgO or Y2O3 are in the 

range 1200 to 1500 MPa and 10 to 12 MPam%74.78-81 

toughness drops to 1.8 MPama and the strength to 170 MPa.73.81 

Strength and toughness values increase from those of cubic zirconia for 

mixtures of the cubic and monoclinic phases (250 MPa and 3.7 MPa-mf) and the 

tetragonal and cubic phases (650 MPa and 7.1 MPa-mf)? Pure monodinic Zr02 and 

mixtures of the monoclinic and tetragonal phases have strength and toughness values 

that fall intermediate between these. 

The weakest, least tough materials are fully stabilized cubic zirconias. The 

Metastable tetragonal hafnia has shown toughness values of the order of 

10 MPa-mi for compositions with 46% tetragonal phase.82-83 (Strength values for 

tetragonal hafnia compositions should be comparable to those of metastable 

tetragonal zirconia.83) As the amount of monoclinic phase increases, the toughness 

decreases similar to the Zr02 situation. As 100% monoclinic Hf02 is approached, the 

toughness drops to approximately 2 MPad.82 This value is much less than that of 

100% monoclinic ZrO,. (The strength of monoclinic HfO, should be much less than 

that of monoclinic Zr02 since extensive microcracking occurs in Hf02.) 

Titanate compositions in the Hf02-Ti02 system were studied for fracture strength 

by Hoagland et. al.84 and Simpson.1 These studies showed that HfTi04 weakens the 

material dramatically. Table 6 presents the data from the two studies. The strengths 

for compositions betwee; 30 and 40 mol% were found to be the lowest. This range 

coincides with the low thermal expansion coefficient range. The low strengths can be 

attributed to the extensive microcracking which results from the anisotropic thermal 

expansion.84 
45 

Table 6. Strength of Compositions in the Hf02-Ti02 System. 

Composition (mol% Ti02) Strength (MPa) Ref e re nce 

30 44.3 84 

30 30.5 1 

37 40.9 84 

37 34.0 1 

40 41.8 84 

45 82.4 84 

50 64.6 84 

The fracture strength of 61 mol% Hf02-3 mol% Zr02-36 mol% Ti02 was found to 

be 36.4 MPa.84 This represents a slight increase from the base composition 

63 mol% Hf02-37 mol% TO2. 

The effect of temperature on the strength and toughness of a typical fluorite 

material (partially stabilized zirconia) is illustrated in Figure 24. A general decrease in 

both the strength and the toughness with increasing temperature is observed. Fluorite 

materials with low thermal expansion coefficients, on the other hand, typically exhibit 

an increase in strength with temperature to an upper critical temperature (e.g., Figure 

25). This increase in strength results from the closing of microcracks. 

2.2.3.2 Measurements 

A number of techniques have been developed to measure fracture strength and 

toughness of brittle ceramics. A summary of the various techniques for determining 

strength has been compiled by Evans86 and Quinn et. aL.87 Techniques for 

determining toughness ca’n be found in a number of studies.88-99 

Strength measurement techniques consider such factors as load direction (Le., 

compression or tension), specimen configuration (e.g., three-point bend, four-point 

bend, double cantilever beam, compact tension, double torsion, tapered cantilever 
700 Test canditions: 

4 pa'nt bending 

dG/d t = 1.75 MfWscc 

surkce as round 

* 

lair 

600. 

500 

Mg-PSZ (ZNCO) 

V 0 

4 1 700 Test canditions: 

4 pa'nt bending 

dG/d t = 1.75 MfWscc 

surkce as round 

lair 

loo 1 

6T - 200 - 400 ' 660 800 1000 

Tcg.cmturc I *C 1 b 

SENB 

I 

'RT 260 ' 400 600 ' 800 ' 1000 

LLmpktwr I 'C I b 

Figure 24. Effect of temperature on (a) strength and (b) toughness of PSZ (Ref. 85). 
47 

12 

0 10 

lu8 

a 

3 

3 ‘6 

a 

0 

v) 

Q 

0 

0 - 

W 

U 

O4 v) 

3 

5 

02 

0 

z 

I I I I I 0 

0 500 1000 1500 2000 2500 30 

TEMPERATURE (“F) 

10 

Figure 25. Effect of temperature on strength of three compositions in the HZT system 

(Ref. 1). Composition values are in weight percent. 

beam, and constant moment),86 specimen size, and surface finish @e., polished and 

indented or notched). Even though no standard method for determining strength has 

been established, variations of the four-point bend method are most often used in the 

literature. Although other methods of measuring strength have advantages, only the 

method and advantages of four-point bend measurements will be discussed here. 

2.2.3.2.1 Fracture Strength 

Figure 26 shows a schematic drawing of the stress application in a four-point 

bend fixture. The top surface of the specimen is put into tension (the bottom in 

compression) when moments are applied to the bar ends. The fracture strength (also 

called the modulus of rupture, four-point bend strength, or fracture stress) is 
48 

< L > 

+X+ 

V v I 

1 

Figure 26. Schematic drawing of the load application in a four-point bend fixture. 

+I+ 

determined from the applied load at fracture, the bar dimensions, and the fixture 

dimensions according to the equation: 

kb-l 

3P(L;1) 

OMOR"2(bh ) 

where GMOR is the fracture strength, P is the applied load, L and I are the inner and 

outer spans between load points, b is the width of the bar, and h is the bar height or 

thickness. This equation can be derived as follows: 

Mc 

om=- I 

where: 

P P (L-l) P(L-I) M=-x-F(x- 2 )= 

h 

C=F 

(47) 

(48) 

(49) 

where M is the applied moment, x is some distance away from the outer span load 

point along the length of the bar, c is the distance away from the top surface to the 

plane where there is zero compression and tension, I is the moment of inertia, y is a 
49 

distance in the direction of the thickness, and A is the cross sectional area of the bar. 

Substituting Equations (48)-(50) into Equation (47) gives Equation (46). 

Factors such as bar sizes, fixture design, fixture dimensions, and bar surface 

finish are still being investigated in an effort to standardize the four-point bend method. 

Sample sizes approximately 5.1 cm x 0.6 cm x 0.3 cm and outer span to inner span 

ratios of 3.0 are most often used in the literature. Size and span ratios have a bearing 

on the measured strength in that they affect the flaw distribution. Fixture designs 

attempt to give uniform point load applications and precise bar alignments. Most 

designs use rods to achieve point load applications and to allow fixture movement for 

uniform loading. 

Surface flaws of suitable size are usually placed on the tensile side of a fourpoint bend specimens in order to give a controlled fracture. Fracture initiates at the 

site of this flaw because it tends to be the worst in the specimen. Fracture strength 

values for specimens with controlled flaws tend to have less scatter than values for 

specimens without controlled flaws. There is, however, a consequent reduction in the 

measured fracture strength when large flaws are intentionally introduced. 

The major disadvantages of the four-point bend method (and flexure test 

methods in general) are: (1) Only small volumes of material are exposed to the full 

applied tensile stress. This minimal exposure may lead to inaccurate strength 

estimates. (2) Strength measurements are highly dependent upon the surface finish of 

the specimens. (3) Biasing of the data can occur as a consequence of which statistical 

analysis method is selected, the inadvertent elimination of weaker specimens during 

sample preparation, and the presence of large flaws such as pores that reduce the 

effective cross sectional area.87 The major advantage of the four-point bend method 

over other flexural methods is that the volume of material exposed to uniform tensile 

stresses is the largest (is., a constant tensile stress is present on the specimen surface 

within the inner loading pin span). 
50 

2.2.3.2.2 Fracture Tough ness 

Fracture toughness measurements are now widely made using indentation 

techniques. These techniques fall into two main categories: (1) tests in which Klc is 

determined from direct measurements of crack size as a function of indentation load 

and (2) tests in which the indentation crack serves as a controlled flaw in a bend 

specimen, so that Klc is determined by a strength measurement.99 Although the 

methods in both categories tend to give reproducible measurements, their accuracy 

levels are about 30 to 40% of the true Klc values for most materials. 

A recently developed indentation technique, however, has proven to be both 

reproducible and accurate. This technique was developed by Cook and Lawn99 and 

is a hybrid of techniques from both categories. Certain problems associated with other 

indentation techniques are by-passed; namely, (1 ) the determination of the elastic 

modulus to hardness ratio (E/H) and (2) the applicability of the E/H ratio to non-ideal 

elastic/plastic materials. Cook and Lawn have found that when two or three Vickers 

indentations are placed within the inner span of a four-point bend specimen (see 

Figure 27), all experience near-identical stress histories during the bend test, but only 

one fails. The critical crack length (Cm) can then be determined from the surviving 

indentations, and using the critical crack stress (a,,,), the fracture toughness can be 

determined from: 

Klc = A omkt + B (51 ) 

1 

where A = 2.02 and b = -0.68 MPa-ms.99 The calibration curve in Figure 28 was used 

to determine A and 8. As the figure shows, the Klc values for a wide variety of 

materials fit the equation. Major benefits of the technique are: (1) it can be used to 

measure Klc at elevated temperatures and (2) fracture strengths are obtained at the 

same time. 

$ 
51 

Figure 27. indentation configuration for technique used to determine fracture 

toughness (Ref. 99). 

OO 0 1 2 3 4 

Figure 28. Calibration curve for fracture toughness method of Cook and Lawn 

(Ref. 99). 
52 

3. EXPERIMENTAL PROCEDURES 

3.1 Compositions and Batching 

The literature search suggested that compositions with 60 mol% or more hafnia 

and moderate concentrations of zirconia and titania would give the best thermal 

expansion, thermal shock resistance, and strengths possible in the system. On the 

basis of this knowledge, the following six compositions (in mol%) were selected for the 

preliminary investigation: 

1. 60% Hf02,35% ZQ, 05% Ti02 

2. 60% Hf02,30% ZQ, 10% Ti02 

3. 60% Hf02,25% ZQ, 15% Ti02 

4. 60% Hf02,20% ZQ, 20% Ti02 

5. 60% Hf02, 15% Zr02, 25% IO2 

6. 60% HfO2, 10% ZrO2,30% Ti02 

Early experiments with materials prepared with these compositions allowed the 

last two compositions to be eliminated from consideration. The remaining 

compositions exhibited the best sintering characteristics and offered the highest 

potential of delivering the properties sought. For the compositions exceeding 20 mol% 

titania, sintering was seriously impaired. 

The starting powders were high purity hafniaa, zirconiab, and titaniac. The 

impurity content of each*is given in Appendix 1. Six hundred gram batches of each 

composition were mixed by milling in plastic containers for one hour. Subsequently, 

a Teledyne Wah Chang Albany, Albany, OR 97321 

Magnesium Elektron, Inc., Fleming, N. J. 08822 

C J. T. Baker Chemical Company, Phillipsburg, N. J. 08865 
53 

three processing techniques were performed on each composition in order to make 

bar samples. 

3.2 Sample Preparation 

The three techniques used to prepare samples are described below. 

