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The early stages of crystallization in alkali-silicate glasses

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Title:
The early stages of crystallization in alkali-silicate glasses
Added title page title:
Alkali-silicate glasses
Creator:
Kinser, Donald L., 1941-
Publication Date:
Copyright Date:
1968
Language:
English
Physical Description:
xiv, 143 leaves. : illus. ; 28 cm.

Subjects

Subjects / Keywords:
Conductivity ( jstor )
Crystallization ( jstor )
Crystals ( jstor )
Dielectric materials ( jstor )
Electrodes ( jstor )
Electron micrographs ( jstor )
Heat treatment ( jstor )
Lithium ( jstor )
Nucleation ( jstor )
Soft drinks ( jstor )
Crystallization ( lcsh )
Dissertations, Academic -- Metallurgical and Materials Engineering -- UF
Glass ( lcsh )
Metallurgical and Materials Engineering thesis Ph. D
Genre:
bibliography ( marcgt )
non-fiction ( marcgt )

Notes

Thesis:
Thesis - University of Florida.
Bibliography:
Bibliography: leaves 138-142.
Additional Physical Form:
Also available on World Wide Web
General Note:
Manuscript copy.
General Note:
Vita.

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University of Florida
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University of Florida
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Copyright [name of dissertation author]. Permission granted to the University of Florida to digitize, archive and distribute this item for non-profit research and educational purposes. Any reuse of this item in excess of fair use or other copyright exemptions requires permission of the copyright holder.
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20143525 ( OCLC )
AFN0638 ( NOTIS )

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Full Text




THE EARLY STAGES OF CRYSTALLIZATION
IN ALKALI-SILICATE GLASSES













By

DONALD L. KINSER


A DISSERTATION PRESENTED TO THE GRADUATE COUNCIL OF
THE UNIVERSITY OF FLORIDA
IN PARTIAL FULFILLMENT OF THE REQUIREMENTS FOR THE
DEGREE OF DOCTOR OF PHILOSOPHY











UNIVERSITY OF FLORIDA
1968




























Dedicated to my wife,

Barbara









ACKNOWLEDGEMENTS


The author would like to acknowledge the assistance

of his advisory committee, Drs. Rhines, DeHoff, Reid and

Carr, and especially his chairman Dr. Hench who contributed

a great deal in the form of discussions, advice and

encouragement throughout the course of this investigation.

The numerous discussions with Dr. Hren are also gratefully

acknowledged.

Thanks are also due to numerous members of the

faculty, staff and students of the Metallurgy Department

for many stimulating discussions and a great deal of

experimental assistance. The aid of Mr. E. J. Jenkins

with electron microscopy techniques and Mr. D. A. Jenkins

with many experimental problems is gratefully acknowledged.

The author is also grateful for the financial

support of the National Science Foundation and the United

States Air Force without whose support this research

would not have been possible.


iii









TABLE OF CONTENTS


ACKNOWLEDGEMENTS ...

LIST OF TABLES .....

LIST OF FIGURES ....

NOMENCLATURE .......

ABSTRACT ...........


Chapter

I.

II.

III.

IV.

V.

VI.

VII.

VIII.


INTRODUCTION ............

EXPERIMENTAL PROCEDURE ...

POLARIZATION .............

EXPERIMENTAL RESULTS .....

DISCUSSION OF RESULTS ....

SUMMARY ..................

CONCLUSIONS ..............

RECOMMENDATION TO FUTURE
INVESTIGATORS ..........


BIBLIOGRAPHY .....................

BIOGRAPHICAL SKETCH ..............


Page

iii

v

vi

xi

xiii


1

19

42

53

108

130

134


136


138

143


............

............

............

............

............


: : : :








LIST OF TABLES


Table Page

1. ANALYSIS OF GLASS RAW MATERIALS ............. 23

2. SUMMARY OF X-RAY LATTICE SPACINGS AND
RELATIVE INTENSITIES FOR 30 MOLE %
LITHIA GLASSES .............................. 54

3. SUMMARY OF X-RAY LATTICE SPACINGS
AND RELATIVE INTENSITIES FOR 30
MOLE % LITHIA-SILICA GLASSES ................ 56

4. SUMMARY OF X-RAY LATTICE SPACINGS
AND RELATIVE INTENSITIES FOR 33
MOLE % SODA-SILICA GLASSES .................. 58

5. SUMMARY OF LATTICE SPACINGS OBTAINED
FROM ELECTRON DIFFRACTION PATTERNS
OF 30 MOLE % LITHIA GLASS AT 4800C .......... 71

6. SUMMARY OF COEFFICIENTS OF LINEAR
EXPANSION AND LOWER TRANSFORMATION
POINTS ..................................... 107








LIST OF FIGURES


Figure Page

1. Schematic melting procedure for 30
mole % lithia-silica glass ................ 20

2. Schematic diagram of x-ray source
and Guinier-DeWolff camera ................ 27

3. Photograph of sample holder employed
in electrical measurements ................ 32

4. Schematic diagram of audio frequency
measurement equipment ..................... 34

5. Schematic diagram of transformer
ratio arm bridge .............................. 35

6. Schematic diagram of radio frequency
measurement equipment ..................... 37

7. Schematic diagram of null detection
system ....................................... 38

8. Schematic diagram of DC measurement
equipment .................................... 40

9. DC polarization curve of 30 mole %
lithia-silica glass in as cast and heat
treated 5 hours at 5000C forms ............ 43

10. Conductivity versus log time curve for
33 mole % lithia-silica glass as cast ..... 45

11. Log DC conductivity versus reciprocal
temperature for 33 mole % lithia-
silica glass as cast ...................... 46

12. Schematic diagram of model used for
polarization calculation .................. 48

13. Replica electron micrograph of 30
mole % lithia-silica glass (30,000x)
as cast .................................... 59







LIST OF FIGURES--Continued


Figure Page

14. Replica electron micrograph of 30
mole % lithia-silica glass (26,000x)
after 5 hours at 5000C ............ ....... 60

15. Replica electron micrograph of 30
mole % lithia-silica glass (31,000x)
after 10 hours at 500C ............... .... 61

16. Replica electron micrograph of 30
mole % lithia-silica glass (31,000x)
after 20 hours at 5000C ................... 62

17. Replica electron micrograph of 30
mole % lithia-silica glass (37,000x)
after 50 hours at 5000C ......... .......... 63

18. Electron diffraction pattern of as
cast 30 mole % lithia-silica glass
at room temperature ................. .. ... 66

19. Electron diffraction pattern of 30
mole % lithia-silica glass at 4800C
after 5 minutes at 4800C .................. 67

20. Electron diffraction pattern of 30
mole % lithia-silica glass at 4800C
after 14 minutes at 4800C ................. 68

21. Electron diffraction pattern of 30
mole % lithia-silica glass at 4800C
after 26 minutes at 4800C ................. 69

22. Electron diffraction pattern of 30
mole % lithia-silica glass at 4800C
after 2.5 hours at 480C .................. 70

23. Electron diffraction pattern of 30
mole % lithia-silica glass at 4800C
after 3.75 hours at 4800C ................. 70

24. Transmission electron micrograph 30
mole % lithia-silica glass at room
temperature as cast (41,000x) ............. 73


vii






LIST OF FIGURES--Continued


Figure Page

25. Transmission electron micrograph of
30 mole % lithia-silica glass at
4800C after 11 minutes at 4800C
(42,000x) ................................. 75

26. Transmission electron micrograph of
30 mole % lithia-silica glass at 4800C
after 35 minutes at 4800C ................ 76

27. Transmission electron micrograph of
30 mole % lithia-silica glass at 480C
after 2.3 hours at 4800C ................. 77

28. Transmission electron micrograph of
30 mole % lithia-silica glass at 4800C
after 4.8 hours at 4800C (23,000x) ....... 78

29. Transmission electron micrograph of
30 mole % lithia-silica glass at 5000C
after 1 hour at 5000C ........... ........ 80

30. Electron diffraction pattern of 30
mole % lithia-silica glass at 5000C
after 1 hour at 5000C ..... ............. 82

31. Transmission electron micrograph of
30 mole % lithia-silica glass at
room temperature after 5 hours at
5000C in bulk form (31,000x) 83

32. Tan 6 versus logio frequency for the
30 mole % lithia-silica glass as cast .... 85

33. Tan 6 versus logio frequency for the
30 mole % lithia-silica glass after
5 hours at 5000C .......................... 86

34. Tan 6 versus logo frequency for the
30 mole % lithia-silica glass after
10 hours at 500 C ...... ................... 87


viii






LIST OF FIGURES--Continued


Figure Page

35. Tan 6 versus logio frequency for the
30 mole % lithia-silica glass after
20 hours at 500 C ........................... 88

36. Tan 6 versus logo frequency for the
30 mole % lithia-silica glass after
50 hours at 5000C .......................... 89

37. Log of frequency maxima versus
reciprocal temperature for various
thermal treatments at 5000C ............... 90

38. Tan 6 versus logo frequency for
the 30 mole % lithia-silica glass
for various thermal treatments
measured at 800C ............................ 91

39. Tan 6 versus loglo frequency for
the 33 mole % lithia-silica glass
for various thermal treatments ............ 92

40. Tan 6 versus log10 frequency for
the 26.4 mole lithia-silica glass
for various thermal treatments ............ 94

41. Tan 6 versus logo frequency for
the 33 mole % soda-silica glass
for various thermal treatments ............ 95

42. Tan 6 versus logo frequency for
the 25 mole % soda-silica glass
for various thermal treatments ............ 96

43. Log DC conductivity versus reciprocal
temperature for the 30 mole % lithia-
silica glass .................................. 98

44. Log DC conductivity versus reciprocal
temperature for the 33 mole % lithia-
silica glass ................................ 100






LIST OF FIGURES--Continued


Figure Page

45. Log DC conductivity versus reciprocal
temperature for the 26.4 mole %
lithia-silica glass ......................... 101

46. Log DC conductivity versus reciprocal
temperature for the 33 mole % soda-
silica glass ................................ 102

47. Log DC conductivity versus reciprocal
temperature for the 25 mole % soda-
silica glass ................................ 104

48. Thermal expansion curve for the 26.4
mole % lithia-silica glass as cast .......... 105

49. Phase diagram for the lithia-silica
system ..................................... 116

50. Free energy-composition diagram for
the lithia-silica system at 500C ........... 117

51. Summary of results for the 30 mole %
lithia-silica glasses ....................... 122

52. Thermal treatment time required to
develop loss peak versus composition
for lithia-silica glasses .................. 123

53. Phase diagram for the soda-silica
system ..................................... 125

54. Free energy-composition diagram for
the soda-silica system at 5500C ............. 126









NOMENCLATURE


Subscripts 1 and 2 refer to the parameters for the

matrix and dispersed phase respectively.


A Shape parameter defined by equation (1-11)

d Interplanar spacing

E Electric field

F Frequency

Q Quench rate

R Resistance

T Absolute temperature

t Sample thickness

t+, t Transferrence number of positive and negative
ion respectively

V Volume fraction of dispersed phase

x Distance parameter defined on Figure 12

tan 6 Tangent of the loss angle 6

e' Real part of complex dielectric constant

E"' Imaginary part of complex dielectric constant

E Static dielectric constant
s
Dielectric constant extrapolated to high
frequency

p Surface charge density

a DC conductivity






NOMENCLATURE--Continued


T



w
tan
tan total


tan 6AC


Period of oscillation

Electric potential

Angular frequency (2nrf)

Tangent of the loss angle


AC component of the loss angle


xii






Abstract of Dissertation Presented to the Graduate Council
in Partial Fulfillment of the Requirements for the
Degree of Doctor of Philosophy


THE EARLY STAGES OF CRYSTALLIZATION
IN ALKALI-SILICATE GLASSES


By

Donald L. Kinser

December, 1968


Chairman: L. L. Hench
Major Department: Metallurgy and Materials Engineering


The initial stages of crystallization in glassy

systems are of importance because of their influence upon

the structure of the fully crystallized material. The

precise nature of the structural changes occurring in the

early stages of crystallization was the subject of this

investigation.

This investigation has shown the appearance of a

metastable transition phase prior to the appearance of

the equilibrium precipitate in four of the five glasses

examined. The metastable phase enables the crystallization

behavior of these glasses to be explained as a process

involving the metastable precipitate as a nucleation site

for the equilibrium lithium disilicate phase.


xiii






The compositions examined include a 26.4, 30 and

33 mole % lithia-silica glass and a 25 and 33 mole % soda-

silica glass. In order to detect the structural changes

occurring during nucleation and crystal growth it was

necessary to apply new x-ray, transmission electron

microscopy, AC and DC electrical measurement techniques

to the glasses.


xiv











CHAPTER I

INTRODUCTION


Crystallization of glasses has been the subject of

scientific interest since the early work of Reaumur (1739).

He crystallized a glass bottle by imbedding it in sand

and heating it to elevated temperatures for an extended

time and found that the resulting crystalline article

was extremely brittle and technically useless. In marked

contrast to this experiment the recent work of Stookey

(1962) led to the development of a class of crystallized

glass-ceramic materials (Pyrocerams) which are quite

strong (20,000 psi-50,000 psi) and as a result are

technically useful materials. The principal difference

between the two products which gives rise to the striking

contrast in mechanical properties is the extremely small

crystallite size of the order of 1 micron in the Pyroceram

materials as opposed to large crystals of the order of

1 millimeter (mm) for the Reaumur product. It is thus

evident that the understanding of the physical processes

which lead to an extremely fine grain size are of

considerable importance both from a technical and a

scientific point of view.


- 1 -





- 2 -


In the case of crystallization in a glass matrix

it has been generally accepted that the number of crystal

nuclei present in the glass determine the ultimate lower

limit to the grain size in the final crystalline product.

Extreme importance is thus necessarily attached to the

question of identity of the nuclei as well as the behavior

of these nuclei under varying conditions.

The objective of the present work is to determine

the nature and behavior of the structural changes occurring

during the nucleation of simple binary alkali silicate

glasses which in combination with other components form

the basis of the Pyroceram crystallized glass-ceramics.

In practice, the nucleation step in the preparation

of a glass-ceramic is carried out at a lower temperature

(nucleation temperature) than is the crystallization or

growth step. Experimental observation of the structural

changes on the size scale necessary to resolve nuclei in

the early stages of crystallization requires an experi-

mental technique capable of detecting and characterizing

small crystals in the glass matrix. This problem is in

some ways analagous to problems encountered in studies of

the precipitation stages in age hardening aluminum alloys.