3.2.1 Reaction-Sintered Powders (TI) 

Two weight percent Polytran FSd in distilled water was added as binder to four 

hundred grams of the mixed as-received high purity powder. The powder was dried 

and sieved through a 200 mesh screen. Pellets 2.54 cm in diameter were then 

pressed uniaxially in a Denison pres9 to 12,000 psi (82 MPa), reaction sintered to 

1700OC for 5 hours in a Remey gas furnacef, crushed, milled in a steel ball mill, 

chemically treated with "03 and HCI to remove metal contamination introduced by 

ball milling, then further ball milled using Zr02 balls for 48 hours. Two hundred grams 

of the reaction-sintered powder were then pressed uniaxially in a tungsten carbide 

lined dies using a Carver pressh into the form of 5.2 cm x 0.6 cm x 0.3 cm bars at 

13,000 psi (87 MPa). Four weight percent Carbowaxi dispersed in acetone was used 

as the binder. The bars were then pressed isostatically to 30,000 psi (207 MPa) in 

rubber bladders using an Autoclave press.i 

The bars were cut in half using a stainless steel razor blade, taken to 5OOOC for 

4 hour to burn out the binder, then fired to 220OOC for 4 hours in a Brew vacuum 

t 

~~~~ ~ 

Owes Laboratories, Inc., chigo, IL 

e The Denison Engineering Company, Columbus, OH 

Remey, manufacturer unknown 

9 Staats and Staats, IN., Irwin, PA 15642 

Fred S. Carver, Inc., Summit, N. J. 

Union Carbii Corporation, N. Y., N. Y. 10017 

J Autoclave Engineers, Inc., Erie, PA 
54 

furnacek at 10-6 torr. A 6 pm diamond finish was put on three sides, f pm finish on the 

top surface, and the edges were bevelled. Three indents were made in most of the 

bars using a Vickers microhardness indenter with a load of 40 Ibs (178 N). Several 

bars were set aside in order to study the reduced state. The rest of the bars were 

oxidized in air at 17OO0C for 12 hours. 

3.2.2 &-Received Powders (T2) 

Four weight percent Carbowax dispersed in acetone was added to 200 grams 

of the mixed as-received powders. The powders were dried, sieved through a 200 

mesh screen, and pressed uniaxially into 5.2 cm x 0.6 cm x 0.3 cm bars at 13.000 psi 

(87 MPa). The bars were pressed isostatically to 30,000 psi (207 MPa), cut in half, 

presintered to 17OO0C in air for 24 hours, then vacuum fired to 2200°C for 4 hours. 

Most of the bars were subsequently oxidized in air at 1200°C for 6 hours. All bars 

were polished and bevelled. Bars used in the fracture studies were indented. 

3.2.3 Plasma-Sprayed Powders (T3) 

Two hundred grams of reaction sintered powder were passed through the high 

temperature plasma (H2 and N2 gas mixture) of an 80 kW METCO 7M plasma spray 

unitm operating at 50% power, then quenched into distilled water to effect rapid 

solidification. The powders were then calcined at 950°C in air for 4 hours and vacuum 

hot pressed" at 2000°C for 2 hours in graphite moldso under a pressure of 6.33 tons 

(27.8 MPa). The 5.2 cm diameter wafers were cut into bars 2.5 cm x 0.3 cm x 0.15 cm 

using a high speed diadond saw and polished. Several bars were oxidized in air at 

1200°C for 6 hours. 

Richard D. Brew and Company, Concord, N. H. 03301 

METCO, Inc., Westbury, L-I, N. Y. 

GVA Corporation / Vacuum Industries Division, Somenrille, MA 

O ATJ grade carbon, Union Carbide Products Division, Cleveland, OH 
55 

3.3 Property Measurements 

Three properties were investigated: (1 ) electrical conductivity, (2) thermal 

expansion, and (3) fracture strength and toughness. Each of these properties is 

important in understanding the effect that titania has in the HZT system and is pertinent 

to applications in heat exchangers and electric heaters. 

I 3.3.1 Electrical Conductivity 

Electrical conductivity measurements were made on bars prepared from asreceived high purity powders. PlatinumP electrodes were painted 0.5 cm apart and 

allowed to dry. Bars approximately 2.5 cm x 0.25 cm x 0.15 cm were mounted in the 

experimental apparatus shown in Figure 29 by wrapping the Pt-30%Rhq leads around 

the electroded ends. The apparatus was then heated to 1700OC. Resistance 

measurements were made while the furnace self-cooled (20 to 40°C/min) using a 

Wayne Kerr Automatic LCR meterr at 1OkHr. The electrical resistance for each 

composition was measured from 170OOC down to approximately 4OOOC in air 

(PO, = 21 kPa), and in premixed gassess of 1 .O% O2 in Ar (Po2= 10 kPa) and 9.55% CO 

in C02 (PO, = 10-2 to 10”‘ kPa). The electrical conductivity was determined using the 

I equation : 

O=(R +)-’ 

where 0 is the total conductivity, R is the measured resistance, A is the bar cross 

sectional area, and I is the distance between electrodes. 

P Electro-science Laboratories, Inc., PI paste No. 5545, King of Pnrssia, PA 19406 

q Electro-science Laboratories, Inc., King of Pnrssia, PA 19406 

Wayne Kerr LCR meter 4225, Woburn, MA 01801 

Air Products and Chemicals, Inc., Tamaqua, PA 18252 
Thermocouple Leads 

Gas Inlet 

Ai r-Ti g h t 

Brass Cap 

Alumina Tube 

Sample 

56 

Pt-30% Rh Leads 

Alumina Tube 

Closed-End 

A1 umi na Tube 

Thermocouple 

HOT ZONE 

Figure 29. Schematic diagram of experimental apparatus used in obtaining the 

electrical resistance of the HZT bars. 
57 

3.3.2 Thermal Expansion 

Thermal expansion measurements were made using the standard dilatometer 

set-up shown in Figure 30. The set-up consisted of a Theta 18 inch alumina specimen 

holder assemblyt mounted to an Orton di1atometer.U Bar samples approximately 

2.5 cm x 0.3 cm x 0.2 cm prepared from as-received powders were mounted vertically 

in the sample holder and upon expansion pushed the center rod up. The transducer 

in the dilatometer, which sensed the rod movement, sent a signal to a Daytronic LVDT 

conditi0ner.v A chart recorder then plotted the sample expansion. The sample holder 

was inserted through the top of a Teresco furnaceY Measurements were made in air 

between room temperature and 1700OC. The heating and cooling rates were 

2OC/min. The apparatus was calibrated by measuring the expansion of a known 

standard; namely, a single crystal sapphire having similar dimensions to the bars. 

Appendix II gives the thermal expansion of the standard. Each sample was cycled 

once to 1700OC prior to being measured. The thermal expansion for each composition 

was determined by adding the expansion of the apparatus determined from the 

standard calibration run to the measured expansion of the bars. 

3.3.3 Fracture Strength and Toughness 

Fracture strength measurements were made on bar specimens prepared 

from reaction-sintered (TI), as-received (T2), and plasma-sprayed (T3) powders. Either 

A1203 or Sic four-point bend load fixtures having inner spans of 6.5 mm and outer 

spans of 19 mm were used. Several bars of each composition in both the oxidized 

and reduced state were roaded using either a floor model Instronx or the load cell 

Theta Industries, Inc., Port Washington, N. Y. 11050 

Orton, Columbus, OH 43201 

Daytronic LVDT Conditioner Model 3130, Miamisburg, OH 45342 

Teresco, Champaign, IL 61820 

lnstron Model 1125, Canton, MA 02021 
58 

A1 umi ne 

Push 

Rod 

Alumina 

Specimen 

Holder 

Sampl 

Orton 

Dila tome ter 

1 x-y Chart Recorder 

rcerl 

I I U II 

A LVDT 

Conditioner 

'" I 

I 

1 

Thermocouple 

HOT ZONE 

* 

Figure 30. Schematic diagram of the dilatometer set-up used to determine the 

thermal expansion of the HZT bars. 
59 

I apparatus shown in Figure 31. Both were equipped with furnaces for hightemperature measurements. The loading rate was about 0.5 cm/min. Room 

temperature measurements were made on both nonindented and triple Vickers I 

I 

indented bars. Measurements at 1000°C and 140OOC were made only on indented 

bars. The fracture toughness was also determined for indented bars at room 

temperature, 1 OOO°C, and 14OOOC. Surviving crack length measurements were made 

using an optical microscope. 

I 

Figure 31. Schematic diagram of fracture load cell apparatus used in fracture 

strength and toughness study. 
60 

3.4 Analyses 

Each of the four compositions prepared using the three different techniques was 

analyzed for starting particle size, density, sintered microstructure, and phases 

present. 

3.4.1 Particle Size and Shape 

Particle size analyses were performed at IMC in Terre Haute, IN using a 

Microtrac Particle Size Ana1yzer.Y The Microtrac uses optical and laser light scattering 

to determine particle sizes. An IS1 DS-130 SEMZ equipped with an EDX was used to 

examine the particle shape and elemental make-up. 

3.4.2 Densities, Mass Changes, and Length Changes 

Densities were obtained by using an ethyl alcohol displacement technique. 

The dry mass and the suspended masses were determined using a Mettler balance.aa 

Mass changes were determined by measuring the mass before and after oxidation. 

Lengths of the green, fired and oxidized bars were also measured. 

3.4.3 Sintered Microstructure 

The sintered microstructure was examined using both optical and scanning 

electron microscopy. The average grain size was determined by the linear intercept 

method. 

Y Leeds and Northrup, North Wales, PA 19454 

aa Mettler Instrument Corp., Hightstown, N. J. 

International Scientific Instruments, Santa Clara, CA 
61 

3.4.4 Phases I 

X-ray analyses were performed using a Debye-Scherrer powder cameraab and 

l a Philips diffractometer.ac Monochromatic 40 kV, 30 mA copper radiation was used 

I for phase analyses. Ni-filtered 35 kV, 15 mA copper Ka radiation with a 11.4 cm I 

I diameter camera was used to determine accurate lattice parameters. 

ab Norelco Type 17010200, Philips Electronic Instruments, Mount Vernon, N. Y. 

ac Philips APD 3520, Philips Electronic Instruments, Inc., Nahwah, N. J. 07430 
62 

4. RESULTS AND DISCUSSION 

The effects of Ti02 on selected properties of Hf02-rich mixed oxides in the HZT 

system were examined in detail. Analyses of the starting powders, sintered 

microstructures, densities, and final phases are presented first, followed by electrical 

conductivity, thermal expansion, and fracture strength and toughness studies. For 

ease of discussion, the following nomenclature is used for sample identification. Each 

group of samples is designated by composition (in mol% Ti02), fabrication technique 

(e.g., made from: (1) reaction-sintered powders, (2) as-received powders, or (3) 

plasma-sprayed powders), and oxidation state (e.g., "ad" represents highly anion 

deficient or "ox," oxidized). For example, 5Tlox is an oxidized composition of 5 mol% 

TO2 made from reaction-sintered powders. 

4.1 Powder Analyses 

The average particle size, shape, and relative composition of the starting 

powders were determined for each composition and fabrication technique. Table 7 

summarizes the particle size analyses. The analyses showed that the reactionsintered (TI) powders consisted of fine particles having an average size of 2 to 4 pm. 

The plasma-sprayed (T3) and as-received (T2) powders consisted of coarse particles 

having an average size of 16 to 23 pm. Figures 32-43 show scanning electron 

micrographs of the twelve starting powders. The T1 and T2 powders had very rough 

jagged surfaces, whereas the T3 powders were round and smooth. EDX analyses 

showed qualitatively that the three sets of powders had the correct proportions of Hf02, 