One of the first experimental techniques employed in

aluminum alloys was DC conductivity studies which gave a

great deal of information about these systems. These





- 3 -


systems have subsequently been investigated by electron

microscopy, x-ray small angle scattering, and several

other techniques. For these reasons and others which

will become evident, electrical properties, electron

microscopy and x-ray diffraction have been used to study

the precipitation or nucleation stage in the crystalli-

zation of glasses. In order to understand and interpret

the results of the experiments described later it is

necessary to understand the conduction mechanisms oper-

ating in these glasses as well as the parameters which

affect them. The historical review which follows is an

attempt to place the present work-in proper perspective

with previous investigations.


DC Conduction


The DC behavior of glasses has been the subject

of study for over 200 years. Early workers in the area

include such eminent scientists as Franklin, Kohlrausch

(1847), Hopkinson (1876), the Curies and Maxwell (1891).

Other lesser known but important workers include

Warburg (1884), Tegetmeier (1890), Fousserau (1883) and

others.

Warburg's early work established the presently

accepted fact that electrical conduction takes place by

the motion of the alkali metal ions in alkali-silicate





- 4 -


glasses. This was established by passing a direct current

through a test tube of "Thuringian"(CaO-MgO-Na20-K20-

A1203-Si02) glass which was filled with either mercury or

sodium amalgam which served as electrodes. Pure mercury

electrodes gave rise to polarization which caused the

current flowing to drop to less than one-thousandth of

its initial value in one hour. In the case of the sodium

amalgam electrodes the current was observed to be constant

over a protracted period of time; thus it was inferred

that the sodium from the amalgam was the charge carrier

in the glass.

Le Blanc and Kerschbaum (1910) repeated Warburg's

work and concluded, perhaps prematurely, that conduction

is due entirely to the motion of sodium ions present in

the glass. Kraus and Darby (1922) recognized that other

conduction mechanisms could also be operating, and confirmed

Faraday's law for charge transport in the glass by a

relatively simple mass transport experiment. The probable

error in the values of the transport number obtained is

too large to allow one to assert conclusively that the

conductivity was entirely ionic. More recently Kirby

(1950) has reported that the transference number of the

sodium ion is greater than 0.995. Russian workers have

also been active in this area as reported by Mazurin

(1965) but no values of the transference number are reported.





- 5 -


Hughes and Isard (1968) have recently made transference

measurements on ternary and higher order systems and

reported values very near unity.

From time to time various workers (Poole 1921,

Cohen 1957 and others) have reported work with many

glasses at high electric field strengths as well as work

in glasses containing transition metal oxides. They have

concluded that electronic conductivity is in fact

operative under either of the two conditions above. The

fact remains, and is almost universally accepted (Owen

1963, Morey 1954, and Mazurin 1965), that in alkali-

silicate glasses alkali ion motion is responsible for

electrical conduction except perhaps at high field

strengths.


Thermal History Effects


Thermal history effects were first recognized by

Fousserau (1883) in the observation of a marked time

dependance of electrical conductivity during thermal

treatments near the annealing range. Extensions of that

work carried out by Fulda (1928), Mulligan, Ferguson and

Rebbeck (1925), Littleton and Morey (1933), Littleton

and Wetmore (1936), Foex (1944), Joyner and Bell (1953)

and others have shown that electrical conductivities in

glasses are extremely sensitive to thermal history. This





- 6 -


strong dependance of conductivity upon thermal history is

generally manifested as a decrease in the conductivity

with increasing thermal treatment. Concurrent with the

conductivity changes,densification is observed. This

observation gives one explanation of the conductivity

behavior. Densification of the glass structure leaves

less void space in the glass so that motion of the alkali

ions is hindered by steric constraints so that conductivity

is reduced during thermal treatments prior to the

appearance of the equilibrium crystalline phases.


Composition Effects


Compositional effects in binary alkali-silicate

glasses are generally well documented and not difficult to

rationalize. Consider first the case of increasing

alkali content in a binary glass. It is evident from

the conduction mechanism that an increased concentration

of the alkali or current carrying species should increase

the conductivity. This behavior has been observed in

the lithia, soda, potash and cesium glasses by Mazurin

(1965) and Owen (1963).

Examination of the molar equivalent composition

glasses in the series lithium, sodium, potassium and

cesium shows that the conductivity decreases in the series

in an inverse relation to the ionic radii. Once again





- 7


this is rather easily rationalized on the basis of steric

considerations.


Electrode Effects


Ionic conduction gives rise to the phenomenon of

electrode polarization or time-dependent conductivity.

This has been recognized by relatively few investigators

since the early work of Warburg. Most investigations

are conducted under conditions which the investigators

apparently assume are non-polarizing.

Proctor and Sutton (1959, 1960) have examined

the polarization in an alkali-lead-silicate glass by

applying a DC potential and measuring the potential

distribution as a function of time. Their results agree

qualitatively with predicted behavior which assumes that

only one ion is mobile and it is blocked at the electrodes.

Conditions which should lead to the smallest

amount of polarization are (1) small field strength,

(2) short time of measurement or AC extrapolations and

(3) non-polarizing electrodes. The conditions above give

rise to serious problems.

The first solution, i.e., that of employing small

field strengths, introduces the problem of measuring

extremely small currents which in turn introduces obvious

experimental difficulties. Short time or AC measurements





- 8 -


require the definition of a time at which the electrode

charge build-up is negligible and thus introduces an

arbitrary variable. Non-polarizing electrodes are

difficult to prepare because of the requirement that the

activity of the charge carrying species in the glass be

equal to that of the corresponding species in the

electrode. This problem is soluble but it is not the

most difficult problem because the "electrode" with the

correct alkali ion activity must then be connected to

some electronic conductor to allow measurements to be

performed. Connection of the electrode to the measuring

circuit then introduces problems similar to the original

one. Electrode polarization in DC measurements has been

avoided in almost all the reported work by allowing only

very small currents to flow in the glass. This has given

rise to considerable experimental difficulty in measuring

the small currents. Some investigators have used AC

measurements of conductivities to circumvent the problem.


AC Behavior


AC electrical properties in the audio frequency

range have been the subject of extensive investigations

for only about 40 years. McDowell and Begeman (1929)

were the first to conduct comprehensive studies of AC

properties of glasses. They examined six glass





- 9 -


compositions including lead glasses, borosilicates and

lead borosilicates over a frequency range of 800 Hz to

1500 KHz and concluded that the dielectric losses could

be satisfactorily rationalized on the basis of the behavior

of individual "molecules and ions" rather than on the

basis of the heterogeneous theories of Maxwell (1892) and

Wagner (1914). Strutt (1931) examined five commercial

glasses including a borosilicate, a soda-lime-silicate

and a "heavy" lead glass and observed a strong correlation

between resistivity and dielectric loss which has been

noted by a great many subsequent workers. Strutt also

formulated an empirical equation (1-1) to describe the

behavior of the dielectric loss.


tan 6 = A exp aT (1-1)


The constants A and alpha in the equation depend upon the

glass composition as well as the measuring frequency.

Other workers, notably Stevels (1946, 1950) and Moore and

DeSilva (1952), have largely corroborated the validity

of Strutt's equation (1-1).

Robinson (1932) examined several glasses using

"non-polarizing electrodes" and concluded that the

observed behavior could be rationalized equally well by

either the Debye dipolar theory or the Maxwell-Wagners

heterogeneous dielectric theory.





- 10 -


In addition to the experimental work cited above,

a great deal of theoretical work was done in the period

since the observation of dielectric absorption by

Hopkinson (1876). Several books (Debye 1929, Smyth 1955

and Frohlich 1958) and innumerable papers treating the

ionic theory of dielectric loss in solids, liquids and

gasses have appeared so that discussion of these theories

is unnecessary. Most of the theories developed give

equations which are reducible in form to the equations

commonly called the Debye equations.


E- E
E =E + s (1-2)
00 1 + W2T2



S" (1-3)
1 + 2T2


t- ( E ]
tan = (1-4)
E + E 2 T2
S

Concurrently with the development of the ionic or

atomistic models, several workers commencing with Maxwell

(1891) have developed macroscopic or heterogeneous

theories to express the behavior of dielectric losses.

Maxwell's initial work developed the frequency dependance

of the dielectric loss for a stratified dielectric model.

This was extended by Wagner (1914) to a uniform dispersion





- 11 -


of spheres and subsequently by Sillars (1937) to spheroids.

A model for ellipsoids was developed by Fricke (1953).

The recent article by van Beek (1967) includes a summary

of equations characterizing the above systems and a great

many other types of dispersions. The equations below

from van Beek characterize a system of dispersed spheres

of conductivity 02 and dielectric constant E2 with volume

fraction V dispersed in a matrix of conductivity oa and

dielectric constant E1.


201 + 02 + 2Vv(o2 01)
S20 + 02 V (02 01) +


(2ai + 02) (62 E1) (2E1 + E2) (02 0i)
x (1-5)
|2ai + 02 V,(o2 u)12


2E1 + E2 + 2VVIE2 E)
E = E (1-6)
O 2E1 + E2 + V (E2 E1)


The relaxation time (T) of the system above is expressed

as:

S= 2 + 2-V (E2 El)
T = O
(1-7)
201 + 02-V (02 o01


The analagous equations for dispersed spheroids

derived by Sillars are considerably more cumbersome than

the sphere model equations. They are given below.




- 12


oC + IA (1 vV + v (C2 1)


I1 + A
a


(1 V ) (02 Ci)


+ V 02
v


x 1 + Aa(2 2 o C a + A (E2-E1) (C2- 1


Ioi + Aa(1 VV) (2 i1 2


8 = -




T =


(1-8)


(1-9)


1 + Ia (1-V)+VI ([2-E8]

81 + A (1-VJ)(C2-E 1)


1 + Aa(1 V)(T' El)


a1 + A (1 V) ](2
a v


(1-10)


- 01]


For prolate spheroids (a > b)


a
A + n{
a ( 2 -) n12 3/2b
F bm


b 2


- li 1/2}


(1-lla)


For oblate spheroids (a < b)


1 5
A = arc cos
a 1- (a)2 1 (a)213/2


For spheres (a = b)


1
A -
a 3


Because of the complexity of the above equations the


8 = 81


(1-11b)


(1-llc)





- 13 -


conditions of small volume fraction and oa >> 02 are

generally imposed in order to simplify the equations to

the form below.



T = o + A 2 (1-12)
A a2


V
S= e 1 + (1-13)


E2 E1

SV E iE + A (E2 1)


Most previous calculations have been made using the

simplified equations but the limitations imposed to derive

these equations must not be neglected.

It is apparent from the above discussion that two

methods of analyzing dielectric losses in materials are

the ionic, atomistic theory or the macroscopic, heterogene-

ous theory. In the past 10 years both types of analysis

have been applied. The work of Taylor (1957), Heroux

(1958), Isard (1962) and Prod'homme (1960) are examples

of the ionic or atomistic type analysis.

Taylor examined the AC properties as functions

of frequency and temperature in three soda-lime-silica

glasses, "commercial sheet glass," and a "Pyrex" type





- 14 -


borosilicate glass. He concluded that the dielectric

relaxation observed was a result of the motion of alkali

ions in the random glass structure. His analysis further

indicated that the distribution of relaxation times was

very similar for all the glasses examined. Prod'homme

essentially verified Taylor's results although there was

some disagreement as to the breadth of the distribution

of relaxation times. Barton (1965) has analyzed his

results in a manner similar to the workers above and has

been able to obtain qualitative results upon the depth

of the ion potential wells which indicate that the

potential well depths depend upon the "occupancy time."

To the knowledge of the author, none of the

workers who attribute dielectric losses to ion motion

have examined the effect of thermal treatments or thermal

history on either the dielectric losses or the distribution

of relaxation times.

The work of Isard (1962), Owen (1961) and Charles

(1963) is of importance because they analyzed their

results on the basis of a heterogeneous dielectric theory.

Owen concluded from analysis of the dielectric loss

behavior of CaO-B203-A1203 glasses that dielectric losses

occurred by the Maxwell-Wagner heterogeneous mechanism.

He further concluded that the borate rich phase was

distributed as a dispersed phase in an alumina rich matrix.





- 15 -


Observations were made on glasses thermally treated in the

annealing-transformation range but no interpretation was

given as to the morphological changes occurring.

Isard's work with a number of glasses has led to

the conclusion that the "classical theory of inhomogenity

can satisfactorily explain the main loss peak in glass."

However, he goes on to say that the high frequency (> 1MHz)

behavior requires an explanation based on the atomistic

theory.

Charles (1963) has examined the dielectric

behavior of a series of lithia-silica glasses with

different thermal histories and reached several conclusions

as to the effect of thermal treatments upon the morphology

of the glass. His results for different quenching

treatments were markedly different because of the

sensitivity of the metastable phase separation to

quenching rates. Analysis of his results gave infor-

mation on the morphological differences and the

connectivity of the various phases which was corrobo-

rated with replica electron microscopy.


Glass Structure


The current conception of the structure of glasses

stems largely from the x-ray diffraction studies of

Zachariasen (1932) and Warren (1937). Their work





- 16 -


indicated that in silicate glasses the average silicon-

oxygen distance is 1.62 A and the average number of

oxygens adjacent to a silicon atom is 4. These two

observations form the basis of a structural model of

tetrahedral SiO4 groups interlocking by sharing of

oxygen ions between adjacent tetrahedra. This model, or

models which closely resemble it, are generally accepted

as representative of most silicate glasses. A great

deal of literature exists in this area and is well

reviewed by Mackenzie (1960). Russian workers in the

glass structure area were first to propose a theory of a

microhetrogeneous glass structure. The early Russian

work exemplified by the work of Lebedev (1940), concludes

that the structure of glass is one of microcrystallites

with dimensions in the same size range as the interatomic

distances in crystals.

Glass structure on a larger scale (20 to 1,000 A)

is an area in which electron microscopy has led to a

detailed structure characterization. Electron

microscopy has shown evidence of liquid-liquid phase

separation in a variety of systems. The lithia-silica

and the soda-silica binaries are two systems which have

been shown to exhibit this type of behavior. The

pioneering work of Slayter (1952) and Prebus and

Michener (1954) has shown that silicate glasses contain





- 17


structural heterogenities in the size range of 20 to 200 A.

Subsequent investigations by Seward et al. (1967) and

Shaw and Uhlmann (1968) have shown that structural

features in the above size range are a general feature

of many glasses.

The work by Vogel and Byhan (1964) in lithia-

silica glasses has shown that most compositions in the

SiO2-Li20-2SiO2 phase field show structural heterogenities

whose existence and general behavior can be rationalized

by a metastable liquid-liquid separation. The existence

of the metastable miscibility gap can be inferred from

the "S" shaped liquidus on the phase diagram determined

by Kracek (1939).

Tran's (1965) work in soda-silica glasses has

shown a similar phase separation in glasses between 9

and 20 mole % soda. No evidence of separation was

observed in glasses of soda content greater than 19 mole %.