Zr02, and Ti02 across the compositional series and that the T1 and T3 particles were 

homogeneous mixtures. No contaminations were detected in any of the powders. 

c 

4.2 Densities, Length Changes, and Mass Changes 

The densities of samples at various stages of fabrication were determined for 
0R.EGINA.E PAGE IS 

OE POOR QUALITY 63 

Table 7. Particle Size Analyses of Starting Powders for HZT Compositions. 

Composition Average Particle Size (pm) 

HfO2 22.83 

za2 1.82 

Ti02 1.20 

5T3 

1 OT3 

20T3 

15T3 

4.1 4 

2.90 

2.57 

3.84 

26.94 

18.29 

20.53 

26.93 

28.53 

16.08 

15.72 

17.43 

Figure 32. Scanning electron micrograph of reaction-sintered powder having 

5 mol% TiO, (5T1). 
.ORIGINAL PAGE IS 

.OF POOR QUALITY 

65 

Figure 35. Scanning electron micrograph of reaction-sintered powder having 

20 mol% Ti02 (2011). 

Figure 36. Scanning electron micrograph of as-received powder having 

5 mol% TiO, (5T2). 

PRECEDLNG PAGE BLANK NOT FILMED 
OFLE(mJN; PAGE IS 

OX;: POOR QUALITX 66 

Figure 37. Scanning electron micrograph of as-received powder having 

10 mol% Ti02 (1 OT2). 

Figure 38. Scanning electron micrograph of as-received powder having 

15 mol% TiO, (1 5T2). 
67 

Figure 39. Scanning electron micrograph of as-received powder having 

20 mol% Ti02 (20T2). 

Figure 40. Scanning electron micrograph of plasma-sprayed powder having 

5 mol% TiO, (5T3). 
68 

Figure 41. Scanning electron micrograph of plasma-sprayed powder having 

10 mol% Ti02 (1 OT3). 

Figure 42. Scanning electron micrograph of plasma-sprayed powder having 

15 mol% TiO, (1 5T3). 
69 

Figure 43. Scanning electron micrograph of plasma-sprayed powder having 

20 mol% Ti02 (20T3). 

each composition and fabrication technique. Table 8 summarizes the results. 

Theoretical mix densities were determined using the rule of mixtures (Le., for each 

composition, the densities of Hf02, Zr02, and Ti02 were added in compositional weight 

proportions to determine the density of the mixed materials). Theoretical x-ray 

densities were determined using lattice parameter measurements of fired materials (in 

the monoclinic form) and assuming that the fluorite unit cell consisted of atoms with 

compositional proportions. 

Isostatically pressed bars in the green state, were found to be approximately 

60% of the theoretical dewity for both the TI and T2 series. Samples that were fired to 

17OOOC in air achieved only 85% theoretical density. Increasing the firing temperature 

to 2200°C increased the densities of the reduced materials to roughly 95% and 93% 

theoretical for the T1 and T2 series, respectively. Upon oxidation of these materials at 

I 
70 

I 

Table 8. Densities, Mass Changes, and Length Changes of HiT Compositions. 

Comp. Theoretical Theoretical Green Fired Fired Mass Length 

Densi (ad) Densi (ox) Change Change 

(g/cm3) (g/m3) (g/m3) (g/cm 3 (yo) (g/cm 3 (yo wt gain) (96 expan.) 

Mix Density X-ray Density Density 

5.34 64 8.14 97 8.09 97 0.47 0.33 

8.54 8.37 4.98 58 7.86 94 7.75 93 0.45 0.46 

5T1 

5% 

5T3 7.97 95 7.99 95 0.28 0.20 

1 OT1 5.26 63 8.06 96 7.99 95 0.61 0.51 

1 OT2 8.55 8.38 4.92 58 7.81 93 7.70 92 0.62 0.57 

1 OT3 8.40 100 8.41 100 0.35 0.31 

f5T1 521 63 7.85 94 7.75 93 0.69 1.40 

1 ST3 8.31 100 8.33 100 0.46 0.54 

5.05 60 7.76 92 733 a7 0.71 2.40 

8.56 8.45 4.76 56 7.57 90 7.42 88 0.74 2.40 

8.37 99 8.45 100 0.55 0.71 

1 5T2 8.55 8.33 4.83 56 7.73 93 7.52 93 0.72 1 .a0 

=TI 

=T2 

2OT3 

1200OC in air, the bars underwent a weight gain and lengthening that increased with 

the concentration of Ti02. A slight reduction in the densities of the oxidized materials 

resulted. Bars that were subsequently oxidized to 170OOC in air suffered a further 

decrease in density reducing to approximately 92% theoretical. The density reduction 

that accompanied the oxidation in the T1 and T2 series materials resulted in part from 

phase changes (e.g., the formation of lower density titanates). Phase analyses of the 

various oxidation states of the T1 and T2 series will be discussed shortly. 

For the T3 series, densities of the vacuum hot pressed samples in the reduced 

state approached theoretical density. Upon oxidation of the samples, the densities 

became 100% of the theoretical values. In addition, lower linear expansions and 

weight gains were observed during oxidation of the T3 compositions as compared with 

the T1 and T2 series. These reduced weight and length changes that accompanied 

oxidation can be attributed to the absence of the cubic phases in the reduced state 

and the titanate phases in the oxidized state of the T3 series. Phase analyses of the T3 

series will be discussed in detail in Section 4.4. (Table 2 shows that the unit cell 
71 

volumes of the titanate and cubic phases are lower than the volumes of the monoclinic 

phases. This volume differential leads to increased volume changes upon oxidation of 

the materials in the T, and T2 series.) The increased densities in the T3 series can 

also be attributed to the absence of the titanate phase in the oxidized materials. 

(Table 2 shows that the titanate phases have lower densities than the monoclinic 

phases.) 

4.3 Microstructure Analyses 

The microstructures of the fired bars were analyzed using both optical and 

scanning electron microscopy. From optical microscopy, large regions (approximately 

16 to 3000 pm) were observed on the polished surfaces of both the reduced and 

oxidized bars. Scanning electron microscopy showed that these large regions 

consisted of very fine grains (-1 pm). A summary of the region and grain sizes is given 

in Table 9. 

Scanning electron micrographs of oxidized, polished surfaces of several 

compositions in the TI, T2, and T3 series are given in Figures 44-52. The micrographs 

(taken at various magnifications) show the relation between the large regions and the 

fine grains. 

A close examination of the region boundaries in the SEM micrographs suggests 

that microcracking occurs in the HZT system. As was pointed out in Section 2.2.2.1, 

microcracking has been observed in the Hf02-Ti02 system. A similar situation would 

be expected in the HZT system. The large (microcrack) regions in the T1 compositions 

dramatically increase in size in going across the series, whereas in the T2 and T3 

series, little change in size is observed. On the other hand, the grain size in the T1 

series decreases, whereas in the T2 and T3 series, it remains constant with 

composition. These observations suggest that the size of the microcracked regions is 
72 

Table 9. Microstructure Analyses of Fired HZT Compositions. 

Composition Average Grain Size (pm) Average Region Size (pm) 

(within Regions) 

5T1 1.3 200 

1 OT1 1.8 160 

15T1 0.8 1000 

20T1 0.7 3000 

5T3 

1 OT3 

20T3 

1 5T3 

1 .o 

1 .o 

1 .o 

0.9 

-1 .o 

-1 .o 

-1 .o 

-1 .o 

27 

34 

36 

32 

31 

20 

17 

16 

inversely proportional to the grain size. Dole et. al.52 found similar effects in their study 

of microcracking in Hf02. 

An examination of the grains in the T1 and T2 series micrographs shows that a 

degree of liquid phase sintering is occurring in the higher Ti02 compositions. The 

possibility that a liquid phase might be present was observed during firing. Several 

bars of the high Ti02 compositions fused together when heated to 2200°C. From the 

phase diagram of the Hf02-Ti02 system (see Figure 7a), a solid plus liquid two-phase 

region exists at approximately 220OOC for compositions above 15 mol% Ti02. A 

similar region apparently exists in the HZT system. 

Figures 44-51 sho*w that a high degree of porosity is present in the sintered 

bodies of the T1 and T2 series. The porosity was observed to concentrate at the grain 

boundaries. Scanning electron micrographs of polished surfaces of 1 OJ30x, 1 5J30x, 

and 20T30x showed the material to be highly dense with very little porosity. Some 

porosity, however, remained in the 5T$X composition after hot pressing. Figure 52 
73 

Figure 44. Scanning electron micrographs of a thermally etched, polished surface 

of a fired 5Tlox bar at (a) high magnification and (b) low magnification. 
74 

Figure 45. Scanning electron micrograph of a thermally etched, polished surface 

of a fired 10710~ bar. 
75 

Figure 46. Scanning electron micrographs of a thermally etched, polished surface 

of a fired 15Tlox bar at (a) high magnification and (b) low magnification. 
76 

Figure 47. Scanning electron micrograph of a thermally etched, polished surface 

of a fired 20Tlox bar. 

Figure 48. Scanning electron micrograph of a thermally etched, polished surface 

of a fired 5T20x bar. 
77 

Figure 49. Scanning electron micrograph of a thermally etched, polished surface 

of a fired 10T20x bar. 

Figure 50. Scanning electron micrograph of a thermally etched, polished surface 

of a fired 15T20x bar. 
78 

Figure 51. Scanning electron micrograph of a thermally etched, polished surface 

of a fired 20T20x bar. 
Figure 52. Scanning electron micrographs of thermally etched, polished 

surface of fired (a) 5T30x and (b) 20T30x bars. 

79 

! 
80 

shows that a degree of the original round particle nature that existed prior to pressing 

was retained. 

Qualitative EDX indicated that all compositions: (1 ) had the correct proportions 

of Hf, Zr, and Ti, (2) were homogeneous, and (3) had no apparent contamination. 

Exceptions to statement (2) above were the 20 mol% Ti02 compositions that had 

titanates. Second phase titanate grains were found that had higher concentrations of 

Ti than surrounding grains. 

X-ray diffractometry suggested that a degree of preferred orientation was 

present within the microcrack region in many of the compositions. Diffractometry 

traces of bars having large microcrack regions tended to have peaks with different 

relative heights depending upon the orientation of the bar with respect to the x-ray 

beam. This, coupled with indications of preferential polishing directions, seems to 

imply that the grains within a microcrack region tend to preferentially align. 

4.4 Phase Analyses 

Qualitative x-ray analyses were performed on each of the four compositions in 

the TI, T2, and T3 series. A variety of phases resulted depending on the composition, 

the fabrication technique, the oxidation state, and the oxidation time and temperature. 

Table 10 gives a summary of the phase analyses. 

4.4.1 Reduced or Highly Anion Deficient State 

All compositions prepared in each of the three techniques were black colored 

after firing in vacuum. This state is referred to as the reduced or highly anion-deficient 

state. In the reduced state, a variation in the stoichiometric ratio and a weight loss 

occur due to oxygen release. The black coloration results from entrapped electrons 

which fill the vacated oxygen sites.43 A modification of the structure can occur 

depending on the degree of anion deficiency (or departure from stoichiometric 

* 
81 

Table 10. X-ray Phase Analyses of the HZT Compositions. 

Comp./Technique ad OX (6-12 hE) ox (4 hr) ox 

(at 1200°C) (at 1300°C) (at 1700OC) 

~TI M M M 

5T2 M M M M 

5% M M 

L 

1 OT1 M+C M 

1 OT2 M+C M 

1 OTS M M 

M 

M M 

15T1 C+Mt M+Ot M 

1 5T2 C+Mt M+Ot M Mi+&f+Ot 

1 5T3 M+C+ M 

20T1 C+Mt M+O M 

20T2 C M+O M MI +M2+0 

20T3 M+C M 

M is the monoclinic phase, MI and M2 refer to two different monoclinic phases in 

solution, C is the cubic phase, and 0 is the orthorhombic titanate phase. A t indicates 

that only trace amounts were observed. 

proportions). For stoichiometric Hf02 and Zr02, the percentages of oxygen are 15.2 

and 26.0 wt%, respectively. Two weight percent oxygen loss can be accommodated 

in Hf02 at 220OOC while still maintaining a solid solution of Hf02-,.100 For Zr02, up to 

three weight percent oxygen loss can be accommodated.101 From the Zr02-rich side 

of the Zr-0 phase equilibrium diagram, at 2200°C, tetragonal ZrO2, is stable up to 

0.19 wt% oxygen loss. From 0.19 to 0.48%, cubic and tetragonal ZrOe,, are stable. 

And from 0.48 to 3.0 wt%aloss, cubic ZrO2, is the stable phase.101 A similar situation is 

found in Hf02,x.100 

In this investigation, the stoichiometric weight percentages of oxygen in the 5, 

10, 15, and 20 mol% Ti02 compositions were 18.5, 18.7, 18.9, and 19.2, respectively. 

For the T1 and T2 compositional series, the weight loss range of the reduced state was 
82 

from 0.45 to 0.74%. For the T3 series, the range was from only 0.28 to 0.55%. The 

lower weight loss in the T3 series is probably due to either the lower firing temperature 

and time (i.e., 20OO0C for 2 hours) or to a difference associated with vacuum hot 

pressing as opposed to sintering in a vacuum furnace. 

A comparison of the observed phases cited in Table 10 with the weight losses 

given in Table 8 suggests that a correlation exists between the degree of oxygen 

weight loss and the resultant phases. For 0 to 0.47 wt% oxygen loss, only the 

monoclinic phase was seen. For 0.47 to 0.74%, the monoclinic and cubic phases 

were observed. Above 0.74%, only the cubic phase existed. These phases 

correspond to a combination of high and low temperature states and follow similar 

trends to those which occur in the Zr-0 system. 

Since fairly rapid cooling rates were employed (20 to 4O0C/min.) in all the 

fabrication techniques, the cubic high temperature phase in the reduced HZT 

, compositions may have been a result of quenching effects. Quenching is known to 

stabilize the high temperature cubic phase in anion deficient 2102.101 In the HZT 

system, the anion deficiency as well as the presence of titanium act to stabilize the 

cubic phase upon cooling. (At low oxygen partial pressures (as occur in vacuum 

sintering), titanium ions enter the lattice structure substitutionally for hafnium and 

zirconium ions. The variability of the titanium ion valence leads to an increased 

degree of anion deficiency in the reduced compositions which becomes even greater 

as increasing amounts of titanium ions are added. A more relaxed stable cubic 

structure is produced as a result of the anion deficiency and the presence of titanium. 

Upon cooling of these materials, the amount of structural reorientation necessary to 

change the high temperature effects of titanium are too great, and thus the cubic phase 

is retained to room temperature.) These occurrences explain the increased amounts 

of cubic phase that are observed in the reduced materials in the HZT system that have 

t 
83 

high Ti02 concentrations. The high temperature tetragonal phase, on the other hand, 

reverts to monoclinic symmetry upon cooling and is not stabilized by titanium thus 

explaining the absence of the tetragonal phase in the reduced state after cooling. 

The unit cell lattice parameters of the monoclinic phase in the reduced HZT 

compositions were approximately: a = 5.009 A; b = 5.134 A; c = 5.272 A; and 

@ = 98.9'. A comparison with the larger parameters of pure monoclinic zirconia and 

hafnia (refer to Table 2) suggests that either the smaller titanium ions are entering the 

lattice or that anion deficiency is reducing the size of the unit cell. The lattice 

parameters of the cubic phases in all of the reduced HZT compositions were 

ca. 5.055 A. This value is lower than that of pure cubic Zr02 (5.07 A) suggesting that 

the cubic unit cell in the HZT compositions is also modified by titanium and/or anion 

deficiency. 

4.4.2 Oxidized State 

The oxidized state was obtained after the reduced materials were annealed in 

air. Depending on the annealing temperature, the bars turned either cream colored 

(1 200OC) or white (1 7OOOC). Also, different phases were observed. 

X-ray analyses showed that materials in the TI and T2 series that were oxidized 

at 1200OC were monoclinic solid solutions for compositions with less than 15 mol% 

Ti02. For compositions with more than 15 mol% Ti02, a titanate phase was observed 

in addition to the monoclinic phase. As Table 10 shows, high Ti02 compositions in the 

T2 series that were annealed at 170OOC contained an additional monoclinic phase, 

whereas identical compositions annealed at 12OO0C did not. (A similar situation 

occurs in the Hf02-Ti02 system.18) Table 11 gives the d-spacings which distinguish 

the two monoclinic phases and the orthorhombic titanate phase in the 15T2ox and 

20T20x materials. The significance of the second monoclinic phase was discussed in 

Section 2.1 -2.2. 

). 
84 

Table 11. X-ray d-spacings from Powder Patterns of T20x Samples Oxidized at 

1700°C. 

Composition d-i 1 1 &(A) d-i 1 1 Ml(A) dl 1 oJA) d~ 1 1 Ml+ &(A) 

(in mol% Ti02) 

15 3.123 3.080 2.952 2.808 

20 3.1 27 3.081 2.963 2.81 0 

For all compositions in the T3 series, only the monoclinic phase was observed 

after oxidation. Even the 20 mol% Ti02 composition showed no presence of the 

titanate phase. This is in agreement with the work of Brown et. al.4 and Schroeder.3 

Powders in the TI and T2 series that were taken to 13OOOC for 4 hour on a hot stage 

x-ray diffractometer and cooled were also observed to be only monoclinic. Both the T3 

series oxidation effects and the T1 and T2 oxidation effects at 1300OC suggest that the 

suppression of the titanate phase is an oxidation related phenomenon. 

X-ray powder diffraction patterns were taken of each composition in the 

monoclinic phase. The results are given in Table 12. A decrease in parameters with 

increasing Ti02 was observed. This confirms that the titanium is entering the unit cell; 

however, whether or not the titanium takes an interstitial or substitutional position is not 

apparent. 

Table 12. X-ray Lattice Parameters of Monoclinic HZT Compositions. 

P 

(degrees) (& 

Composition 

(in mol% Ti02) I 

5 5.109 5.158 5.288 99.2 137.6 

10 5.087 5.127 5.273 99.0 135.8 

15* 5.071 5.104 5.258 98.8 134.6 

20* 5.030 5.050 5.240 . 98.6 131.3 

*Not stable form (i.e., 100% monoclinic phase with no titanate). 
4.5 Electrical Conductivity 

Electrical resistance measurements were made on bars in the T2 compositional 

series at temperatures over the range 4OOOC to 17OOOC in air and in gas mixtures of 

1% oxygen in argon and 9.55% CO in C02. Measurements were made to determine 

the effects of titania on the conductivity of compositions in the HZT system as a function 

of temperature and oxygen partial pressure. The T2 series was chosen for 

examination since testable materials were produced for each composition within the 

series. Similar electrical behavior, however, would be expected for identical 

compositions in the TI and T3 series as well. 

The electrical conductivity as a function of temperature is shown in Figure 53 for 

the four compositions. Figure 54 shows the same data plotted so that the pressure 

dependence for each composition can be seen more clearly. Several observations 

can be made from the electrical conductivity data: (1) At high temperatures and low 

PO,, the conductivity increases with increasing Ti02 concentration. (2) At low 

temperatures and high PO,, the conductivity decreases with increasing Ti02, while at 

low PO,, it increases. The observed conductivity trends coupled with the phase 

analyses suggest that titanium is incorporated (1) interstitially at low temperatures and 

high PO, and (2) substitutionally with a valency change at low PO,. 

The probable defects in the HZT system are summarized in Table 13. The 

I 

defect formation reactions can be written as: 

F 

KZ -bL 

Kt 

K4 

Ks 

K, 

K7 

00 * ojn+v; 

Oo t *02 + Vow + 2e' 

ev * e' + h" 

s * VM"" + 2vo"" 

FM * F;"" + VM"" 

FM FM' + h' 

FM' * FM" + h" I 
86 

-1 I 

5% 

10% 

15% 

0 20% 

0.4 0.6 0.8 1.0 1.2 1.4 1.6 

1 OOOn (1 PK) 

(a) 

5% 

+ 1OYo 

15% 

0 20% 

0.4 0.6 0.8 1 .o 1.2 1.4 

1 OOOn (1 PK) 

-1 7 

0.8 1.2 1.6 2.0 

1 OOOn (1 PK) 

Figure 53. The electrical conductivity of four compositions in the HZT system 

measured as a function of temperature in (a) air, (b) 1 !Ao O2 in Ar, and 

(c) 9.55% CO in C02. The amount of TiO, in mol% is indicated on the 

right. 
-2 - 

3- 

-4- 

-5 - 

-6 - 

-2 - 

3- 

-4- 

-5 - 

-6- 

-7 - 

5 moWo Titania 

Q Air 

l%Oxygen 

e 

e 

0 

-5 - 

-6 - 

0.4 0.6 0.8 1.0 1.2 1.4 1.6 

1OOO/r (1PK) 

(a) 

- 

-an 

'"my *a a 

** a 9 ,a I. 

0.4 0.6 0.8 1.0 1.2 1.4 1.6 1.8 

10OOlT (1PK) 

n 

E 

.c 8 

E 

1 

n 

Y 

c 

Io 

3 

g 

8 

J 

a7 

-1 7 

J 

-2 - 

3- 

-4- 

-5 - 

-6- 

10 mol% Titania 

Air 

a 9.55% co 

1VoOxygen 

10 mol% Titania 

Air 

a 9.55% co 

1VoOxygen 

+a 

%g.' 8 

m 6 

0.4 0.6 0.8 1.0 1.2 1.4 1.6 

1ooo/r (1PK) 

(b) 

-I I 

20 mol% Titania 

b a Air 

L l%Oxygen 

I a 9.55%co 

-8' . I. 1.1 

0.4 0.8 1.2 1.6 2.0 

1 OOO/T (1 PK) 

Figure 54. The electrical conductivity measured as a function of temperature in air, 

1 % O2 in Ar, and 9.55% CO in C02 for compositions in the HZT system 

having (a) 5, (b) 10, (c) 15, and (d) 20 mol% Ti02. 
88 

Table 13. 

Defect Description 

00 Oxygen in Oxygen Site 

Oi" Interstitially Incorporated Oxygen Ion v0- Doubly-Ionized Oxygen Vacancy 

VM1,'* Tetravalent Cation Vacancy 

FM Substitutionally Incorporated Ti4+ Cation on a Tetravalent Cation Site 

FM, Substitutionally Incorporated Ti3+ Cation on a Tetravalent Cation Site 

FM" Substitutionally Incorporated Ti2+ Cation on a Tetravalent Cation Site 

Interstitially Incorporated Ti4+ Cation 

S Schottky Pair 

ev Electron in Valence Band 

e' Electron in Conduction Band 

h" Hole in Valence Band 

Probable Defects in the HZT System. 

Flwhere the reactions in Equations (56)-(59) represent the effect of interstitially and 

substitutionally incorporating titanium into the fluorite structure. The reactions in 

Equations (53)-(55) are equivalent to those given in Equations (2)-(4) for pure fluorite 

materials. A more direct relation between hole and electron formation, however, is 

given by Equation (55). 

The mass action law for each becomes, respectively: 

K1 = Io;] P0O-I (60) 

(61 1 

1 

K2 = n* PQ~ 0/0"3 

where [ ] indicates defect concentrations, and n and p are the electron and hole 

concentrations, respectively. 
89 

Conservation of titanium requires that the sum of the concentratiohs of 

interstitially and substitutionally incorporated titanium is equal to the total amount of 

titanium added. Also, electroneutrality requires that the concentration of all positive 

charge carriers equal that of all negative charge carriers. These two conditions are 

then: 

[Til = [Fd [FMI [FMI + [FIT (67) 

n + 4[V~*"l+ [FM~ + ~[FM'*] + 2[Oi7 = p + 4[F<"] + 2p,"q (68) 