Subsequent work in both binary systems involving thermal

treatments in the annealing-transformation range has

shown the coarsening of the liquid-liquid separation

prior to the appearance of crystalline phases (Aver'yanov

and Koshits 1966).

The mechanism of liquid-liquid separation has

been investigated by various authors who have generally

agreed upon the nature of the separation. Depending upon





- 18 -


the composition of the glass, separation takes place

either by nucleation and growth or by spinodal decom-

position. Cahn (1968) has recently summarized the

thermodynamic arguments for spinodal decomposition as

well as discussed the various systems, both oxide and

metal, in which the mechanism has been reported. Cahn

and Charles (1965) have summarized the theory of phase

separation and applied their results to various glass

systems. Haller (1965), McCurrie and Douglas (1967) and

others have examined many of the systems discussed in

the above works and concluded that the observed structures

could be explained by a random nucleation and growth

process.











CHAPTER II

EXPERIMENTAL PROCEDURE

Glass Preparation

The objective of this work was to examine the

effect of thermal treatment upon the electrical properties

of alkali-silicate glasses and to correlate those

properties with structural changes occurring in the

glasses. To accomplish this objective it is necessary

that the initial "state" of the glasses be well defined or

at least invariant insofar as the parameters used to

characterize the glasses could distinguish. With this in

mind an experiment to examine the possible glass prepa-

ration variables such as melting time, temperature of

melting, pouring temperature, alkali vaporization, quenching

rate, annealing time and annealing temperature was designed.

The above variables were examined by pouring a

series of 30 mole % Li2O-SiO2 glass samples as shown

schematically in Figure 1. The diagram depicts the

process of melting a glass for 24 hours at 13500C,

pouring several samples at two quenching rates followed

by annealing of these samples for periods of time from

1 to 12 hours. Next the melt temperature was increased

to 14500C and held for 24 hours and a similar set of


- 19 -





- 20 -


2 moad moia

b. 300'C
(Annealing)
E

I I I
24 48 72
Time (hours)





Figure l.--Schematic melting procedure for 30 mole %
lithia-silica glass.





- 21-


samples was poured. Following this the melt temperature

was reduced to 13500C, held for 24 hours, and a similar

set of samples poured.

Samples were selected from the above melt and

electrical property measurements were carried out by

methods discussed later. Each of the possible types of

samples exhibited AC and DC electrical properties which

were identical within the limits of experimental error

with all others. From this result it is concluded that

none of the above-mentioned variables,over the range

examined, affects the initial electrical properties of the

30 mole % lithia-silica glass.

The possibility that the initial state of the

glass is different under the above-described conditions,

but is not detected by the techniques employed, was

examined by the following technique. Samples from each

class of specimens described above were thermally treated

and their electrical properties were re-examined. Once

again the electrical properties of the various types of

samples were identical even though the treatment had

changed the properties of the entire group. The property

changes of the group are discussed later. It is thus

concluded that variations of the variables set forth

above within the range examined do not affect the initial

structure of the glass.





- 22 -


X-ray fluorescence examination of the samples

melted for a total of 72 hours in a platinum crucible

showed no evidence of platinum.

In consideration of the above conclusions,

preparation of glasses for the remainder of this study

were made in the following manner.

1. Glass batches from materials of purity shown
in Table 1 and total weight of approximately
1 kilogram were weighed to an accuracy of
0.1 gram, giving an expected composition
accuracy of at least 1 part in 103.

2. The batch was then mixed in a jar mill without
balls for a minimum of 1 hour.

3. A 150 milliliter platinum crucible was then
filled from the batch and placed in an
electrically heated silicon carbide element
furnace at 13500C 50C or 14500C 50C,
depending on the glass being melted.

4. After approximately 15 minutes the crucible
was removed and refilled as necessary until
the crucible was full.

5. The crucible was then covered with a platinum
lid and the melt held at temperature for a
minimum of 24 hours.

6. After 24 hours the lid was removed from the
crucible and the crucible was replaced in the
furnace to allow it to come back to 13500C
prior to glass pouring.

7. The glass was then poured in a 17.5 mm diameter
steel mold. A tightly fitted plunger was
pressed into the molten glass, giving a sample
thickness between 3 mm and 8 mm and a diameter
of 17.5 mm.





- 23 -


TABLE 1

ANALYSIS OF GLASS RAW MATERIALS



Compound Weight %

Sodium Carbonate'
Na2C03 99.8
Chloride .0005
Nitrogen .0005
Phosphate .0005
Sulfate .001
Arsenic .0001
Calcium and Magnesium .005
Iron .0002
Potassium .001
Silica .005
Heavy Metals .0002

Lithium Carbonate2
Li2Cos 99.3
Na2Co3 0.2
Iron .0008
Sulfate 0.3
Chloride .0003
Calcium .0003
Phosphate .0001

Silicon Dioxide3
Si02 99.91
Iron .019
Alumina .08
Titania .009
Calcium and Magnesium Trace


IBaker Reagent.

2Foote Mineral Company.


3Pennsylvania Glass Sand Company.





- 24 -


8. The resulting glass button was quickly removed
(after approximately 30 seconds for cooling)
and placed in a furnace at 3000C 50C where
it was held for 1 hour and air cooled.

9. After each sample was poured, the remaining
glass was placed in the furnace and allowed
to come back to the melting temperature, while
the steel mold was chilled in tap water to
keep it at the same temperature for each
sample.

It was found that 1 hour at 3000C was the minimum

treatment which allowed the samples to be cooled to room

temperature without breakage. The residual stresses

remaining after this treatment were measured by standard

birefringence techniques and found to be 3,000 to 5,000 psi.


Electrodes


For electrical measurements, the faces of the

samples were ground parallel to within 0.05 mm with

silicon carbide metallographic paper and polished with

600 grit metallographic paper. This surface was then

cleaned with distilled water and quickly dried. In order

to vapor deposit the desired double guard ring gold

electrodes on the faces of the samples, paper masks were

affixed to the sample faces. These paper masks with an

inner diameter of 12.7 mm and an outer diameter of 15.9 mm

were cut from a gummed paper label with a modified

machinist's compass. The resulting masks were moistened




- 25 -


with distilled water and carefully affixed to the samples

so that they were concentric with the sample itself. The

samples were then placed in a specially designed holder in

a standard vacuum metallizer and the vacuum system was

pumped down to less than 1 micron. The samples were then

plated with gold from a tungsten basket placed about

100 mm from the sample. Total electrode thickness was

approximately 2 10-6 mm as determined from weighing

samples before and after plating. The samples were then

removed from the evaporator and placed in a furnace at

3000C 50C for 1 hour to increase the adherence of the

gold to the sample. Samples were removed from the

furnace, air cooled and the charred paper masks were

carefully removed, leaving the sample with gold electrodes

in a double guard ring configuration.


X-Ray


Samples for examination in the Guinier (1956) and

DeWolff (1947) x-ray camera were selected from the samples

poured for electrical measurements in order to insure that

their thermal history was identical to that for the

electrical samples. These samples were heat treated,

broken with a hammer and anvil, then ground in an alumina

mortar and pestle to pass a Tyler 200 mesh screen. The

ground glass-crystals were stored in a closed container

to prevent moisture pickup and resulting hydration.





- 26 -


The Nonius Guinier-DeWolff x-ray camera and copper

x-ray tube used for these experiments are shown schematically

in Figure 2. The Guinier-DeWolff camera, a vacuum path,

focusing quartz crystal monochromated powder camera, is

capable of examining 4 samples simultaneously. The samples

were held in 4 slots in a flat sample holder approximately

0.025 mm thick with Scotch #810 tape. The powder pattern

was recorded on Kodak Type NS double emulsion x-ray film,

developed 8 minutes at 230C in Kodak Type D-76 developer

and fixed 3 minutes at 230C in Kodak "Rapid Fix." It was

found that this technique gave the highest line intensity

with a tolerable fog level for a given exposure. Using

the above techniques, it was found that 0.1 weight %

lithium metasilicate crystal in a prepared lithium

disilicate glass standard could be detected. Detection

of crystals with this small weight fraction required a

50 hour exposure at 40 kilovolts and 20 milliamperes. The

resulting patterns were analyzed using a reader which

allowed the crystal lattice spacings to be read directly

from the film.


Thermal Expansion and Softening Points


Samples for thermal expansion measurements were

prepared by melting the glass in the same manner as for

the electrical samples. The dilatometric samples were





- 27 -


Film


Monochromator


X-ray /
Source Sample











Figure 2.--Schematic diagram of x-ray source and
Guinier-DeWolff camera.





- 28 -


poured in a steel mold which tapered from a diameter of

12.7 mm to 15.9 mm in its five centimeter length. This

taper was necessary to allow rapid removal of the sample

from the mold to prevent breakage. Following pouring,

the samples were held in a furnace at 300C for 1 hour

to prevent breakage, then removed and air cooled. Samples

prepared in this manner were too highly strained to be

cut with a cutoff wheel so they were all heat treated

1/2 hour at their subsequent heat treatment temperature

to further remove strains. After this treatment the

samples were cut to slightly over 50.8 mm length with a

water cooled silicon carbide wheel and the ends were

polished to 50.80 mm .02 mm on 180 silicon carbide

metallographic paper.

The thermal expansions of the various thermally

treated glasses were measured in an Orton recording

dilatometer which is a quartz tube and push rod apparatus

with a linear variable differential transformer transducer

for measuring the expansion. Samples were separated from

the quartz tube and push rod with 0.025 mm platinum foil

to prevent reaction with the quartz. The apparatus was

set up to plot the thermal expansion curve as a continuous

function of the temperature.




- 29 -


Electron Microscopy


Replica Preparation

Samples to be examined by electron microscopy were

prepared in the same manner as the electrical samples and

heat treated in the bulk state. After thermal treatments,

the samples were fractured to expose a fresh surface from

the interior of the sample. The fracture surface thus

obtained was etched 1 minute in an aqueous 5 volume %

hydrofluoric acid solution. The samples were then placed

in an evaporator and a platinum preshadow was applied,

followed by a carbon film. The carbon film replica was

removed from the glass by immersing the sample in an

aqueous 2% hydrofluoric acid solution and allowing the

replica to float off the sample. This replica was placed

on a copper microscope grid and washed several times in

distilled water to remove the acid. Replicas were

examined in a Phillips EM 200 electron microscope using

standard techniques.


Transmission Preparation

Samples for examination by transmission electron

microscopy were poured and heat treated in the same manner

as the electrical samples. Samples were mechanically

thinned to approximately 0.2 mm prior to chemical thinning.

The mechanical thinning was accomplished by grinding and





- 30 -


polishing one face of the poured button, cementing the

polished face to a small flat piece of steel and grinding

to the final thickness. Samples of 3 mm diameter were

then cut from the thinned material while it was still

fixed on the steel block. The samples were cut by using

a hollow copper drill (3 mm ID) with a slurry of 400 grit

silicon carbide in a drill press. The resulting sample

blanks were removed from the steel block and the cement

was removed using ethylene dicloride in an ultrasonic

cleaner.

The chemical thinning was carried out by dimpling

the sample in the center, followed bya final thinning

operation. The dimpling operation was accomplished by

masking the outer edges of the samples with a lacquer

("Microstop") which did not allow the edges of the sample

to be attacked. The solution used for this operation was

made up of 10 parts hydrofluoric acid (48%), 5 parts

nitric acid and 14 parts acetic acid by volume. After the

dimpling operation, the masking material was removed in

acetone and the sample was carefully washed. At this

point the sample has a relatively thick edge with a thin

interior region which greatly facilitated handling. Final

thinning was accomplished by alternately dipping the

sample in hydrofluoric acid for short periods of time

(30 to 60 seconds) and examining the center portion for holes.





- 31 -


When the first visible hole developed, etching was stopped

and the sample was washed several times in distilled water.

The sample was then placed in an evaporator and a thin

film of carbon was evaporated on one surface to prevent

charging of the sample by the electron beam. The sample

was then placed in the heating stage holder (Phillips

PW 6560) with a large platinum aperture to facilitate

heat conduction.

The sample holder was then placed in the Phillips

EM-200 electron microscope with a rotating-tilting stage

and heated to the required observation temperature. The

remainder of the electron microscopy was then conducted

by standard techniques.


Electrical Measurements


Sample Chamber

DC and AC properties were measured over a range of

temperatures in a vacuum environment. The requirements

of electrical shielding in AC measurements and guarding in

DC measurements was of major importance in the sample

holder design. The measurements require that the leads

and contacts be made in a coaxial configuration. The

sample holder design shown in Figure 3 used in these

experiments utilizes a coaxial arrangement of leads as

far as possible. Where this is not possible, high quality




- 32 -


3f
:i


N


Figure 3.--Photograph of sample holder
employed in electrical measurements.




- 33 -


insulation with resistivity greater than 1 1014 ohm cm

at room temperature was used. Electrical contacts to the

samples and leads out of the sample chamber were platinum

to avoid oxidation and thermoelectric problems. The sample

chamber was vacuum sealed to allow all electrical

measurements to be conducted in a vacuum of less than

1 micron. The sample chamber and sample were heated by

means of an external nichrcme heating element in con-

junction with a Variac for temperature control. The

sample temperature, as measured with a chromel-alumel

thermocouple placed approximately 1 mm from the sample,

was constant over a series of electrical measurements to

10C of the set point


AC Measurements

Discussion of the electrical equipment is sub-

divided into the audio frequency (AF) equipment, the radio

frequency (RF) equipment and the null detection system.

The AF equipment, shown schematically in Figure 4, consisted

of a Wayne-Kerr B-221 transformer radio arm bridge in

conjunction with a Hewlett-Packard 651-A oscillator. A

schematic diagram of the bridge circuit is shown in

Figure 5 where ZS and Z are the standard and unknown

impedances respectively. The balance condition is

satisfied, as indicated by a null on the detector, when





- 34 -


AUDIO FREQUENCY MEASUREMENT APPARATUS

20HZ TO 20KHZ



HP-651A WAYNE-KERR
--- B-221 NULL SIGNAL
OSCILLATOR A.F BRIDGE



SAMPLE
CHAMBER








Figure 4.--Schematic diagram of audio frequency
measurement equipment.




- 35 -


SOURCE


Figure 5.--Schematic diagram of transformer ratio
arm bridge.





- 36 -


equal currents flow in each half of the center tapped

transformer (T2). When this condition is satisfied, the

potential on the primary will be zero and the right hand

terminals of the unknown and standard will be at neutral

potential. The same voltage is applied to both unknown

and standard and for equal currents to flow in each half

of the primary of transformer T2; the real and imaginary

parts of the unknown impedance must be equal to those of

the standard. The instrument is designed to allow values

of the resistive and capacitative component to be read

directly from the instrument dials.