The defect concentrations as a function of PO, can then be determined for the 

following six regions. 

In Region I (very low PO,), the dominating electroneutrality condition is 

n = 2[V,'q, and the prevailing titanium conservation condition is [FM"] - [Ti), giving: 

1 

p = (gy Po," 

- 1 w,"] Po,'= 
90 

In Region II (low PO,), the electroneutrality condition is [Vo""] = [FM**] and the 

titanium conservation condition is [FM"] = vi], giving: 

D/o"l= VI 

K1 [O?] -- 

- [Til 

and the equations for [FM], [FM*], [FM**], and [Fr] are the same as those in Region 1. 

In Region 111 (medium Po2), the electroneutrality condition is [FM*'] = 2[F("""], and 

the titanium conservation condition requires that [FM**] =$[Ti] and [F("""] = $.[Ti], giving: 

and the equations for [FM], [FM*], [FM'*], and [Fro"] are the same as those in Regions I 

and II. 
91 

In Region IV (medium high PO,), the electroneutrality condition is [FM'] = 4[Fj7, 

and the titanium conservation condition is [FM] = vi], giving: 

and the equations for [FM], [FM'], [FM"], and [Fr] are the same as those in Regions I, ll 

and 111. 

In Region V (high PO,), the electroneutrality condition is p = [FM'] and the 

titanium conservation condition is [FM] = [Ti], giving: 

and the equations for [FM], [FM'], [FM'~, and [FI"'] are the same as those in Regions 1,11, 

111, and IV. 
92 

In Region VI (very high PO,), the electroneutrality condition is p = 4[VM1'"], and 

the titanium conservation condition is [FM] = [Ti], giving: 

and th quation 

Ill, IV and V. 

for [FM], [FM'], [FM"], and [Fr] are the same as those in Regi ns I, II 

The proposed Kroger-Vink diagram which shows the defect concentration 

dependence on the oxygen partial pressure is shown in Figure 55. The titanium 

concentration was chosen arbitrarily to be 10 mol%. Equilibrium constants (K) were 

chosen to give reasonable defect concentrations over the specified pressure regions. 

The diagram given in Figure 55 is one of a number of possibilities based on 

assumptions of equilibrium constants and titanium concentrations. In addition, the 

pressure ranges given in the diagram are only relative to actual pressure regions. 

The electrical conductivity measurements from this work were plotted as a 

function of PO, at various temperatures (Figure 56). (The PO, of 9.55% CO in C02 has 

the temperature dependence shown in Figure 57.) A negative PO, dependence is 

observed at high temperatures under low pressures for HZT compositions with high 

Ti02 concentrations. This negative slope is indicative of n-type conductivity. A positive 

slope is observed at low temperatures under high pressures for all compositions, and 

at high temperatures for the 5 mol% Ti02 composition. This positive slope is indicative 

of p-type conductivity. A general tendency to go from n-type to p-type conductivity 
93 

I NTRl NSlC DEFECTS 

f Y 

4 

30 

25 

20 

15 

10 

5 

0 

-5 - 

-60 -40 -20 0 20 40 60 

LOG OXYGEN PARTIAL PRESSURE 

EXTR I N SIC D E FECTS 

-60 -40 -20 0 20 40 60 

LOG OXYGEN PARTIAL PRESSURE 

t 

(b) 

Figure 55. Kr6ger-Vink diagram showing defect concentration dependence on 

PO, for KO2 doped M02 for (a) intrinsic and (b) extrinsic defects. 
94 

5 mol% Titania 

-2 - 

-3 - 

-4- 

-5 - 

10 mol% Titania ' 

I? "I 

-8 -6 -4 -2 0 2 4 6 -8 -6 -4 -2 0 2 4 6 

LOG OXYGEN PARTIAL PRESSURE (Pa) LOG OXYGEN PARTML PRESSURE (Pa) 

looO~C 

130OOC 

1400°C 

0 15oooc 

1600% 

15 mol% Titania 

8l-A-L-J J -6 

-8 -6 -4 -2 0 2 4 6 

LOG OXYGEN PARTIAL PRESSURE (Pa) 

i. 

(a 

20 mop/, Titania 

-8 -6 -4 -2 0 2 4 6 

LOG OXYGEN PARTIAL PRESSURE (Pa) 

Figure 56. Electrical conductivity dependence on PO, for compositions in the HZT 

system: (a) 5 mol% TiO,, (b) 10 mol% TO2, (c) 15 mol% TO2, and 

(d) 20 mol% TO,. 
95 

0: 

9.55% Carbon Monoxide/90.45% Carbon Dioxide 

.;a 

a_ lo t 

Figure 57. The temperature dependence of the oxygen partial pressure of 

9.55% CO in C02. 

occurs as oxygen partial pressure is increased. This tendency is suppressed as TO2 

is added. That is, at high pressures, titanium is incorporated interstitially reducing 

p-type conductivity with increasing Ti02 concentration. At low pressures, titanium is 

incorporated substitutionally with a valency change increasing n-type conductivity with 

increasing Ti02 concentration. The trends observed in the data support the proposed 

defect model for the HZT system. 

Since defect mobilities were not considered in the present model and since 

equilibrium constants were chosen rather than determined, detailed correlations 

between the conductivities predicted by the model and experimental values cannot be 

conclusively made. Furthermore, accurate determinations of the slopes in Figure 56 

are not possible due to the limited nature of the data. Conductivity measurements at 

several additional partial pressures need to be made in order to properly determine 

the slopes and assess the model. The primary purpose of plotting the data in 

Figure 56 was to show the general pressure dependence on conductivity that occurs 

in the HZT system and to correlate the dependence to a possible defect model. 

t 
96 

A comparison with the literature shows that unlike the Zr02-based electrolytes 

that are doped with CaO or MgO, the HZT compositions have a rather high electronic 

conductivity component. Worrell and WeppnerM found a similar effect in Ti02 doped 

Y203-stabilized Zr02. They also concluded that the high electronic conductivity was 

due to titanium entering into interstitial positions. They were unable, however, to 

determine whether the titanium valency changed leading to increased conductivity via 

electron hopping or if the oxygen vacancies were causing the increase in conductivity. 

4.6 Thermal Expansion 

Thermal expansion measurements were made on T20x bars from room 

temperature to 170OOC in air. Measurements were made to determine the effects of 

titania on the thermal expansion behavior in the HZT system. The T20x series was 

chosen for examination since all the compositions in the series produced testable 

bars. Extensive microcracking occurred in the Tlox series making them impossible to 

thermally cycle to 1700°C repeatedly without breaking. As will be explained shortly, 

low thermal expansion coefficients (< 4 x lO-S/OC) would be expected for these 

materials since microcracking was so prevalent. Compositions in the T30x series were 

also unable to survive repeated cycling to 170OOC and hence were not tested. Again, 

low thermal expansion coefficients would be expected based on literature findings for 

monoclinic materials in the HZT system.2-4 Since the testing apparatus was not set up 

for thermal expansion measurements in reducing environments, thermal expansion 

measurements of reduced HZT compositions were not made. A study of the thermal 

. 

expansion behavior of lhese materials would certainly be enlightening and is 

recommended for future studies of the HZT system. 

4.6.1 Thermal Expansion Coefficient 

The fractional thermal expansions (AL/L) versus temperature are shown in 
97 

Figures 58-61 for the four T20x compositions. The thermal expansion coefficient of 

each composition is given in Table 14 for both the monoclinic and tetragonal phases. 

A general decrease in the coefficient occurs with increasing Ti02 concentration. For 

the monoclinic phase, a range from 6.2 to 5.0 X 10-6/oC was observed for the 

compositional range 5 to 20 mol% TiO,. As pointed out in Table 4, the calculated 

linear thermal expansion coefficients determined from lattice parameter data for pure 

Hf02 and Zr02 in the monoclinic phase are 6.9 and 7.6 x lO-S/OC, respectively. The 

measured coefficients for compositions of Hf02-Zr02 solid solutions having about the 

same proportions as in this investigation are ca. 7.2 X 10-6/OC.21 The measured 

thermal expansion coefficient values in this study fall below each of these values 

implying that Ti02 is indeed reducing the thermal expansion. 

From studies of the Zr02-Ti02 and Hf02-Ti02 systems, the thermal expansion 

coefficient was also found to reduce with increasing TiO,. For 20 mol% Ti0, added to 

ZrO,, the reduction was 11%.19 For 25 mol% TiO, added to Hf02, the thermal 

expansion coefficient was reduced by 60%.56 In this study, the reduction was 'double 

the amount found in the Zr02-Ti02 system. As pointed out in Section 2.2.2.1, the 

titanate that forms in the Hf02-Ti02 system is the cause of the extreme reduction in the 

thermal expansion coefficient. The titanate that forms in the Zr02-Ti02 system does not 

reduce the thermal expansion coefficient by nearly as much. This suggests that the 

titanates that form in the HZT system have an intermediate effect on the the reduction 

in thermal expansion as compared with HfTi04 and ZrTi04. But indeed, the titanate 

does lower the thermal expansion coefficient as is seen by the large drop in Table 14 

in going from the 15 to 26 mol% compositions. 

A discrepancy exists between the measured coefficient of the 20 mol% TO2 

composition obtained in the present study (5.0 x lO-S/OC) and that found for an 

identical composition studied by Brown et. al. (3.1 x 1 O-S/OC).4 These anomalous 

results can be explained by a number of factors (e.g., fabrication and phase 
98 

0.010 I 1 

3 

a 

0 

Y 

z 

<n 

z 

X 

W 

2 

5 

3 

I 

I- 

-0.002 

0 400 800 1200 1600 2000 

TEMPERATURE (“C) 

+ Heating 

-0 Cooling 

Figure 58. Linear thermal expansion of the 5 mol% Ti02 composition in the 

HZT system. 

0.008 1 Omol% Titania 

0.006 

0.004 

0.002 

0.000 

~ + Heating 

’ -O Cooling 

-0.002 

0 400 800 1200 1600 2000 

TEMPERATURE (“C) 

Figure 59. Linear thermal expansion of the 10 mol% Ti02 composition in the 

HZT system. 

- 
99 

0.010 , 1 

3 

Y 

a 

z 

v) 

z 

X 

W 

0 

a n 

$ 

CT 

W 

I 

I15mol% Titania 0.008 

0.006 

0.004 

0.002 

-0.000 

-0.002 

0 400 800 1200 1600 2000 

TEMPERATURE (“C) 

Heating 

Cooling 

Figure 60. Linear thermal expansion of the 15 mol% TiOe composition in the 

HZT system. 

3 a 

Y 

0.006 

.0.004 

0.002 

-0.000 

-0.002 

TEMPERATURE (“c) 

Heating 

Cooling 

Figure 61. Linear thermal expansion of the 20 mol% Ti02 composition in the 

HZT system. 
100 

Table 14. Thermal Expansion Coefficients of the Monoclinic and Tetragonal Phases 

in the HZT Compositions. 

Composition 

(mol% TiO,) 

Thermal Expansion Coefficient (1 O$/OC) 

Monoclinic Phase Tetragonal Phase 

5 6.22 8.31 

10 6.1 2 8.29 

15 6.05 8.26 

20 5.03 7.96 

differences, temperature range differences, and measurement rate differences). 

Schroeder, in his work with the HZT system, found that the fabrication process 

had a critical effect on the reproducibility of measurements and that the reproducibility 

was rather p00r.2*3 As this study has already pointed out, phase differences can be 

expected for the same composition depending on the fabrication and oxidation 

process. in the work of Brown et. al.,4 the 20 mol% composition was fabricated by hotpressing disks from reaction-sintered powders and then sintering the disks at 20OO0C 

for 1 hour. In the present study, the fabrication involved sintering cold pressed bars to 

2200°C in vacuum for 4 hours and then reoxidizing to 1700°C. In the former case, 

only the monoclinic phase was observed. In the latter case, two monoclinic phases 

and a titanate phase resulted. One factor that may explain the discrepancies in the 

thermal expansion coefficient values, therefore, is that different phases due to the 

fabrication processes were being examined even though the compositions were the 

same. 

A second factor is the range over which the measurements were made. In the 

Brown et. al. study, the thermal expansion coefficient was measured from room 

temperature to only 1000°C. The thermal expansion coefficient for the same range in 

the present study is approximately 3.9 x 10-V°C. 

A third factor might be the rate at which the measurements were made. In the 

present study, the thermal expansion was measured at a rate of 2OC/min which 
101 

allowed the sample plenty of ime to achieve its equilibrium length at all temperatures. 

Although Brown et. ai. did not state their heating rate, had the measurement been 

made at a faster rate, a lower expansion might be expected. 

For the tetragonal phase, the thermal expansion coefficients for the 

compositions examined in this study were found to decrease with increasing TiO2 

(refer to Table 14). The values were much lower than the coefficients for tetragonal 

Hf02 and Zr02 (1 0.0 x 1 O-VOC and 10.6 x 1 OVOC, respectively). Since the range over 

which the measurements were made was small, the observed values might be a bit 

lower than the true values. Nonetheless, the reduced values from the pure phases 

further suggests that Ti02 lowers the thermal expansion. 

4.6.2 Transformation Tempe rat u res 

The monoclinic to tetragonal transformation temperatures can easily be 

identified in the thermal expansion curves shown in Figures 58-61. As the cutves 

suggest, the transformation is spread over a temperature range and has a thermal 

cycle hysteresis (Le., the transformation temperatures for the heating and cooling 

curves are not the same). Such behavior is typical of the Hf02 and Zr02 fluorite 

structures. The strain energy due to the diffusionless transformation (Le., atom 

displacements of less than an interatomic distance) in the fluorite structure adds an 

additional degree of freedom and consequently causes the transition to occur over a 

temperature range rather than at a single temperature. The large strains that are set 

up as the temperature increases are not relieved by atom migration, but rather by 

twinning.102 This upsetsthe energy balance such that the reaction is not reversible at 

the same temperature thereby producing a thermal cycle hysteresis.lm 

The phase diagram of the Hf02-Zr02 system (Figure 6) shows that when ZQ is 

added to Hf02, the transformation temperature range increases, while the 

transformation temperature decreases. The phase diagrams of the Hf02-Ti02 and the 
102 

Zr02-Ti02 systems (Figures 7a and 8b, respectively) show that a similar behavior 

occurs when TiO, is added up to the point where the titanate begins to form. Beyond 

this point, a levelling off of the lower monoclinic to tetragonal transformation 

temperature is observed. In the present study, the thermal expansion curves show that 

additions of Ti02 lower the transformation temperature and increase the transformation 

temperature range in a similar manner to that observed in the HT and ZT systems. 

Table 15 summarizes the transformation temperatures for the four compositions. It 

should be pointed out that at the same time that TiO, is being added in the present 

HZT system, ZrO, is being removed. This suggests that the effect of TO2 on the 

transformation temperature is greater than the effect of ZrO, since the trends due to the 

latter are totally reversed. It should also be pointed out that the levelling off of the 

transformation temperature (similar to that which is observed in the HT and ZT 

systems) occurs at the 15 mol% TiO, composition in the present HZT system. This 

corresponds to the composition at which the titanate begins to form and suggests that 

the solubility limit of TiO, in the present study is between 10 and 15 mol% 

(ca. 12 mol%). 

The transformation that occurs in going from the reduced state to the oxidized 

state was also observed using both a hot stage x-ray diffractometer and a dilatometer 

in air. This inversion involves a cubic to monoclinic phase change. For each 

composition, the transformation temperature was found to be approximately 600°C. 

Figure 62 shows the thermal expansion trace. 

The following observations were made for the 15T2 composition: (1) The initial 

thermal expansion of' the reduced state material was found to be low 

(a = 5.0 X 10-6/"C). (2) A large increase in length was observed during the cubic to 

monoclinic transformation. (3) The thermal expansion of the upper temperature region 

of the heating curve and the cooling curve were the same as that of the 15T20x thermal 

expansion in Figure 60. 
103 

Table 15. Transformation Temperatures and Ranges for the Monoclinic to 

Composition Transformation Range Transformation Temperature 

Tetragonal Inversion in the HZT System. 

(mol% TiO,) ("C) ("C) 

5 1387 - 1510 1469 

10 

15 

20 

1551 - 1632 

1346 - 1449 

1510 - 1592 

1286 - 1429 

1408 - 1592 

1286 - 1450 

1408 - 1592 

1591 

1388 

1571 

1367 

1551 

1367 

1551 

3 

Y 

a 

I0.03 , 1 

15mol% Titania 

0.02 

+ Heating 

0 Cooling 

0 400 800 1200 1600 2000 

I TEMPERATURE ("C) 

Figure 62. Linear thermal expansion of the 15 mol% TiOp composition upon 

oxidizing. 
104 

Several conclusions can be drawn from the oxidation curve. First, a rapid initial 

oxidation occurs at 600°C. Second, beyond 6OO0C, the oxidation is much more 

gradual. And third, once the oxidation is complete, the cubic phase no longer exists. 

One possible mechanism is a rapid oxidation at 600°C corresponding to the point at 

which a sufficient amount of energy is introduced to produce significant oxygen 

mobility, followed by an equilibrium mobility of oxygen as oxidation continues. 

4.6.3 Microcracki ng 

The effects of extensive microcracking can be observed in the 15 and 20 mol% 

Ti02 thermal expansion curves. A hysteresis between the hes?ing and cooling cuwes 

is evident. A correlation between the size of the hysteresis and the size of the 

microcracks can be made. The microstructures in Figures 44-52 show that the low 

TiOp compositions have very fine microcracks (less than one grain wide), whereas the 

microcracks in the high compositions are very large (approximately seven grains 

wide). Similarly, the hystereses in the thermal expansion curves of the high Ti02 

compositions are quite pronounced whereas in the low TiO, composition, the 

hysterses are negligible. The conclusion can then be made that large microcracks 

lead to the large hystereses in the high TiO, compositions. 

4.7 Fracture Strength and Toughness 

This section is divided into three parts. The first part examines the phase and 

microcracking dependencies of the strengths of reduced and oxidized compositions in 

the HZT system. Specifidally, the effects of the cubic phase on the strength of reduced 

materials and the effects of the titanate phase on the strengths of oxidized materials 

are discussed. Also, compositions with various degrees of microcracking are 

examined to illustrate the effects of microcracking on strength. The second part of this 

section examines the effects of composition on fracture toughness. Materials in both 
105 

the oxidized and reduced states are investigated. The final part of this section 

examines the effects of temperature on the fracture strength and toughness of the 

oxidized compositions in the HZT system. A similar examination of the reduced 

compositions was not made due to testing apparatus limitations and limited material 

supplies. Further examination of these materials to determine the high temperature 

strength and toughness in reducing environments is warranted. 

4.7.1 Fracture Strength 

Fracture strength values were determined by placing bars in a four-point bend 

load fixture. The tensile side of the test bars was polished to a $ pm diamond finish to 

minimize the number and size of surface flaws. The edges were bevelled to reduce 

the stress concentration effects of sharp corners. Measurements were taken of the 

load at fracture and the bar dimensions at the site of fracture. Using Equation (46), the 

modulus of rupture (or fracture strength) was calculated for each sample. 

The modulus of rupture for both indented and nonindented bars were 

determined. As discussed in Section 2.2.3.2.1 , controlled flaw fractures (as occur from 

preindented specimens) generally give reproducible MOR values. However, the 

measured values do not always give an adequate assessment of the material strength 

(Le., the measured MOR values tend to be low). On the other hand, fractures from 

nonindented specimens tend to give high MOR values but with large degrees of 

scatter. Therefore, MOR measurements for both indented and nonindented bars were 

made in order to give both an accurate assessment of the strengths of HZT 

compositions and to detefmine the reproducibility of the measured trends. 

Between three and twelve bars of each composition for each series were tested. 

Due to the low number of bars tested, a Weibull plot could not be made. Instead, the 

statistical nature of the fracture strength was given as a 90% confidence interval as 
106 

monoclinic phase. Figures 63 and 64a suggest that a similar situation occurs in the 

determined from Student's t distribution curves.104 Such curves are generally used for 

low sample populations. The strength was therefore reported as: 

(1 03) - SX 

X f0.10/2. n-1 - 6 

where TT is the average of the n measured values (e.g., MOR values), t0.10/2, n-1 is the 

value from the Student's t distributions for a 90% confidence interval for n-1 values, 

and sx is the standard deviation of the n values. 

4.7.1.1 Nonindented Bars 

Modulus of rupture values of nonindented bars were determined for each 

composition in the TI, T2, and T3 series. Both the reduced and the oxidized states 

were examined. The effects of titania on the strengths of the HZT compositions are 

shown in Figure 63. The data is replotted in Figure 64 to show the relative strengths 

for the different fabrication techniques. From Figures 63 and 64, the effects of phase 

changes and microcracking on fracture strength can be seen. 
107 

Figure 63. 

Large Microcrack Variations 

500 

Nonindented 

a- Tlad 

C+Mt -+ Tlox 

M+Ot 

-Pm law 

0 5 10 15 20 25 

COMPOSITION (mPh Titania) 

0 

(a) 

Similar Microstructures 350, 1 

150 "I 100 

50 

Nonindented 

C 

M+O 

0 5 10 15 20 25 

COMPOSITION (moPh Tiitania) 

(b) 

Minimal Second Phases i 

~ - Nonindented I 

Q T2ad 

+ T2ox 

/ I - 

M M M I 

100 j 31 w I ... 

Fracture strengths of nonindented highly anion deficient (ad) and 

oxidized (ox) HZT compositions made from (a) reaction-sintered (T1), (b) 

as-received (T2), and (c) plasma-sprayed powders (T3). Phases are 

indicated by the symbols M (monoclinic), C (cubic), and 0 

(orthorhombic). (A t indicates trace amounts.) Average microcrack 

region sizes are placed above each composition. 
108 

400- 

300 - 

200 - 

100 - 

500 1 1 