The RF measuring equipment, shown in Figure 6,

consisted of a Wayne Kerr B-601 bridge and the same

oscillator as was used with the AF equipment. The bridge

design is, in theory, similar to the AF bridge except that

the transformers and standards used are designed for use

in the RF range.

The null detection system, shown in Figure 7,

consisted of a General Radio 1232-A null detector used in

conjunction with a General Radio 1232-Pl RF mixer and a

Wayne Kerr 0-22-D beat frequency oscillator. The 1232-A

covers the frequency range 20 Hz to 100 KHz directly, so

that the AF and RF null signals up to 100 KHz were

detected directly. At frequencies greater than 100 KHz,

it was necessary to use a beat frequency technique to





- 37 -


RADIO FREQUENCY MEASUREMENT APPARATUS


20KHZ TO


IOMHZ


NULL SIGNAL


Figure 6.--Schematic diagram of radio frequency
measurement equipment.





- 38 -


DETECTOR SYSTEM


GENERAL RADIO
1232-A
NULL DETECTOR


A.F NULL SIGNAL


Figure 7.--Schematic diagram of null detection system.




- 39 -


reduce the signal frequency to the range of the 1232-A

detector. This was accomplished by mixing a local

oscillator signal of frequency 100 KHz greater than the

measurement frequency with the signal to be detected.

This gives a beat frequency signal proportional to the

original signal at a frequency of 100 KHz which can be

detected by the 1232-A detector.


DC Measurements

DC measurements were made with the equipment shown

schematically in Figure 8. The short time measurements

were conducted by displaying a signal proportional to the

current flowing in the sample on a Hewlett-Packard 140 A

oscilloscope with a 1420 A time base and 1402 A dual trace

amplifier and photographing the oscilloscope trace with

a Tektronix C-12 camera. This technique allowed the

sample conductivity to be measured in a time of less than

5 milliseconds in most cases. The minimum time depended

upon the range of the Keithly 416 high speed picoammeter

and no measurements were made in times less than the

response time of the picoammeter.

The DC potential for the measurements was furnished

by a Hewlett-Packard 6217 power supply which has a voltage

stability of less than 0.10% + 5 millivolts in 8 hours

with less than 200 microvolt AC ripple. The connections





- 40 -


D C MEASUREMENT APPARATUS


TEKTRONIX C12
cSCILLOSCOP'E
CAMERA


r ----------------- I
HP-412-A I
VACUUM TUBE VOLTMETERI
_CURRENT MODE) _












Figure 8.--Schematic diagram of DC measurement
equipment.




41 -




to the sample and picoammeter were made with guarded

leads as indicated in Figure 8. Sample currents greater

than 3 10-5 amps were measured on a Hewlett-Packard

412-A vacuum tube voltmeter and because of the large

time constant of this instrument, no measurements could

be made in times less than 1 second.











CHAPTER III

POLARIZATION


The objective of the polarization experiments was

to establish the behavior of the sample under a DC field

in order that the electrode polarization could be elimi-

nated as a cause of relaxation phenomena in the AC

measurements. This objective requires that the time

required to polarize the sample electrodes be greater

than the maximum period (minimum frequency) used in the

AC measurements.

Polarization results from a 30 mole % lithia glass

in the as cast state are presented in Figure 9. These

results are a combination of DC and AC conductivity

measurements. The "time" for the AC measurements was

taken as the reciprocal of the measurement frequency. This

definition of the time parameter results in a reasonably

good fit in the region of overlap of the two curves and is

thus considered to be a valid definition. The agreement

of the two curves is in fact remarkably good considering

the fact that in the overlapping region both measurement

techniques are approaching their time limits and the

accuracy is subject to larger errors in that time domain.


- 42 -





- 43 -


Time ISecondsl


Figure 9.--DC polarization curve of 30 mole % lithia-
silica glass in as cast and heat treated 5 hours at
5000C forms.




- 44 -


The general shape of the curves for the as cast

glass, as expected, shows increasing conductivity in the

short time (10-4 to 10-5 second) range, a flat intermediate

region (10-s to 10-1 second) and a region of electrode

polarization for times in the 10-1 to 10+3 second range.

The curve corresponding to the sample heat treated

5 hours at 5000C exhibits more structure than the as cast

glass. The short time and initial flat region is quite

similar to the as cast behavior up to about 2 10-3

second. At that point a dispersion appears and at longer

times the flat tries to reappear but is masked by the

appearance of electrode polarization. Discussion of the

dispersion in the heat treated sample will be presented

in the AC properties section for reasons that will be

evident later.

The appearance of the polarization problem is

illustrated by the series of conductivity-time plots in

Figure 10. It is evident that at low temperatures, when

the conductivity is low, the problem is unimportant but

at high temperatures the problem is very serious. The

effect of the polarization upon the measured DC conductivity

is shown in a more conventional form, the log conductivity

versus reciprocal temperature plot, for two arbitrarily

chosen times of measurement (approximately 1 second and

60 seconds) as shown in Figure 11. The problem of




- 45 -


TIME (Seconds)


Figure 10.--Conductivity versus log time
curve for 33 mole % lithia-silica glass
as cast.





- 46 -


lol T IC)


Figure ll.--Log DC conductivity versus
reciprocal temperature for 33 mole %
lithia-silica glass as cast.





- 47 -


polarization is less important in the low temperature

region but at higher temperatures the problem in its

most serious state causes an inversion in the slope of

the curve. It is thus evident that DC measurements made

without consideration of the problem of polarization are

virtually meaningless if they are measured over a

temperature-time range where the phenomeron shown above

is significant.

In order to develop a working criterion for

determining the effect of polarization on conductivity

measurements, a model consisting of a plane parallel

capacitor with mobile positive charges in the dielectric

and totally blocking electrodes was chosen. This model

is approximately the situation in an alkali silicate

glass capacitor. With a potential applied, this capacitor

(Figure 12) will have a surface charge equal to the

negative of the volume charge assumed to be uniformly

dispersed in the dielectric. It is assumed that total

polarization occurs when the back voltage due to the

surface charge is equal to the applied voltage and the

apparent conductivity is zero.

To allow comparison of the surface charge calcu-

lated from the model above with an experimentally derived

value, it was necessary to define electrode polarization

on the basis of the conductivity time curves. The criterion




- 48 -


CATHODE
UNIFORM POSITIVE
SURFACE CHARGE
UNIFORM NEGATIVE
VOL UME CHARGE
ANODE









Figure 12.--Schematic diagram of model used for
polarization calculation.




- 49 -


chosen was to assume that a conductivity drop of one full

order of magnitude represented total polarization of the

sample. The charge transported during the time required

for polarization to occur will be calculated from the

experimental results and compared with that predicted from

the model.

The electric field in the capacitor due to the

uniform surface charge on one electrode is given by

equation (3-1) from Page and Adams (1958).


E P (3-1)


The field due to the uniform negative charge of equal

magnitude but opposite sign to the surface charge is

given by equation (3-2) (Page and Adams 1958).


P x 2)
E = (3-2)


In order to obtain the potential between the two plates,

the two fields above are added and integrated over the

electrode spacing. As a result of symmetry, the uniform

volume charge does not contribute to the potential thus

the integration of (3-1) gives equation (3-3).


pt (3-3)
E: E:





- 50 -


Equation (3-3) gives the magnitude of the back potential

corresponding to a given polarization charge on or adjacent

to the cathodic electrode. It is thus possible to calcu-

late from equation (3-3) the magnitude of the surface

charge from the applied potential, sample thickness and

the dielectric constant obtained from AC measurements.

Substitution of the above values from the sample used to

obtain the polarization curves in Figure 10 yields a

surface charge of 9.80 10-4 coulombs/cm2.

In order to obtain an experimental value of the

surface charge, it is assumed that all the charge

transported is left on the electrodes. The charge

transported is the time integral of the current over the

time of polarization. Because the analytical form of the

current time behavior is not known, the integral above is

approximated by assuming that the current drops from its

initial value to zero in the time interval considered.

This assumption can be justified by reference to

Figure 10. The conductivity-time and current-time curves

have the same shape because of Ohm's law; hence Figure 10

can be looked upon as a current-time plot. The value of

the current transported during the polarization depicted

in the 1800C curve of Figure 10 calculated in the manner

described above yields a surface charge of 1.33 10-3


coulombs/cm2.





- 51 -


Comparing this value with the value calculated

from the model above, it is evident that the agreement is

good considering the approximations made. There are, at

least, two possible reasons for the differences. The

electrodes have been assumed to be totally blocking to

the charge carrier but it is possible that this is not

the case and part of the charge transported is not on the

surface of the dielectric. Another possible source of

the difference is the approximation of the shape of the

current-time curve. The approximation of the curve shape

in Figure 10 yields a value of the current which is too

high. These two factors are sufficient to explain the

observed difference and the experimental measurement.

The results above indicate that the problem of

polarization can be approached by two methods. The first

and most direct consists of measuring the polarization

behavior of each sample and analyzing the conductivities

directly from the polarization results. The second and

faster method is to calculate the conductivity required

to cause polarization in the measurement time used. This

would allow prediction of cases in which polarization

difficulties will arise. Clearly both techniques have some

inherent difficulties but either one or both methods should

be used to eliminate or reveal polarization problems in

DC measurements.





52 -




The DC measurements reported herein were made by

measuring representative polarization curves and taking

the conductivity values from the time independent portions

of those curves.











CHAPTER IV

EXPERIMENTAL RESULTS


X-Ray


The objective of the x-ray studies was to detect

and identify the crystals appearing in the initial stages

of crystallization. The initial x-ray studies were

conducted to determine the lower limits of detection of

crystalline phases in a glassy matrix. Examination of a

series of lithium disilicate glasses mixed with 5.0, 1.0,

0.5 and 0.1 weight % crystalline lithium metasilicate

showed that the crystals in a 0.1 weight % standard could

be observed. A similar experiment with crystalline

lithium disilicate in lithium disilicate glass gave the

same lower limit of detection.

Results of the x-ray examination of the 30 mole %

lithia-silica glass following various heat treatments at

500C are presented in Table 2. The actual patterns are

not presented because the lines of interest are very weak

even after 50 hour x-ray exposures and photographic

reproduction is difficult. The phases present change with

thermal treatment at 5000C from the glassy material in

the as cast form, to a glassy material plus a phase

tentatively identified as crystalline lithium metasilicate


- 53 -





- 54


TABLE 2


SUMMARY OF


X-RAY LATTICE SPACINGS AND RELATIVE INTENSITIES
FOR 30 MOLE % LITHIA GLASSES


30 mole % Li2O 30 mole % Li2O
Standard Standard Heat Treated Heat Treated
Li20O2SiO2 Li20OSiO2 5 hrs at 5000C 50 hrs at 5000C

1 (A) Intensity d (A) Intensity d (A) Intensity d (A) Intensity


2

10

1(M.S.)



1

10

10

10

1 (M.S.)

1



5



1 (M.S.)

5

5

2


4.70











3.30









2.70



2.35


2.08


4.70

4.35

4.18







3.30



2.91



2.81

2.70


5.5







3.75

3.65

3.58


LA L .1


5.45


$.75

$.65

L.58

[.30

?.95


?.39


?.27





- 55 -


with a weight fraction of approximately 0.1 weight %.

This 5 hour x-ray pattern shows seven lines, three of

which correspond to the stronger lithium metasilicate

crystal lines. The other four lines are weaker and, at

present, not conclusively identified. It appears that

the unidentified lines correspond to a transition phase

of lower symmetry than the orthorhombic metasilicate.

Indexing is difficult, if not impossible, because of the

small number of lines observed. Examination of the 10

and 20 hour samples show only a diffuse peak character-

istic of the glassy state with no evidence of crystalline

phases. The 50 hour sample exhibits the four strongest lines

corresponding to the equilibrium lithium disilicate crystal.

X-ray results for the 33 mole % lithia (lithium

disilicate) glass are shown in Table 3. The reaction

sequence observed in this glass is that of the glassy

material transforming to the equilibrium lithium

disilicate crystal with thermal treatment. The equilibrium

lithium disilicate precipitate is observed after 50 hours

at 5000C.

X-ray examination of the 26.4 mole % lithia glass

did not reveal the presence of any crystalline phases with

thermal treatments up to 50 hours at 5000C.





- 56 -


TABLE 3

SUYVMARY OF X-RAY LATTICE SPACINGS AND RELATIVE
INTENSITIES FOR 33 MOLE % LITHIA-SILICA
GLASSES


Standard Crystal 33 mole % Li20 33 mole % Li20
Li20-2SiO2 50 hrs at 5000C 100 hrs at 5000C

d (A) Intensity i d (A) Intensity d (A) Intensity

7.4 2 7.4 2

5.45 10 5.45 1 5.45 3

4.18 1

3.75 10 3.75 2 3.75 10

3.65 10 I 3.65 2 3.65 10

3.58 10 3.58 2 3.58 10

2.95 1

2.90 5 2.90 3

2.39 5 2.39 3

2.35 5 2.35 1 2.35 3

2.27 2





- 57 -


The results of x-ray examination of the 33 mole %

soda (sodium disilicate) glass are presented in Table 4.

These results show the appearance of the equilibrium

sodium disilicate crystal after 10 hours at 5500C.

Further treatments at this temperature only increase the

amount of crystalline sodium disilicate present in the

glass.

The x-ray results of the 25 mole % soda glass do

not show evidence of crystallization with treatments of

50 hours at 5500C.

Electron Microscopy

The objective of the electron microscopy was to

identify and determine the morphology of the phases

appearing during the initial stages of crystallization.

Replica Techniques

Replica electron micrographs of the 30 mole %

lithia glass are presented in Figures 13 to 17. The

micrograph of the as cast glass (Figure 13) exhibits a

droplike structure with separated regions in the size

range 0.1 to 0.5 micron.

It is known (Vogel and Byhan 1964) that silica

rich glasses are attacked by hydrofluoric acid much less

rapidly than lithia rich silicate glasses. From this

information and a knowledge of the shadowing direction,

it is evident that the drop regions in Figure 13 are richer

in silicon than the surrounding matrix.




- 58


TABLE 4

SUMMARY OF X-RAY LATTICE SPACINGS AND RELATIVE
INTENSITIES FOR 33 MOLE % SODA-SILICA GLASSES


Standard Crystal 33 mole % Na20 33 mole % Na20
Na20-2SiO2 50 hrs at 5000C 10 hrs at 550C

d (A) Intensity d (A) Intensity d (A) Intensity


10.8
9.5
6.1
5.90
4.95
4.50
4.20
4.16
3.94
3.85
3.77
3.71
3.42
3.30
3.26
3.22
3.10
3.03
2.95
2.93


6.0
5.85
4.90


4.18





3.75


3.40


3.28
3.22


2.99


2.93


3.75





3.28








2.54
2.41







- 59 -


Figure 13.--Replica electron micrograph of 30 mole %
lithia-silica glass (30,000x) as cast.