Nonindented 

C+Mt 

L 

COMPOSITION (mop/, Titania) 

200 - 

100 - 

M 

M+O 

Q Tlad 

+ T2ad 

4 T3ad 

Q Tlox 

-+ T2ox 

4 T30X 

Figure 64. Fracture strengths of nonindented HZT compositions in the TI, T2, and T3 

series in the (a) reduced and (b) oxidized states. Phases are indicated 

by M (monoclinic), C (cubic), and 0 (orthorhombic). 
109 

monoclinic phase had relatively high strengths (>250 MPa). 

The T2ad series best illustrates the effects of the cubic phase on strength. In this 

series, the microstructure was essentially the same for all compositions (Le., the grain 

size, microcrack region size, and porosity were basically identical). Consequently, the 

strength was primarily affected by compositional or phase effects. A very decisive 

decrease in strength was observed in the T2ad series with increasing cubic phase 

concentration (see Figure 63b). Hence, the strength was determined to be inversely 

proportional to the amount of cubic phase present. 

The T3ad series best illustrates the effects that suppressing the cubic phase has 

on strength. In this series, concentrations of the cubic phase were minimal (Le., 

essentially no cubic phase was observed in the 5T3ad and 15T3ad compositions, and 

only small concentrations were observed in the 20T3ad composition). Consequently, 

the effects of titania on the strength of HZT materials in the reduced state could be 

determined. Figure 63c shows that the resultant strengths of the materials were 

relatively high (~200 MPa) and increased with Ti02 concentration. A decrease in 

strength occurred at the 20 mol% composition corresponding to the emergence of the 

cubic phase. 

The Tlad series also showed a decrease in strength with increasing cubic 

phase concentration (approximately a 61 % reduction). More importantly though, the 

series illustrated the effects of microcracking on the strength. From microstructural 

analyses, the microcrack region sizes of all the compositions in the TI series were 

found to be high (~160 pm). A large amount of fluctuation in the sizes also occurred. 

From an examination of Figure 63a, a correlation between strength and microcrack 

region size can be seen. The compositions with large microcrack sizes tended to have 

low strengths. This correlation is further substantiated by the extremely low strength of 

the 20Tlad composition. (For this composition, all samples broke prior to testing due 

to the load exerted on them during polishing.) The 20Tlad compositions had very 
, 

I 110 

large microcrack regions (of the order of 3000 pm) and tended to break readily along 

region boundaries. 1 

4.7.1.1 -2 Oxidized State 

Phase analyses of the TI and T2 series showed that for HZT compositions in the 

oxidized state, the low TiO, compositions were essentially monoclinic and the high 

compositions were a mixture of monoclinic and titanate phases. As was pointed out in 

Section 2.2.3.1, the titanate phase tends to drastically reduce the strength of materials 

in the Hf02-Ti02 system (e.g., the MOR drops to approximately 40 MPa when HfTiO, 

forms). This weakening effect can be associated with the low thermal expansion and 

the high degree of microcracking that occurs in the system when the titanate is 

present. The following conclusions were made: (1) The anisotropic expansion of the 

titanate phase leads to extensive microcracking which gives low thermal expansion. 

(2) The extensive microcracking makes the material weak. 

The effects of the titanate phase on the thermal and mechanical properties of 

the present compositions in the HZT system were similar to those found in the HT 

system. Figures 50 and 51 show that a large degree of microcracking occurred in the 

high TiO, HZT compositions. Also, the thermal expansion coefficients were found to 

be moderately low (refer to Table 14). Consequently, the material was weak (Figures ! 

63a and b). 

I The T30x series, illustrates the effects that suppressing the titanate phase has 

on the strength of oxidized compositions in the HZT system. In this series, all the 

compositions were essentially pure monoclinic solid solutions with the exception that 

trace amounts of the titanate phase were found in the 20T30x composition. As pointed 

out in Sections 4.2 and 4.3, the bars approached 100% theoretical density and had no 

apparent impurities. Fracture surfaces suggest that indeed little porosity existed in 

these compositions (see Figures 65-68). Also, Figure 52 suggests that the severity of 

i 

f 

I 
111 

Figure 65. Scanning electron micrograph of the fracture surface of a 5T30x bar. 

Figure 66. Scanning electron micrograph of the fracture surface of a 10T30x bar. 
Figure 67. 

Figure 68. 

112 

Scanning electron micrograph of the fracture surface of a 15T30x bar. 

Scanning electron micrograph of the fracture surface of a 20T30x bar. 
113 

the microcracking was minimal. These factors imply that the resultant strengths are 

indicative of the true effect that Ti02 has on the fracture strength of monoclinic solid 

solutions in the system. Figure 63c indicates that the effect of TiO, is to increase the 

strength as long as the titanate phase is suppressed. 

4.7.1.2 Indented Bars 

Modulus of rupture values of indented bars were determined for each 

composition in the T1 and T2 series. The T3 series was not examined because not 

enough bars could be produced of each composition to give an adequate assessment 

of the indented strength. (The 10T3 ana 15T3 compositions suffered thermal shock 

damage during the cooling cycle of the hot press runs such that only a few one inch 

bars could be cut from the two inch disks. The bars that were produced were used in 

the nonindented fracture study.) Room temperature measurements of strength were 

made for both the reduced and oxidized states. Figure 69 summarizes the results. 

Trends in the strength measurements of the indented samples were similar to those 

observed in the nonindented samples. That is, the fracture strengths of the highly 

anion deficient and oxidized bars decreased with increasing Ti02 concentration due to 

the cubic and titanate phases, respectively. 

The strengths of the indented bars were about half that of nonindented bars. 

Section 2.2.3.2 pointed out that in indentation techniques, a decrease in strength of 

this magnitude is expected. The reproducibility, however, was much better, as the 

error bars in Figure 69 suggest. 

A difference in the hature of the indents for the T1 and T2 series resulted due to 

the difference in the method of indentation. In the T1 series, indents were first made on 

the bars in the reduced state and subsequently, the bars were oxidized. In the T2 

series bars were first oxidized and then indented. The effect on strength, from Figure 

69, was apparently minimal since the magnitudes of the MOR values of the Ti and T2 
114 

50 - 

Large Microcrack Variations 

250 > 

1 

27pm 34pm 36pn 32pm 

M M+C 

Similar Microstructures 

200 

C+M 

1 00 

M+O 

M+O 

4 Tlad 

+ Tlox 

Q T2ad 

+ T2ox 

Figure 69. Fracture strengths of indented reduced and oxidized HZT compositions 

made from (a) reaction-sintered (TI) and (b) as-received powders (T2). 

Phases and microcrack region sizes for each composition are given. 
115 

I 

series are about the same. The appearance of the indents, on the other hand, was not 

the same. In the T20x series, the indents and crack extensions appeared sharp and 

well defined whereas those in the TIox series were diffuse and irregular. Grain 

boundary etching and strain relief movement in the TI series due to annealing 

probably accounts for the difference in appearance. 

All breaks occurred at one of the three indents for all the compositions except 

the 5T1ox. For this composition, fractures never occurred at an indent and rarely 

occurred within the region where indents were placed. Both the appearance of the 

indented region and the fracture nature suggest that microcrack toughening occurred 

for this composition. A further examination of this material is warranted. 

The fracture surfaces of the indented bars in the TI and T2 ox series are shown 

in Figures 70-76. These figures confirm that a high degree of porosity was present in 

the compositions and that fracture occurred along grain boundaries. In the high Ti02 

compositions, evidence of fracture along microcrack region boundaries was apparent 

and microcracks within the inner regions of the bars could be seen (see Figure 76b). 

4.7.2 Fracture Toughness 

The fracture toughness was determined for indented bars in the T1 and T2 

series. Klc values were calculated using Equation (51) and using both the measured 

MOR values and surviving crack lengths of each sample. Plots of the results are given 

in Figure 77. 

In general, the fracture toughness decreased with increasing Ti02 for both the 

reduced and oxidized states. Trends similar to those found in the strength 

measurements were observed. In the reduced state, the toughness values of 

compositions possessing the cubic phase were found to be lower than those which 

were mostly or all monoclinic. In the oxidized state, the toughness was observed to 

decrease with increasing titanate phase concentration. 

t 
116 

Figure 70. Scanning electron micrographs of the fracture surface of a 5Tlox bar. 

L 

, 

'* 

.. 

. 

.. 

Figure 71. Scanning electron micrographs of the fracture surface of a 10T ox bar. 
117 

Figure 72. Scanning electron micrographs of the fracture surface of a 15Tlox bar. 

Figure 73. Scanning electron micrographs of the fracture surface of a 5T20x bar. 
118 

I 

Figure 74. Scanning electron micrographs of the fracture surface of a 10T20x bar. 

Figure 75. Scanning electron micrographs of the fracture surface of a 15T20x bar. 
119 

Figure 76. Scanning electron micrographs of the fracture surface of a 20T20x bar 

at (a) low magnification and (b) high magnification. 
120 

n 

E 

2 

a 

61 1 J 

5- 

4- 

3- 

2- 

M 

M 

M+Ot 

0 5 10 15 20 25 

COMPOSITION (moPh Tiiania) 

C 

M+O 

Figure 77. 

Q T2ad 

+ T20x 

0 5 10 15 20 25 

COMPOSITION (mol% Titania) 

c 

(b) 

Q Tlad 

-+ Tlox 

Fracture toughness of reduced and oxidized compositions in the 

(a) T1 and (b) T2 series. Phases and microcrack region sizes are given. 
121 

A comparison of the TI and T2 series shows that the large microcracks in the 

high Ti02 compositions of the T1 series act to accelerate crack propagation thereby 

reducing the toughness. The high Ti02 compositions of the T2 series (which do not 

have large microcracks) have higher Klc values. In the low Ti02 compositions, 

microcracking had little effect on the Klc except in the 5Tlox composition. Annealing 

of the indents in this composition apparently led to microcrack a hardening. Crack 

propagation was decelerated due to the stress relief which resulted from the alignment 

of microcracks. The bars never broke at indents and rarely broke within the indented 

region. Consequently, the measured strength as well as the measured fracture 

toughness values of this composition are most likely lower than the true material 

values. 

IA comparison with literature values suggests that the Klc values for the highly 

cubic reduced HZT compositions were of the same order of magnitude as those of fully 