- 60 -


**~;P M


Figure 14.--Replica electron micrograph of 30 mole %
lithia-silica glass (26,000x) after 5 hours at
5000C.




- 61 -


Figure 15.--Replica electron micrograph of 30 mole %
lithia-silica glass (31.,000x) after 10 hours at
5000C.




- 62 -


Figure 16.--Replica electron micrograph of 30 mole %
lithia-silica glass (31,000x) after 20 hours at
5000C.




- 63 -


Figure 17.--Replica electron micrograph of 30 mole %
lithia-silica glass (37,000x) after 50 hours at
5000C.





- 64 -


The micrograph of the 30 mole % sample heat

treated 5 hours at 5000C is shown in Figure 14. This

micrograph exhibits the separated drop regions character-

istic of the as cast sample and further shows a secondary

drop-like separation in the matrix surrounding the primary

drops. This secondary separation in the matrix is the

separation of the lithium metasilicate in the matrix.

Figure 15 corresponds to the 30 mole % glass heat

treated at 5000C for 10 hours. This micrograph shows

some evidence of the' initial droplets but they are not as

pronounced as in the earlier thermal treatments. The

matrix exhibits a fine scale separated drop morphology

similar to that in the 5 hour treatment. Thus, the drop

regions are beginning to dissolve with thermal treatment.

The micrograph of the 30 mole % glass heat treated

20 hours at 5000C is shown in Figure 16. This micrograph

exhibits an overall uniformity of structure with only

traces of the primary drop separation.

The micrograph of a sample heat treated 50 hours

at 5000C is shown in Figure 17. This micrograph shows

remnants of the structure present at 5 hours but it is

generally more homogeneous than all the other glasses

examined in the above series. The raised regions are

the equilibrium lithium disilicate precipitate.





- 65 -


Transmission Electron Microscopy

Diffraction Patterns

The electron diffraction patterns taken during a

thermal treatment at 4800C are presented in Figures 18

to 23. The patterns initially show no crystallinity

followed by the appearance of lithium metasilicate and

lithium disilicate crystals. The interplanar spacings

obtained from the diffraction patterns and tabulated in

Table 5 indicate that during the early stages of

crystallization the crystalline lithium metasilicate

predominates, but in the latter stages the lithium

disilicate predominates.

The lattice spacings taken from the ASTM card file

(Smith 1967) are those corresponding to room temperature,

while those discussed above correspond to 4800C. The

results of Glaser (1967) with high temperature x-ray

diffraction on the lithium disilicate indicate that the

expected expansion in the {130} lattice spacing for example

is less than one percent. This is within the expected

experimental error in the electron diffraction measurements

so the comparison with room temperature lattice spacings

is a valid one.

Transmission Electron Micrographs

Figure 24 is a micrograph of the 30 mole % glass

before thermal treatment. This micrograph exhibits a

network structure which is undoubtedly a result of the




- 66 -


Figure 18.--Electron diffraction pattern of as cast
30 mole % lithia-silica glass at room temperature.




- 67 -


Figure 19.--Electron diffraction pattern of 30 mole %
lithia-silica glass at 4800C after 5 minutes at
4800C.




- 68 -


Figure 20.--Electron diffraction pattern of 30 mole %
lithia-silica glass at 480C after 14 minutes at
4800C.




- 69 -


Figure 21.--Electron diffraction pattern of 30 mole %
lithia-silica glass at 4800C after 26 minutes at
4800C.




- 70 -


Figure 22.--Electron diffraction pattern of 30 mole %
lithia-silica glass at 4800C after 2.5 hours at 4800C.




























F-_.ure 23.--Electron diffraction pattern of 30 mole %
lithia-silica glass at 480QC after 3.75 hours at 4800C.





- 71 -


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EH









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Ir 0 --
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ri-
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Ln co
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U~C) C) ) ) ) C


t t


rc5.
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H ii
cr f
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- 72 -


u~ i~ u~ ci u c~ i~ i I I
aaa~aaI I Il


C -

1 4
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;' (


on H on
-D 0 0
H 0O
*H *


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rH

4-1 Q
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0m
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CD OrI m


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c<


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2S S22 S3 :S >


m m m m m U1 m m
Q a Q Q a z z





- 73 -


-'b6








i
















Figure 24.--Transmission electron micrograph 30 mole %
lithia-silica glass at room temperature as cast
(41,000x).





- 74 -


carbon deposited on the glass surface. The carbon layer

partially masks the glass structure but the diffraction

pattern leaves no doubt that some of the structure present

is crystalline metasilicate. The carbon surface film is

discussed further in the following section. Figure 25 is

a micrograph taken at 480C after 11 minutes at 480C. It

shows the appearance of black spots which are apparently

crystals in a glassy matrix. The diffraction pattern at

this point in time shows that the crystals present are

primarily lithium metasilicate.

Bright and dark field micrographs taken after 35

minutes at 480C are presented in Figure 26. The dark

field micrograph was taken from the {170} diffraction ring

of the lithium disilicate so that the bright crystals in

the dark field image are lithium disilicate crystals.

Micrographs taken after 2.3 hours at 4800C are shown in

Figure 27. The dark field image corresponds to the {170}

diffraction ring in the lithium disilicate pattern hence

all the bright areas in the dark field micrograph are

lithium disilicate crystals. The disilicate crystal size

has increased in the time elapsed between Figure 26b and

Figure 27b, indicating the growth of the lithium disilicate

crystals in the glassy matrix. The last micrograph in

this series at 480C, shown in Figure 28, was taken after

4.8 hours. This micrograph shows the development of an

elongated morphology from the previously equiaxed morphology.




- 75 -


Figure 25.--Transmission electron micrograph of
30 mole % lithia-silica glass at 4800C after 11
minutes at 4800C (42,000x).




- 76 -


la


Figure 26.--Transmissicn electron micrograph of
30 mole % lithia-silica glass at 4800C after 35
minutes at 480C. (a) Bright field (b) Dark
field (46,000x).


,ye ... .', .
;f



t, %1'?^ .- *o
.I. '-- .. *';4
+<.'^ <^*f.
*.; -%; A ?d M




- 77 -


Figure 27.--Transmission electron micrograph of
30 mole % lithia-silica glass at 4800C after 2.3
hours at 4800C. (a) Bright field (b) Dark field
(31,000x).





- 78 -


Figure 28.--Transmission electron micrograph of
30 mole % lithia-silica glass at 4800C after 4.8
hours at 480C (23,000x).





- 79 -


The effect of the carbon film on the surface of

these glasses was investigated by evaporating a carbon

film similar to that used on the glasses on a thin mica

crystal and examining the evolution of the structure with

thermal treatment. The carbon film on mica exhibited an

initial structure similar to that shown in Figure 24.

Heating of the mica-carbon film to 480C and following

the microstructural changes revealed a lower volume

fraction of essentially the same black spots character-

istic of the glass micrographs in Figures 25 and 26a.

It is thus concluded that some of the black spots must

be carbon.

It was not possible to obtain a dark field image

from the weak and diffuse lithium metasilicate rings

during the sequence. This difficulty indicates that the

observed black spots may be crystalline carbon but no

carbon diffraction rings were observed. This leaves the

identity of the black spots open to question but it

appears that some of the black spots must be lithium

metasilicate. It is possible that the carbon masks the

lithium metasilicate crystals and hence they are not

observed except in the diffraction pattern.

Figure 29a and 29b are micrographs taken at 500C

after 1 hour at 5000C. The dark field image is taken

from the area circled on the diffraction pattern in




- 80 -


I-
































Figure 29.--Transmission electron micrograph of
30 mole % lithia-silica glass at 5000C after 1
hour at 5000C. (a) Bright field (b) Dark field
(31,000x) .




- 81 -


Figure 30. The important feature of these micrographs is

the extremely long crystals growing in the glass matrix.

The growth of the whiskers out of the glass matrix is

also an interesting feature.

A room temperature transmission micrograph of a

glass heat treated in the bulk form for 5 hours at 5000C

is shown in Figure 31. There is a pronounced similarity

between this micrograph and the replica micrograph of a

similar sample shown in Figure 14.


AC Results


The objective of the AC measurements was to monitor

the initial stages of crystallization and characterize the

structure and structural changes occurring during this

period.


30 Mole % Lithia-Silica Glasses

The results of the AC measurements are presented

in the form of the AC loss angle (tan 6AC) as a function

of frequency and temperature. The tan 6AC values were

calculated from equation (4-1) following the method of

Charles (1963).


GTotal 1 1 1
tan 6 a- = tan 6AC + tan 6 +
Total WC AC DC C RAC RDC


(4-1)




- 82 -


Figure 30.--Electron diffraction pattern of 30 mole %
lithia-silica glass at 500C after 1 hour at 5000C.




- 83 -


Figure 31.--Transmission electron micrograph of
30 mole % lithia-silica glass at room temperature
after 5 hours at 500C in bulk form (31,000x)





- 84 -


The tangent of the loss angle is shown as a function of log

frequency for the various heat treatments at 5000C in

Figures 32-36. It can be noted from Figure 32 that the

as cast glass is free of loss peaks, but after a 5 1/2

hour heat treatment (Figure 33) at 5000C, large loss peaks

have appeared. The tolerance on the time of appearance for

the loss peaks was established by heat treating a sample in

1/2 hour increments and measuring the AC properties following

each heat treatment. Further heat treatments at the same

temperature for times up to 50 hours (Figures 34, 35 and

36) cause the magnitude of tan 6AC to decrease and the

peak location to shift to higher frequencies.

The temperature dependence of the frequency maxima

of the tan 6AC curves is shown in Figure 37. The frequency

maxima curves have the Arrhenius form typical of a thermally

activated process. The activation energy of the frequency

maxima-reciprocal temperature curves is 14.7 kilocalories

mole in all cases shown in Figure 37. The equivalence of

the loss process activation energy and the DC conductivity

activation energy indicates that the loss process is the

result of ionic motion. The loss behavior results are

summarized in Figure 38.

33 Mole Z Lithia-Silica Glasses

Results for the 33 mole % lithia glass are summarized

in Figure 39. The loss spectrum of this glass shows the

appearance of the loss peak between 2 and 5 hours heat




- 85 -


|

SNO HEAT TFEAT

122 --o- -
32
02 -o--- I


2i i







SI.
24 I
so




s V

4 1
\l I.






_______ I 0" -.-.. .. .. ... -.-- ....-
2 1 4 6 e

LOG, FREQUENCY







Figure 32.--Tan 6 versus logo frequency for the
30 mole % lithia-silica glass as cast.




Full Text

PAGE 1

THE EARLY STAGES OF CRYSTALLIZATION IN ALKALI-SILICATE GLASSES By DONALD L. KINSER A DISSERTATION PRESENTED TO THE GRADUATE CdUNCIL OF THE UNIVERSITY OF FLORIDA IN PARTIAL FULFILLMENT OF THE REQUIREMENTS FOR THE DEGREE OF DOCTOR OF PHILOSOPHY UNIVERSITY OF FLORIDA 1968

PAGE 2

UNIVERSITY OF FLORIDA 3 1262 08552 3305

PAGE 3

Dedicated to Piy wife, Barbara

PAGE 4

ACKNOWLEDGEMENTS The author would ].ike to acknowledge the assistance of his advisory committee, Drs. Rhines, DeHoff, Reid and Carr, and especially his chairman Dr. Hench v7ho contributed a great deal in the form of discussions, advice and encouragement throughout the course of this investigation. The numerous discussions with Dr. Hren are also gratefully acknowledged . Thanks are also due to numerous members of the faculty, staff and students of the Metallurgy Department for many stimulating discussions and a great deal of experimental assistance. The aid of Mr. E. J. Jenkins with electron microscopy techniques and Mr. D. A. Jenkins with many experimental problems is gratefully acknowledged. The author is also grateful for the financial support of the National Science Foundation and the United States Air Force without whose support this research would not have been possible. Ill

PAGE 5

TABLE OF CONTENTS Page ACKNOWLEDGEMENTS iii LIST OF TABLES V LIST OF FIGURES vi NOMENCLATURE xi ABSTRACT xiii Chapter I . INTRODUCTION 1 II . EXPERIMENTAL PROCEDURE 19 III. POLARIZATION 42 IV. EXPERIMENTAL RESULTS 53 V. DISCUSSION OF RESULTS 108 VI . SUM>L?iRY 130 VII . CONCLUSIONS 134 VIII. recommendation:^ to future INVESTIGATORS 136 BIBLIOGRAPHY 138 BIOGRAPHICAL SKETCH 143 IV

PAGE 6

LIST OF TABLES Table Page 1 . ANALYSIS OF GLASS RAW MATERIALS 23 2. SUMMARY OF X-RAY LATTICE SPACINGS AND RELATIVE INTENSITIES FOR 3 MOLE % LITHIA GLASSES 54 3. SUMMARY OF X-RAY LATTICE SPACINGS AND RELATIVE INTENSITIES FOR 3 MOLE % LITHIA-SILICA GLASSES 56 4. SW-IMARY OF X-RAY LATTICE SPACINGS AND RELATIVE INTENSITIES FOR 33 MOLE % SODA-SILICA GLASSES 58 5. SUMMARY OF LATTICE SPACINGS OBTAINED FROM ELECTRON DIFFRACTION PATTERNS OF 30 MOLE % LITHIA GLASS AT 480°C 71 6. SUMiMARY OF COEFFICIENTS OF LINEAR EXPANSION AND LOWER TRANSFORMATION POINTS 107

PAGE 7

LIST OF FIGURES Figure Page 1. Schematic melting procedure for 30 mole % lithia-silica glass 20 2. Schematic diagram of x-ray source and Guinier-DeWolf f camera 27 3. Photograph of sample holder employed in electrical measurements 32 4. Schematic diagram of audio frequency measurement equipment 34 5. Schematic diagram of transformer ratio arm bridge 35 6. Schematic diagram of radio frequency measurement equipm.ent 37 7. Schematic diagram of null detection system 38 8. Schematic diagram of DC measurement equipment 4 9. DC polarization curve of 3 mole % lithia-silica glass in as cast and heat treated 5 hours at 500°C forms 43 10. Conductivity versus log time curve for 33 mole % lithia-silica glass as cast 45 11. Log DC conductivity versus reciprocal temperature for 3 3 mole % lithiasilica glass as cast 46 12. Schematic diagram of model used for polarization calculation 48 13. Replica electron micrograph of 30 mole % lithia-silica glass (30,000x) as cast 59 VI

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LIST OF FIGURES — Continued Figure Page 14. Replica electron micrograph of 30 mole % lithia-silica glass (26,000x) after 5 hours at 500°C 60 15. Replica electron micrograph of 30 mole % lithia-silica glass (31,000x) after 10 hours at 500°C 61 16. Replica electron micrograph of 30 mole % lithia-silica glass (31,000x) after 20 hours at 500°C 62 17. Replica electron micrograph of 30 mole % lithia-silica glass (37,000x) after 50 hours at 500°C 63 18. Electron diffraction pattern of as cast 30 mole % lithia-silica glass at room temperature 6 6 19. Electron diffraction pattern of 30 mole % lithia-silica glass at 480°C after 5 minutes at 480°C 67 20. ElecTiron diffraction pattern of 30 mole % lithia-silica glass at 480°C after 14 minutes at 480°C 68 21. Electron diffraction pattern of 30 mole % lithia-silica glass at 480°C after 26 minutes at 480°C 69 22. Electron diffraction pattern of 30 mole % lithia-silica glass at 480°C after 2.5 hours at 480°C 70 23. Electron diffraction pattern of 30 mole % lithia-silica glass at 480°C after 3.75 hours at 480°C 70 24. Transmission electron micrograph 30 mole % lithia-silica glass at room temperature as cast (41,000x) 73 VI 1

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LIST OF FIGURES — Continued Figure Page 25. Transmission electron micrograph of 30 mole % lithia-silica glass at 480°C after 11 minutes at 480°C (42,000x) 75 26. Transmission electron micrograph of 30 mole % lithia-silica glass at 480°C after 35 minutes at 480°C 76 27. Transmission electron micrograph of 30 mole % lithia-silica glass at 480''C after 2.3 hours at 480°C 77 28. Transmission electron micrograph of 30 mole % lithia-silica glass at 480°C after 4.8 hours at 480°C (23,000x) 78 29. Transmission electron micrograph of 30 mole % lithia-silica glass at 500°C after 1 hour at 500°C 80 30. Electron diffraction pattern of 30 mole % lithia-silica glass at 500°C after 1 hour at 500°C 82 31. Transmission electron micrograph of 30 mole % lithia-silica glass at room temperature after 5 hours at 500°C in bulk form (31,000x) 83 32. Tan 6 versus log i o frequency for the 30 mole % lithia-silica glass as cast .... 85 33. Tan 6 versus logio frequency for the 30 mole % lithia-silica glass after 5 hours at 500°C 86 34. Tan 5 versus log i o frequency for the 30 mole % lithia-silica glass after 10 hours at 500°C 87 Vill

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LIST OF FIGURES — Continued Figure Page 35. Tan 6 versus logio frequency for the 30 mole % lithia-silica glass after 20 hours at 500°C 88 36. Tan 6 versus log i o frequency for the 30 mole % lithia-silica glass after 50 hours at 500°C 89 37. Log of frequency maxima versus reciprocal temperature for various thermal treatments at 500°C 90 38. Tan 6 versus logio frequency for the 30 mole % lithia-silica glass for various thermal treatments measured at 80°C 91 39. Tan 6 versus logio frequency for the 33 mole % lithia-silica glass for various thermal treatments 92 40. Tan 6 versus logio frequency for the 26.4 mole lithia-silica glass for various thermal treatments 94 41.' Tan 5 versus logio frequency for the 33 mole % soda-silica glass for various thermal treatments 95 42. Tan 6 versus logio frequency for the 25 mole % soda-silica glass for various thermal treatments 96 43. Log DC conductivity versus reciprocal temperature for the 3 mole % lithiasilica glass 98 44. Log DC conductivity versus reciprocal temperature for the 33 mole % lithiasilica glass 100 IX

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LIST OF FIGURES — Continued Figure Page 45. Log DC conductivity versus reciprocal temperature for the 26.4 mole % lithia-silica glass 101 46. Log DC conductivity versus reciprocal temperature for the 33 mole % sodasilica glass 102 47. Log DC conductivity versus reciprocal temperature for the 25 mole % sodasilica glass 104 48. Thermal expansion curve for the 26.4 mole % lithia-silica glass as cast 105 49. Phase diagram for the lithia-silica system 116 50. Free energy-composition diagram for the lithia-silica system at 500°C 117 51. Summary of results for the 30 mole % lithia-silica glasses 122 52. Thermal treatment time required to develop loss peak versus composition for lithia-silica glasses 123 53. Phase diagram for the soda-silica system 125 54. Free energy-composition diagram for the soda-silica system at 550°C 126 X

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NOMENCLATURE Subscripts 1 and 2 refer to the parameters for the matrix and dispersed phase respectively. A Shape parameter defined by equation (1-11) d Interplanar spacing E Electric field F Frequency Q Quench rate R Resistance T Absolute temperature t Sample thickness ' + ' ;_ Transferrence number of positive and negative ion respectively V Volume fraction of dispersed phase V X Distance parameter defined on Figure 12 tan 6 Tangent of the loss angle 6 £ ' Real part of complex dielectric constant e ' ' Imaginary part of complex dielectric constant e. Static dielectric constant s e Dielectric constant extrapolated to high oo frequency p Surface charge density a DC conductivity XI

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NOMENCLATURE — Continued T Period of oscillation (() Electric potential 0) Angular frequency (27t/) tan 6 , , Tangent of the loss angle tan 6 _ AC component of the loss angle Xll

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Abstract of Dissertation Presented to the Graduate Council in Partial Fulfillment of the Requirements for the Degree of Doctor of Philosophy THE EARLY STAGES OF CRYSTALLIZATION IN ALKALI-SILICATE GLASSES By Donald L. Kinser December, 1968 Chairman: L. L. Hench Major Department: Metallurgy and Materials Engineering The initial stages of crystallization in glassy systems are of importance because of their influence upon the structure of the fully crystallized material. The precise nature of the structural changes occurring in the early stages of crystallization was the subject of this investigation. This investigation has shov/n the appearance of a metastable transition phase prior to the appearance of the equilibrium precipitate in four of the five glasses examined. The metastable phase enables the crystallization behavior of these glasses to be explained as a process involving the metastable precipitate as a nucleation site for the equilibrium lithium disilicate phase. Xlll

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The compositions examined include a 26.4, 3 and 33 mole % lithia-silica glass and a 25 and 33 mole % sodasilica glass. In order to detect the structural changes occurring during nucleation and crystal growth it was necessary to apply new x-ray, transmission electron microscopy, AC and DC electrical measurement techniques to the glasses. XIV

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CHAPTER I INTRODUCTION Crystallization of glasses has been the subject of scientific interest since the early work of Reaumur (1739) . He crystallized a glass bottle by imbedding it in sand and heating it to elevated temperatures for an extended time and found that the resulting crystalline article was extremely brittle and technically useless. In marked contrast to this experiment the recent work of Stookey (1962) led to the development of a class of crystallized glass-ceramic materials (Pyrocerams) which are quite strong (20,000 psi — 50,000 psi) and as a result are technically useful materials. The principal difference between the two products which gives rise to the striking contrast in mechanical properties is the extremely small crystallite size of the order of 1 micron in the Pyroceram materials as opposed to large crystals of the order of 1 millimeter (mm) for the Reaumur product. It is thus evident that the understanding of the physical processes which lead to an extremely fine grain size are of considerable importance both from a technical and a scientific point of view. 1

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2 In the case of crystallization in a glass matrix it has been generally accepted that the number of crystal nuclei present in the glass determine the ultimate lower limit to the grain size in the final crystalline product. Extreme importance is thus necessarily attached to the question of identity of the nuclei as v/ell as the behavior of these nuclei under varying conditions. The objective of the present work is to determine the nature and behavior of the structural changes occurring during the nucleation of simple binary alkali silicate glasses which in combination with other components form the basis of the Pyroceram crystallized glass-ceramics. In practice, the nucleation step in the preparation of a glass-ceramic is carried out at a lower temperature (nucleation temperature) than is the crystallization or growth step. Experimental observation of the structural changes on the size scale necessary to resolve nuclei in the early stages of crystallization requires an experimental technique capable of detecting and characterizing small crystals in the glass matrix. This problem is in some ways analagous to problems encountered in studies of the precipitation stages in age hardening aluminum alloys. One of the first experimental techniques employed in aluminum alloys was DC conductivity studies which gave a great deal of information about these systems. These

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3 systems have subsequently been investigated by electron microscopy, x-ray small angle scattering, and several other techniques. For these reasons and others which will become evident, electrical properties, electron microscopy and x-ray diffraction have been used to study the precipitation or nucleation stage in the crystallization of glasses. In order to understand and interpret the results of the experiments described later it is necessary to understand the conduction mechanisms operating in these glasses as well as the parameters which affect them. The historical review which follows is an attempt to place the present work in proper perspective with previous investigations. DC ConduGtion The DC behavior of glasses has been the subject of study for over 200 years. Early workers in the area include such eminent scientists as Franklin, Kohlrausch (1847), Hopkinson (1876), the Curies and Maxwell (1891). Other lesser known but important workers include Warburg (1884) , Tegetmeier (1890) , Fousserau (1883) and others . Warburg's early work established the presently accepted fact that electrical conduction takes place by the motion of the alkali metal ions in alkali-silicate

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4 glasses. This was established by passing a direct current through a test tube of "Thuringian" (CaO-MgO-Na20-K20Al203— Si02) glass which was filled with either mercury or sodium amalgam which served as electrodes. Pure mercury electrodes gave rise to polarization which caused the current flowing to drop to less than one-thousandth of its initial value in one hour. In the case of the sodium amalgam electrodes the current was observed to be constant over a protracted period of time; thus it was inferred that the sodium from the amalgam was the charge carrier in the glass. Le Blanc and Kerschbaum (1910) repeated Warburg's work and concluded, perhaps prematurely, that conduction is due entirely to the motion of sodium ions present in the glass. Kraus and Darby (1922) recognized that other conduction mechanisms could also be operating, and confirmed Faraday's law for charge transport in the glass by a relatively simple mass transport experiment. The probable error in the values of the transport number obtained is too large to allow one to assert conclusively that the conductivity was entirely ionic. More recently Kirby (1950) has reported that the transference number of the sodium ion is greater than 0.995. Russian workers have also been active in this area as reported by Mazurin (1965) but no values of the transference number are reported.

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5 Hughes and Isard (1968) have recently made transference measurements on ternary and higher order systems and reported values very near unity. From time to time various workers (Poole 1921, Cohen 1957 and others) have reported work with many glasses at high electric field strengths as well as work in glasses containing transition metal oxides. They have concluded that electronic conductivity is in fact operative under either of the two conditions above. The fact remains, and is almost universally accepted (Owen 1963, Morey 1954, and Mazurin 1965), that in alkalisilicate glasses alkali ion motion is responsible for electrical conduction except perhaps at high field strengths. Thermal History Effects Thermal history effects were first recognized by Fousserau (1883) in the observation of a marked time dependance of electrical conductivity during thermal treatments near the annealing range. Extensions of that work carried out by Fulda (1928) , Mulligan, Ferguson and Rebbeck (1925) , Littleton and Morey (1933) , Littleton and Wetmore (1936) , Foex (1944) , Joyner and Bell (1953) and others have shown that electrical conductivities in glasses are extremely sensitive to thermal history. This

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6 strong dependance of conductivity upon thermal history is generally manifested as a decrease in the conductivity with increasing thermal treatment. Concurrent with the conductivity changes, densification is observed. This observation gives one explanation of the conductivity behavior. Densification of the glass structure leaves less void space in the glass so that motion of the alkali ions is hindered by steric constraints so that conductivity is reduced during thermal treatments prior to the appearance of the equilibrium crystalline phases. Composition Effects Compositional effects in binary alkali-silicate glasses are generally well documented and not difficult to rationalize. Consider first the case of increasing alkali content in a binary glass. It is evident from the conduction mechanism that an increased concentration of the alkali or current carrying species should increase the conductivity. This behavior has been observed in the lithia, soda, potash and cesium glasses by Mazurin (1965) and Owen (1963) . Exam.ination of the molar equivalent composition glasses in the series lithium, sodium, potassium and cesium shows that the conductivity decreases in the series in an inverse relation to the ionic radii. Once again

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7 this is rather easily rationalized on the basis of steric considerations. EZeotvode Effects Ionic conduction gives rise to the phenomenon of electrode polarization or time-dependent conductivity. This has been recognized by relatively few investigators since the early v/ork of Warburg. Most investigations are conducted under conditions which the investigators apparently assume are non-polarizing. Proctor and Sutton (195 9, 19 60) have examined the polarization in an alkali-lead-silicate glass by applying a DC potential and measuring the potential distribution as a function of time. Their results agree qualitatively with predicted behavior which assumes that only one ion is mobile and it is blocked at the electrodes, Conditions which should lead to the smallest amount of polarization are (1) small field strength, (2) short time of measurement or AC extrapolations and (3) non-polarizing electrodes. The conditions above give rise to serious problems. The first solution, i.e., that of employing small field strengths, introduces the problem of measuring extremely small currents which in turn introduces obvious experimental difficulties. Short time or AC measurements

PAGE 23

require the definition of a time at which the electrode charge build-up is negligible and thus introduces an arbitrary variable. Non-polarizing electrodes are difficult to prepare because of the requirement that the activity of the charge carrying species in the glass be equal to that of the corresponding species in the electrode. This problem is soluble but it is not the most difficult problem because the "electrode" with the correct alkali ion activity must then be connected to some electronic conductor to allow measurements to be performed. Connection of the electrode to the measuring circuit then introduces problems similar to the original one. Electrode polarization in DC measurements has been avoided in almost all the reported work by allowing only very small currents to flow in the glass. This has given rise to considerable experimental difficulty in measuring the small currents. Some investigators have used AC measurements of conductivities to circumvent the problem. AC Behavior* AC electrical properties in the audio frequency range have been the subject of extensive investigations for only about 40 years. McDowell and Begeman (192 9) v/ere the first to conduct comprehensive studies of AC properties of glasses. They examined six glass

PAGE 24

9 compositions including lead glasses, borosilicates and lead borosilicates over a frequency range of 800 Hz to 1500 IC-iz and concluded that the dielectric losses could be satisf accorily rationalized on the basis of the behavior of individual "molecules and ions" rather than on the basis of the heterogeneous theories of Maxwell (1892) and Wagner (1914) . Strutt (1931) examined five commercial glasses including a borosilicate , a soda-lime-silicate and a "heavy" lead glass and observed a strong correlation between resistivity and dielectric loss which has been noted by a great many subsequent workers. Strutt also formulated an empirical equation (1-1) to describe the behavior of the dielectric loss. tan 6 = A exp aT (1-1) The constants A and alpha in the equation depend upon the glass composition as well as the measuring frequency. Other workers, notably Stevels (1946, 1950) and Moore and DeSilva (1952) , have largely corroborated the validity of Strutt' s equation (1-1). Robinson (1932) examined several glasses using "non-polarizing electrodes" and concluded that the observed behavior could be rationalized equally well by either the Debye dipolar theory or the Maxwell-Wagners heterogeneous dielectric theory.