stabilized cubic zirconia (Le., 2 to 4 MPa-m; as compared to 1.8). The monoclinic HZT 

compositions had Klc values of the order of 5 to 6 MPa-mi which 'is comparable to 

monoclinic HfO, and Zr02. 

4.7.3 Temperature Dependence of MOR and Klc 

The fracture strengths and toughnesses were determined as a function of 

temperature for each composition in the T1 and T2 series. Indented bars that were 

oxidized at 1200°C, were heated to 1000°C and to 1400OC in air and broken at 

temperature. Figures 78 and 79 show the-results. In general, a decrease in strength 

and toughness 0ccurred"for samples fractured at 1000°C, as compared with room 

temperature values, and an increase occurred for bars fractured at 14OOOC. 

The literature suggests that a decrease in strength and toughness with 

increasing temperature is expected for Hf02- and ZrO2-based materials (refer to Figure 

24). On the other hand, an initial increase in strength with temperature is expected for 
122 

200 

150 

100 

50 

0 - 

0 400 800 1200 1600 

TEMPERATURE (“C) 

(a) 

250 1 

T c t 

150 

100 

50 

J. 

0 1...1...1... 

Q 5T1 

+ lOT1 

4 15T1 

Q 5T2 

+ 10T2 

4 15T2 

-9 20T2 

- 

0 400 800 1200 1600 

TEMPERATURE (“C) t 

Figure 78. Temperature dependence of the fracture strength for compositions 

in the HZT system in the (a) T1 and (b) T2 series. 
123 

I 

t 

4 5A 

3-i . 

0 400 800 1200 1600 

TEMPERATURE (“C) 

6 

5 

4 

3 

2 

1 

0 

0 400 800 1200 1600 

TEMPERATURE (“C) 

-b 5T1 

+ lOTl 

4 15T1 

Q 5T2 

-+ 10T2 

4 15T2 

-0 20T2 

Figure 79. Temperature dependence of the fracture toughness for compositions 

in the HZT system in the (a) T1 and (b) T2 series. 
124 

materials which have microcracks (refer to Figure 25). In the present study of the HZT 

system, the results were mixed. A slight decrease in strength (several percent) 

occurred at 1000°C which was offset by an increase at 1400OC. As was pointed out in 

Section 2.2.3, strength tends to be invariant (or increase slightly) with temperature in 

the range where fracture is elastic. Preexisting flaws and microcracking dominate the 

fracture behavior in this temperature range. 

The increased strengths at 14OOOC suggest that microcracking effects (Le., 

microcrack healing) are present in the HZT compositions. (The higher temperature 

elicits a material expansion that closes microcracks and consequently increases the 

strength.) The lowered strengths at 1000°C suggest a situation exists whereby 

microcracking does not dominate the fracture behavior. 

The lower values of strength and toughness found in the T1 series as compared 

with the T2 series confirm the correlation between microcrack size and fracture 

behavior presented in the previous sections. In the T2 series, microcracks were small 

(approximately 30 pm) giving high strength and toughness values. In the T1 series, 

microcracks were larger giving slightly lower values. 

The effects of microcrack healing on strength and toughness can be associated 

with microcrack size. Microcracks in the T1 series were found to be many times wider 

than those in the T2 series (refer to Figures 44-51). The enhancement in strength 

associated with the closing of the narrow cracks in the T2 series was greater than that 

of the wider cracks in the T1 series. A comparison of the curves in Figures 78 and 79 

illustrates the magnitude of the strength improvement effects. 

In addition to mic;ocrack healing effects, oxidation effects were also involved. 

The materials studied in this investigation were initially oxidized at 12OOOC. 

Subsequently, the materials were tested at temperatures above and below this 

temperature. The indentation procedures for the two series, however, were not the 

same. As was pointed out in Section 4.7.1.2, bars in the T1 series were indented and 
125 

then oxidized, whereas bars in the T2 series were oxidized then indented. An 

examination of Figures 78 and 79 shows that the severity of the reduction in strength 

and toughness with temperature is less in the T1 series than in the T2 series. This 

difference suggests that during the oxidation of the indented T1 bars, an annealing or 

crack healing of the indents occurred. Furthermore, the indented T2 bars which were 

not annealed prior to fracture did not experience crack healing. 

In summary, microcrack healing in both the T1 and T2 series accounted for the 

increase in strength and toughness with temperature. Those compositions which had 

noticeable increases tended to be in the T2 series where microcracks were relatively 

small. Large microcracks in the T1 series led to less strength improvement. 

t 

I 
I 126 

5. CONCLUSIONS 

I (1) The solubility limit of Ti02 in the studied HZT system was determined to be 

between 10 and 15 mol%. Phase analyses showed that below 15 mol%, the HZT 

compositions were monoclinic solid solutions, while above 15%, the materials were 

I titanates in a monoclinic solid solution. 

I (2) The oxidation states and fabrication processes were found to significantly 

affect the final phases and microstructure of the compositions examined in the HZT 

system. X-ray analyses showed that as the titania concentration was increased, an 

increase in the concentration of the cubic phase in the reduced state was observed. 

The degree of anion deficiency produced during fabrication determined the amount of 

cubic phase present. In the oxidized state, an increase in the titanate phase 

~ 

I 

concentration was observed. Incomplete oxidation tended to suppress the titanate 

phase and extend the monoclinic solid solution limit. Microstructural analyses showed 

that hot pressed materials were highly dense, whereas sintered materials had up to 

8% porosity. While microcrack sizes varied greatly with composition and fabrication 

technique, grain sizes were approximately 1 pm for all compositions and varied only 

slightly. 

I 

I 

(3) A model was proposed for the PO, dependence on the electrical 

conductivity of TiO, doped Hf02-rich mixed oxides. In the model, the titanium was 

incorporated interstitially at high PO, and substitutionally at low PO,. The measured PO, 

dependence showed similar trends to those suggested by the model. 

(4) An explanation of the low thermal expansion of materials in the HZT 

system was given. De6ending on the oxidation state, materials had low thermal 

expansion coefficients due to either the presence of the titanate phase or the presence I 

I 

i of the high Ti02 monoclinic phase. 

Thermal expansion coefficients were shown to decrease with increasing titania 

concentration for HZT compositions. Microcracks were also shown to reduce the 
127 

thermal expansion. Extensive microcracking (resulting from the anisotropic nature of 

the titanate phase) occurred in the high Ti02 compositions leading to low thermal 

I expansion coefficients and producing thermal cycle hystereses as well. 

I 

I 

I (5) Fracture strength and toughness increased with Ti02 when the titanate 

phase was suppressed. When the titanate phase formed, the strength and toughness I 

I 

I 

I was reduced substantially. Microcracking also weakened the high TiOp compositions 

and strengthened the low compositions under certain oxidation conditions. 

I 

I 

The temperature dependence of the strength and toughness went through a 

minimum and then increased. The increase in strength was attributed to microcrack 

healing at temperatures above the oxidation temperature. 
128 

6. SUGGESTIONS FOR FUTURE WORK 

(1) Phase analyses of fully oxidized materials for a wider range of titania 

compositions need to be performed. 

(2) Variations in the concentrations of hafnia and zirconia need to be made in 

order to more fully study the effects they have in the HZT system. 

(3) A ternary phase equilibrium diagram of the HZT system needs to be 

constructed. 

(4) The mechanism which inhibits sintering and grain growth in TiO, doped 

hafnia-rich mixed oxides needs to be determined. 

(5) Electrical conductivity measurements at several more oxygen partial 

pressures need to be made in order to fully prove the defect model. 

(6) Linear thermal expansion data needs to be taken to higher temperatures in 

order to accurately determine the tetragonal thermal expansion coefficient. 

(7) Fracture strength and toughness measurements need to be made above 

1400°C to fully determine the high temperature strength effects. 

(8) The effects of various atmospheres (e.g., reducing) on the thermal 

expansion and strength of compositions in the HZT system need to be examined. 

(9) The effect of titania on the thermal conductivity, creep, and elastic modulus 

in the HZT system needs to be determined. 
129 

APPENDICES 

APPENDIX 1. Impurity Content of Hafnia, Zirconia, and Titania Used in the HZT 

Experiments. 

Table 1.1 Impurity Content of Teledyne Wah Chang Albany HfO,. 

Lot Analvsis in Dm 

AI <35 

B 0.4 

Ca/Hf02 <20 

Cd <0.2 

co e5 

Cr <20 

cu <25 

Fe 40 

Mg <1 0 

Mn <25 

Mo <25 

Na/Hf02 <20 

Ni <25 

Pb <1 0 

Si <1 00 

Sn <1 0 

Ti e25 

V e1 0 

W 4 0 

Zn <50 

Zr 0.98% 

c Element Results 

U/Hf02 12 

Impurities on Hf metal basis unless otherwise noted. 
130 

Table 1.2 Impurity Content of Magnesium Elektron Zr02 Development Grade 

XZ0544/ 1. 

bnaly& 

Properties Results 

Si02 0.1 2% 

Ti02 0.12% 

Fe203 <0.02% 

so3 0.07% 

Particle Size (d50 Coulter) 

Surface Area (Quantasorb) 3.8 m2/g 

1.9 pm 

Table 1.3 Impurity Content of J.T. Baker Chemical Company Ti02 Baker Analyzed 

Reagent. 

Analvsls 

Element Results 

As <.00005% 

Fe 0.01 % 

Pb 0.005% 

Zn O.Ol0/0 

€ 
131 

APPENDIX II. Thermal Expansion of Standard Single Crystal Sapphire. 

pintinnal pureau of Btnndnrde 

&xiifit ate 

Standard Reference Material 732 

Skgle Crystal Sapphire - Thermal Expansion 

T.A. €fa& 

TbcrmJ Expanrbn and Expuuivi~ u a Funaioo d Tempenium (I-) 

~ 

0 x IO* S.31 x IO4 K-‘ 

38 S.48 

3m 6.12 

647 6.61 . 

9n 1.00 

1346 1.32 

1719 1.59 

21w 1.82 

wo tu2 

7##. h20 

3320 L31 

3743 Ls2 . 

I dt 

1 Ln3 =-G,m 

9001: 4608 x IO- 8.79 I IO4 K-’ 

IO00 5m 9.0s 

I loo 6417 9.29 

1200 1357 9.52 

1300 a320 9.14 

1406 9306 9.97 

ldQ0 1134s 10.42 

1100 12399 10.66 

1800 rwn IO.@ 

1944 tcsn 11.14 

2ooo 15704 11.38 

Ixy) 10314 1o.m 

Wuhingwn.0.C 20234 

occ~bcr J, 19n 
132 

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141 

VITA 

Paul Russell Staszak was born in Chicago, Illinois on February 21, 1960. He 

received a B.S. degree in 1982 and an M.S. degree in 1985 from the University of 

Illinois at Urbana-Champaign in Ceramic Engineering. He served as both a research 

and teaching assistant during his graduate career at the University of Illinois. His 

teaching duties included laboratory courses in ceramic x-ray analysis and glass. His 

graduate research work included studies of the structure of germanate glasses and 

high temperature ceramics. He also spent two summers as a student research 

assistant at Argonne National Laboratory working on the laser pyrolysis of amorphous 

metal powders. He has two publications, one on his work with laser pyrolysis, the 

other on a space group calculation of magnesium spinel. 

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