PAGE 25

-loin addition to the experimental work cited above, a great deal of theoretical work was done in the period since the observation of dielectric absorption by Hopkinson (1876). Several books (Debye 1929, Smyth 1955 and Frohlich 1958) and innumerable papers treating the ionic theory of dielectric loss in solids, liquids and gasses have appeared so that discussion of these theories is unnecessary. Most of the theories developed give equations which are reducible in form to the equations commonly called the Debye equations. e e ' = e^ + -^ — (1-2) " 1 + W^T^ 1 + W^T^ (1-3) tan (S = ^ = — ^ (1-4) e + £ CO^T^ Concurrently with the development of the ionic or atomistic models, several workers commencing with Maxwell (1891) have developed macroscopic or heterogeneous theories to express the behavior of dielectric losses. Maxwell's initial work developed the frequency dependance of the dielectric loss for a stratified dielectric model. This was extended by Wagner (1914) to a uniform dispersion

PAGE 26

11 of spheres and subsequently by Sillars (1937) to spheroids, A model for ellipsoids was developed by Fricke (1953) . The recent article by van Beek (1967) includes a summary of equations characterizing the above systems and a great many other types of dispersions. The equations below from van Beek characterize a system of dispersed spheres of conductivity Oz and dielectric constant £2 with volume fraction V dispersed in a matrix of conductivity Oi and dielectric constant ei. 2a 1 + 2 + 2V (a 2 oi) S ^^ 2ai + Oz -TT—/ r + 3V ai V (oz Oi) V [2oi + 02) (£2 £1) (2£i + £2) (a2 oi] \2oi + Oz Vy (02 ai) |2 (1-5) 2£i + £2 + 2Vy (£2 £1) 2£i + £2 + V (£2 £1} (1-6) The relaxation time (x) of the system above is expressed as: 2£i + £2-V [£2 £1) T = £0 2ai + 02-V (02 oi) (1-7) The analagous equations for dispersed spheroids derived by Sillars are considerably more cumbersome than the sphere model equations. They are given below.

PAGE 27

12 Ol + e = Ci A fl V ) + V [oz Ol] a, + A^ [1 V J (02 Qi] + V 02 V a " V 1 + A^(a2 Ol) [^2 ei]-iei + A^(e2-ei) [02-O1] i^i + A^(l Vj[a2 Qi] El + £ =^ £ 1 00 ' A fl-V ]+V fe2-ei] ^' "" ^a^^'^uH^^"^^) (1-8) (1-9) T = £0 Ol + A^(l Vy) (02 Ol) (1-10) For prolate spheroids {a > b) a b ii)' 1 ii)' 1 ^n{(j. \m li^/^ 3/2 "'^i ' ^& (1-lla) For oblate spheroids (a < b) (f) a 1(f)' Ii {lYV'' arc cos (7-) (1-llb) For spheres (a = b) a 3 (1-llc) Because of the complexity of the above equations the

PAGE 28

13 conditions of small volume fraction and a i >> 02 are generally imposed in order to simplify the equations to the form below. Eo £1 e^ ei,{l + V, El + \[^2

PAGE 29

14 borosilicate glass. He concluded that the dielectric relaxation observed was a result of the motion of alkali ions in the random glass structure. His analysis further indicated that the distribution of relaxation times was very similar for all the glasses examined. Prod'homme essentially verified Taylor's results although there was some disagreement as to the breadth of the distribution of relaxation times. Barton (1965) has analyzed his results in a manner similar to the workers above and has been able to obtain qualitative results upon the depth of the ion potential wells which indicate that the potential well depths depend upon the "occupancy time." To the knowledge of the author, none of the v/orkers v;ho attribute dielectric losses to ion motion have examined the effect of thermal treatments or thermal history on either the dielectric losses or the distribution of relaxation times. The work of Isard (1962) , Owen (1961) and Charles (1963) is of importance because they analyzed their results on the basis of a heterogeneous dielectric theory. Owen concluded from analysis of the dielectric loss behavior of CaO-B203-Al203 glasses that dielectric losses occurred by the Maxwell-Wagner heterogeneous mechanism. He further concluded that the borate ri:,h phase was distributed as a dispersed phase in an alumina rich matrix.

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15 Observations were made on glasses thermally treated in the annealing-transformation range but no interpretation was given as to the morphological changes occurring. Isard's work with a number of glasses has led to the conclusion that the "classical theory of inhomogenity can satisfactorily explain the main loss peak in glass." However, he goes on to say that the high frequency {> IMHz ) behavior requires an explanation based on the atomistic theory. Charles (1963) has examined the dielectric behavior of a series of lithia-silica glasses with different thermal histories and reached several conclusions as to the effect of thermal treatments upon the morphology of the glass. His results for different quenching treatments were markedly different because of the sensitivity of the metastable phase separation to quenching rates. Analysis of his results gave information on the morphological differences and the connectivity of the various phases which v;as corroborated with replica electron microscopy. Glass Structure The current conception of the structure of glasses stems largely from the x-ray diffraction studies of Zachariasen (1932) and Warren (1937). Their work

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16 indicated that in silicate glasses the average siliconoxygen distance is 1.62 A and the average number of oxygens adjacent to a silicon atom is 4. These two observations form the basis of a structural model of tetrahedral SiOi* groups interlocking by sharing of oxygen ions between adjacent tetrahedra. This model, or models which closely resemble it, are generally accepted as representative of most silicate glasses. A great deal of literature exists in this area and is well reviewed by Mackenzie (1960) . Russian workers in the glass structure area were first to propose a theory of a microhetrogeneous glass structure. The early Russian work exemplified by the work of Lebedev (1940) , concludes that the structure of glass is one of microcrystallites with dimensions in the same size range as the interatomic distances in crystals. Glass structure on a larger scale (20 to 1,000 A) is an area in which electron microscopy has led to a detailed structure characterization. Electron microscopy has shown evidence of liquid-liquid phase separation in a variety of systems. The lithia-silica and the soda-silica binaries are two systems which have been shown to exhibit this type of behavior. The pioneering work of Slayter (1952) and Prebus and Michener (1954) has shown that silicate glasses contain

PAGE 32

17 structural heterogenities in the size range of 20 to 200 A. Subsequent investigations by Seward et al . (1967) and Shav; and Uhlmann (19 68) have shown that structural features in the above size range are a general feature of many glasses. The work by Vogel and Byhan (1964) in lithiasilica glasses has shown that most compositions in the Si02-Li20' 2Si02 phase field show structural heterogenities v/hose existence and general behavior can be rationalized by a metastable liquid-liquid separation. The existence of the metastable miscibility gap can be inferred from the "S" shaped liquidus on the phase diagram determined by Kracek (1939) . Tran's (1965) v/ork in soda-silica glasses has shown a sim.ilar phase separation in glasses between 9 and 20 mole % soda. No evidence of separation was observed in glasses of soda content greater than 19 mole %. Subsequent work in both binary systems involving thermal treatments in the annealing-transformation range has shown the coarsening of the liquid-liquid separation prior to the appearance of crystalline phases (Aver'yanov and Koshits 1966) . The mechanism of liquid-liquid separation has been investigated by various authors who have generally agreed upon the nature of the separation. Depending upon

PAGE 33

IS the corr.position of the glass, separation takes place either by nucleation and growth or by spinodal decomposition. Cahn (1968) has recently summarized the thermodynamic arguments for spinodal decomposition as well as discussed the various systems, both oxide and metal, in which the mechanism has been reported. Cahn and Charles (1955) have summarized the theory of phase separation and applied their results to various glass systems. Haller (1965), McCurrie and Douglas (1967) and others have examined many of the systems discussed in the above works and concluded that the observed structures could be explained by a random nucleation and growth process.

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CHAPTER II EXPERIMENTAL PROCEDURE Glass Preparation The objective of this work was to examine the effect of thermal treatment upon the electrical properties of alkali-silicate glasses and to correlate those properties with structural changes occurring in the --""^ glasses. To accomplish this objective it is necessary that the initial "state" of the glasses be well defined or at least invariant insofar as the parameters used to characterize the glasses could distinguish. With this in mind an experiment to examine the possible glass preparation variables such as melting time, tem.perature of melting, pouring temperature, alkali vaporization, quenching rate, annealing time and annealing temperature was designed. The above variables were examined by pouring a series of 30 mole % Li20-Si02 glass samples as shown schematically in Figure 1. The diagram depicts the process of melting a glass for 24 hours at 1350°C, pouring several samples at tv/o quenching rates followed by annealing of these samples for periods of time from 1 to 12 hours. Next the melt temperature was increased to 1450°C and held for 24 hours and a similar set of 19 -

PAGE 35

20 e 3 s «> o. E U50*C (Pouring) 1350'C steel \\ graphite mold*! \mo!d 300*C 1350* C Q,\\Q. OAQ (Annealing) m TTl TTT 24 Time (hours) 48 72 Figure 1. — Schematic melting procedure for 30 mole % lithia-silica glass.

PAGE 36

21 samples was poured. Following this the melt temperature was reduced to 1350°C, held for 24 hours, and a similar set of samples poured. Samples were selected from the above melt and electrical property measurements were carried out by methods discussed later. Each of the possible types of samples exhibited AC and DC electrical properties which were identical within the limits of experimental error with all others. From this result it is concluded that none of the above-mentioned variables, over the range examined, affects the initial electrical properties of the 30 m.ole % lithia-silica glass. The possibility that the initial state of the glass is different under the above-described conditions, but is not detected by the techniques employed, was examined by the following technique. Samples from each class of specimens described above were thermally treated and their electrical properties were re-examined. Once again the electrical properties of the various types of samples v/ere identical even though the treatment had changed the properties of the entire group. The property changes of the group are discussed later. It is thus concluded that variations of the variables set forth above within the range examined do not affect the initial structure of the glass.

PAGE 37

22 X-ray fluorescence examination of the samples melted for a total of 72 hours in a platinum crucible showed no evidence of platinum. In consideration of the above conclusions, preparation of glasses for the remainder of this study v;ere made in the following manner. ' 1. Glass batches from materials of puriry shown in Table 1 and total weight of approximately 1 kilograrri were weighed to an accuracy of 0.1 gram, giving an expected composition accuracy of at least 1 part in 10^. 2. The batch was then mixed in a jar mill without balls for a minimum of 1 hour. 3. A 150 milliliter platinum crucible was then filled from the batch and placed in an electrically heated silicon carbide element furnace at i350°C ± 5°C or 1450°C ± 5°C, depending on the glass being melted. 4. After approxim.ately 15 minutes the crucible was removed and refilled as necessary until the crucible was full. 5. The crucible v/as then covered with a platinum lid and the melt held at temperature for a minimum of 24 hours. 6. After 24 hours the lid was removed from the crucible and the crucible was replaced in the furnace to allow it to come back to 1350°C prior to glass pouring. 7. The glass v/as then poured in a 17.5 mm diameter steel m.old. A tightly fitted plunger was pressed into the molten glass, giving a sample thickness between 3 mm and 8 rrim and a diameter of 17.5 mm.

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23 TABLE 1 ANALYSIS OF GLASS RAW MATERIALS Comoound Weight % Sodium Carbonate ^ NaoCOs Chloride Nitrogen Phosphate Sulfate Arsenic Calcium and Magnesium Iron Potassium Silica Heavy Metals Lithium Carbonate ^ LizCos Na2Co3 Iron Sulfate Chloride Calci^um Phosphate Silicon Dioxide ^ Si02 Iron Alumina Titania Calcium and Magnesium 99.8 .0005 .0005 .0005 .001 .0001 .005 .0002 .001 .005 .0002 99.3 0.2 0003 0.3 0003 0003 ,0001 99.91 .019 .08 .009 Trace 'Baker Reagent. ^Foote Mineral Company. ^Pennsylvania Glass Sand Company,

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24 8. The resulting glass button was quickly removed (after approximately 30 seconds for cooling) and placed in a furnace at 300°C ± 5°C where it was held for 1 hour and air cooled. 9. After each sample was poured, the remaining glass was placed in the furnace and allowed to come back to the melting temperature, while the steel mold was chilled in tap water to keep it at the same temperature for each sample. It was found that 1 hour at 300°C was the minimum treatment which allowed the samples to be cooled to room temperature without breakage. The residual stresses remaining after this treatment were measured by standard birefringence techniques and found to be 3,000 to 5,000 psi, Electrodes For electrical measurements, the faces of the samples were ground parallel to within 0.05 mm with silicon carbide metallographic paper and polished with 600 grit metallographic paper. This surface was then cleaned with distilled water and quickly dried. In order to vapor deposit the desired double guard ring gold electrodes on the faces of the samples, paper masks were affixed to the sample faces. These paper masks with an inner diameter of 12.7 mm and an outer diameter of 15.9 mm were cut from a gummed paper label with a modified machinist's compass. The resulting masks were moistened

PAGE 40

25 with distilled water and carefully affixed to the samples so that they were concentric with the sample itself. The samples were then placed in a specially designed holder in a standard vacuum metallizer and the vacuum system was pumped down to less than 1 micron. The samples were then plated with gold from a tungsten basket placed about 100 mm from, the sample. Total electrode thickness was approximately 2 • 10"^ mm as determined from weighing samples before and after plating. The samples v/ere then removed from the evaporator and placed in a furnace at 300°C ± 5°C for 1 hour to increase the adherence of the gold to the sample. Samples were removed from the furnace, air cooled and the charred paper masks were carefully removed, leaving the sample v;ith gold electrodes in a double guard ring configuration. X-Ray Samples for examination in the Guinier (1956) and DeWolff (1947) x-ray camera were selected from the samples poured for electrical measurements in order to insure that their thermal history was identical to that for the electrical samples. These samples were heat treated, broken with a hammer and anvil, then ground in an alumina mortar and pestle to pass a Tyler 200 mesh screen. The ground glass-crystals were stored in a closed container to prevent moisture pickup and resulting hydration.

PAGE 41

26 The Nonius Guinier-DeWolf f x-ray camera and copper x-ray tube used for these experiments are shown schematically in Figure 2, The Guinier-DeVJolf f camera, a vacuum path, focusing quartz crystal monochromated powder camera, is capable of examining 4 samples simultaneously. The samples were held in 4 slots in a flat sample holder approximately 0.025 mm thick with Scotch #810 tape. The powder pattern v;as recorded on Kodak Type NS double emulsion x-ray film, developed 8 minutes at 23°C in Kodak Type D-76 developer and fixed 3 minutes at 23°C in Kodak "Rapid Fix." It was found that this technique gave the highest line intensity with a tolerable fog level for a given exposure. Using the above techniques, it was found that 0.1 weight % lithium metasilicate crystal in a prepared lithium disilicate glass standard could be detected. Detection of crystals with this small weight fraction required a 50 hour exposure at 40 kilovolts and 20 milliamperes. The resulting patterns were analyzed using a reader which allowed the crystal lattice spacings to be read directly from the film. Thermal Expansion and Softening Points Samples for thermal expansion measurements were prepared by melting the glass in the same manner as for the electrical samples. The dilatometric samples were

PAGE 42

27 IWonochromator X-ray Source Sample \ Film Figure 2 . --Schematic diagram of x-ray source and Guinier-DeWolf f camera.

PAGE 43

28 poured in a steel mold which tapered from a diameter of 12.7 mm to 15.9 mm in its five centimeter length. This taper was necessary to allow rapid removal of the sample from the mold to prevent breakage. Following pouring, the samples were held in a furnace at 3 00°C for 1 hour to prevent breakage, then removed and air cooled. Samples prepared in this m.anner were too highly strained to be cut with a cutoff wheel so they were all heat treated 1/2 hour at their subsequent heat treatment temperature to further remove strains. After this treatment the samples v/ere cut to slightly over 50.8 mm length with a water cooled silicon carbide wheel and the ends were polished to 50.80 mm ± .02 rccm. on 180 silicon carbide metallographic paper. The thermal expansions of the various thermally treated glasses were measured in an Orton recording dilatometer which is a quartz tube and push rod apparatus with a linear variable differential transformer transducer for measuring the expansion. Samples were separated from the quartz tube and push rod with 0.025 mm platinum foil to prevent reaction with the quartz. The apparatus was set up to plot the thermal expansion curve as a continuous function of the temperature.

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29 Electron Microscopy Replica Preparation Samples to be examined by electron microscopy were prepared in the same manner as the electrical samples and heat treated in the bulk state. After thermal treatments, the samples were fractured to expose a fresh surface from the interior of the sample. The fracture surface thus obtained was etched 1 minute in an aqueous 5 volume % hydrofluoric acid solution. The samples were then placed in an evaporator and a platinum preshadow was applied, followed by a carbon film. The carbon film replica was removed from the glass by immersing the sample in an aqueous 2% hydrofluoric acid solution and allowing the replica to float off the sample. This replica was placed on a copper microscope grid and washed several times in distilled water to remove the acid. Replicas were examined in a Phillips EM 200 electron microscope using standard techniques. Transmission Preparation Samples for examination by transmission electron microscopy were poured and heat treated in the same manner as the electrical samples. Samples were mechanically thinned to approximately 0.2 mm prior to chemical thinning, The mechanical thinning was accomplished by grinding and

PAGE 45

30 polishing one face of the poured button, cementing the polished face to a small flat piece of steel and grinding to the final thickness. Samples of 3 mm diameter were then cut from the thinned material while it was still fixed on the steel block. The samples were cut by using a hollow copper drill (3 mm ID) with a slurry of 400 grit silicon carbide in a drill press. The resulting sample blanks were removed from the steel block and the cement was removed using ethylene dichloride in an ultrasonic cleaner. The chemical thinning was carried out by dimpling the sample in the center, followed by a final thinning operation. The dimpling operation v/as accomplished by masking the outer edges of the samples with a lacquer ("Microstop" ) which did not allow the edges of the sample to be attacked. The solution used for this operation was made up of 10 parts hydrofluoric acid (48%) , 5 parts nitric acid and 14 parts acetic acid by volume. After the dimpling operation, the masking material was removed in acetone and the sample was carefully washed. At this point the sample has a relatively thick edge with a thin interior region which greatly facilitated handling. Final thinning was accomplished by alternately dipping the sample in hydrofluoric acid for short periods of time (30 to 60 seconds) and examining the center portion for holes,

PAGE 46

31 When the first visible hole developed, etching was stopped and the sample was washed several times in distilled water. The sample was then placed in an evaporator and a thin film of carbon was evaporated on one surface to prevent charging of the sample by the electron beam. The sample was then placed in the heating stage holder (Phillips PW 6560) with a large platinum aperture to facilitate heat conduction. The sample holder was then placed in the Phillips EM-200 electron microscope with a rotating-tilting stage and heated to the required observation temperature. The remainder of the electron microscopy was then conducted by standard techniques. Electrioat Measurements Sam-pie Chamber DC and AC properties were measured over a range of temperatures in a vacuum environment. The requirements of electrical shielding in AC measurements and guarding in DC measurements was of major importance in the sample holder design. The measurements require that the leads and contacts be made in a coaxial configuration. The sample holder design shown in Figure 3 used in these experiments utilizes a coaxial arrangement of leads as far as possible. Where this is not possible, high quality

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32 Figure 3. — Photograph of sample holder employed in electrical measurements.

PAGE 48

33 insulation with resistivity greater than 1 • 10^"* ohm • cm at room temperature was used. Electrical contacts to the samples and leads out of the sample chamber were platinum to avoid oxidation and thermoelectric problems. The sample chamber was vacuum sealed to allow all electrical measurements to be conducted in a vacuum of less than 1 micron. The sample chamber and sample were heated by means of an external nichrome heating element in conjunction with a Variac for temperature control. The sample temperature, as measured with a chrome 1-alumel thermocouple placed approximately 1 mm from the sample, was constant over a series of electrical measurements to ± 1°C of the set point AC Measurements Discussion of the electrical equipment is subdivided into the audio frequency (AF) equipment, the radio frequency (RF) equipment and the null detection system. The AF equipment, shown schematically in Figure 4, consisted of a Wayne-Kerr B-221 transformer radio arm bridge in conjunction with a Hewlett-Packard 651-A oscillator. A schematic diagram of the bridge circuit is shown in Figure 5 where Z and Z are the standard and unknown impedances respectively. The balance condition is satisfied, as indicated by a null on the detector, when

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34 CUDIO FREQUENCY MEASUREMENT APPARATUS 20HZ TO 20KHZ HP-65IA OSCILLATOR WAYNEKERR B-221 A.F BRIDGE SAMPLE CHAMBER NULL SIGNAL Figure 4. — Schematic diagram of audio frequency measurement equipment.

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35 00 TURNS Figure 5. — Schematic diagram of transformer ratio arm bridge.

PAGE 51

35 equal currents flov/ in each half of the center tapped transformer (T2) . When this condition is satisfied, the po-cential on the primary will be zero and the right hand teriTiinals of the unknov/n and standard will be at neutral potential. The same voltage is applied to both unknown and standard and for equal currents to flow in each half of the primary of transformer T2; the real and imaginary parts of the unknown impedance must be equal to those of the standard. The instrument is designed to allow values of the resistive and capacitative component to be read directly from, the instrument dials. The RF measuring equipment, shown in Figure 5, consisted of a Wayne Kerr B-601 bridge and the same oscillator as was used with the AF equipment. The bridge design is, in theory, similar to the AF bridge except that the transformers and standards used are designed for use in the RF range. The null detection system, shown in Figure 7, consisted of a General Radio 1232-A null detector used in conjunction with a General Radio 1232-Pl RF mixer and a Wayne Kerr 0-22-D beat frequency oscillator. The 1232-A covers the frequency range 20 Hz to 100 KHz directly, so that the AF and RF null signals up to 100 KHz were detected directly. At frequencies greater than 100 KHz, it v/as necessary to use a beat frequency technique to

PAGE 52

37 RADIO FREQUENCY MEASUREMENT APR^RATlB 20KHZ TO lOMHZ KP-65IA OSCILLATOR

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38 DETECTOR SYSTEM Rf NUU. SlONAL GENERAL RADIO 1232PI R.F. MIXER WAYNE-KERR 022D LOCAL OSCILLATOR A F NULL SIGNAL GENERAL RADIO 1232-A NULL DETECTOR n Figure 7. — Schematic diagram of null detection system.

PAGE 54

39 reduce the signal frequency to the range of the 1232-A detector. This was accomplished by mixing a local oscillator signal of frequency 100 KHz greater than the measurement frequency with the signal to be detected. This gives a beat frequency signal proportional to the original signal at a frequency of 100 K?Iz which can be detected by the 1232-A detector. DC Measurements DC measurements were made with the equipment shown schematically in Figure 8. The short time measurements were conducted by displaying a signal proportional to the current flowing in the sample on a Hewlett-Packard 140 A oscilloscope with a 1420 A time base and 1402 A dual trace amplifier and photographing the oscilloscope trace with a Tektronix C-12 camera. This technique allowed the sample conductivity to be measured in a time of less than 5 milliseconds in most cases. The minimum time depended upon the range of the Keithly 416 high speed picoammeter and no measurements were made in times less than the response time of the picoammeter. The DC potential for the measurements was furnished by a Hewlett-Packard 6217 pov;er supply which has a voltage stability of less than 0.10% + 5 millivolts in 8 hours with less than 200 microvolt AC ripple. The connections

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40 D C MEASUREMENT APPARATUS HP-62I7A DC POWER SUPPLY lOVPC SAMPLE CHAMBER TEKTRONIX C-12 OSCILLOSCOPE CAMERA HP-I40A osauoscoPE TllltiCK •tPUT KEITHLEY 416 HIGH SPEED PICOAMMETER voLTAae IWUT HP-4I2-A I VACULIM TUBE VOLTMETERj (CURRENT MODE) | Figure 8. — Schematic diagram of DC measurement equipment.

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41 to the sample and picoainineter were made with guarded leads as indicated in Figure 8. Sample currents greater than 3 • 10"^ amps were measured on a Hewlett-Packard 412-A vacuum tube voltmeter and because of the large time constant of this instrument, no measurements could be made in times less than 1 second.

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CHAPTER III POLARIZATION The objective of the polarization experiments was to establish the behavior of the sample under a DC field in order that the electrode polarization could be eliminated as a cause of relaxation phenomena in the AC measurements. This objective requires that the time required to polarize the sample electrodes be greater than the maximum period (minimum frequency) used in the AC measurements. Polarization results from a 30 mole % lithia glass in the as cast state are presented in Figure 9. These results are a combination of DC and AC conductivity measurements. The "time" for the AC measurements was taken as the reciprocal of the measurement frequency. This definition of the time parameter results in a reasonably good fit in the region of overlap of the two curves and is thus considered to be a valid definition. The agreement of the two curves is in fact remarkably good considering the fact that in the overlapping region both measurement techniques are approaching their time limits and the accuracy is subject to larger errors in that time domain. 42 -

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43 7 ' E

PAGE 59

44 The general shape of the curves for the as cast glass, as expected, shows increasing conductivity in the short time (10""* to 10"^ second) range, a flat intermediate region (10~^ to 10"^ second) and a region of electrode polarization for times in the 10~^ to lO"*"^ second range. The curve corresponding to the sample heat treated 5 hours at 500°C exhibits more structure than the as cast glass. The short time and initial flat region is quite similar to the as cast behavior up to about 2 • 10"^ second. At that point a dispersion appears and at longer times the flat tries to reappear but is masked by the appearance of electrode polarization. Discussion of the dispersion in the heat treated sample will be presented in the AC properties section for reasons that will be evident later. The appearance of the polarization problem is illustrated by the series of conductivity-time plots in Figure 10. It is evident that at low temperatures, when the conductivity is low, the problem is unimportant but at high temperatures the problem is very serious. The effect of the polarization upon the measured DC conductivity is shown in a more conventional form, the log conductivity versus reciprocal temperature plot, for two arbitrarily chosen times of measurement (approximately 1 second and 60 seconds) as shown in Figure 11. The problem of

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45 10" 10E «io3Q JO 20 30 *Q 50 60 70 80 90 100 HO TIME (Seconds) Figure 10. — Conductivity versus log time curve for 33 mole % lithia-silica glass as cast.

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46 10 E i. c o I u U o o T = 1 second 4 T= 69 seconds 10 _L J \ J., =L J.2 2.4 2.6 2.8 30 3.2 3>« lo'/T(°C) 36 Figure 11. --Log DC conductivity versus reciprocal temperature for 33 mole % lithia-silica glass as cast.

PAGE 62

47 polarization is less important in the low temperature region but at higher temperatures the problem in its most serious state causes an inversion in the slope of the curve. It is thus evident that DC measurements made without consideration of the problem of polarization are virtually meaningless if they are measured over a temperature-time range where the phenomenon shown above is significant. In order to develop a working criterion for determining the effect of polarization on conductivity measurements, a model consisting of a plane parallel capacitor with mobile positive charges in the dielectric and totally blocking electrodes was chosen. This model is approximiately the situation in an alkali silicate glass capacitor. With a potential applied, this capacitor (Figure 12) v/ill have a surface charge equal to the negative of the volume charge assumied to be uniformly dispersed in the dielectric. It is assumed that total polarization occurs when the back voltage due to the surface charge is equal to the applied voltage and the apparent conductivity is zero. To allow comparison of the surface charge calculated from the model above with an experimentally derived value, it was necessary to define electrode polarization on the basis of the conductivity time curves. The criterion

PAGE 63

48 UNIFORM POSITIVE SURFACE CHARGE UNIFORM NEGATIVE VOLUME CHARGE CATHODE ANODE Figure 12 . --Schematic diagram of model used for polarization calculation.

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49 chosen was to assume that a conductivity drop of one full order of magnitude represented total polarization of the sample. The charge transported during the time required for polarization to occur will be calculated from the experimental results and compared with that predicted from the model. The electric field in the capacitor due to the uniform surface charge on one electrode is given by equation (3-1) from Page and Adams (1958) . E = .5-^— (3-1) The field due to the uniform negative charge of equal magnitude but opposite sign to the surface charge is given by equation (3-2) (Page and Adams 1958) . E = -^ ^ (3-2) EEO t In order to obtain the potential between the two plates, the two fields above are added and integrated over the electrode spacing. As a result of symmetry, the uniform volvime charge does not contribute to the potential thus the integration of (3-1) gives equation (3-3) . 4, = ,P^ • (3-3)

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50 Equation (3-3) gives the magnitude of the back potential corresponding to a given polarization charge on or adjacent to the cathodic electrode. It is thus possible to calculate from equation (3-3) the magnitude of the surface charge from the applied potential, sample thickness and the dielectric constant obtained from AC measurements. Substitution of the above values from the sample used to obtain the polarization curves in Figure 10 yields a surface charge of 9.80 • 10"** coulombs/cm^. In order to obtain an experimental value of the surface charge, it is assumed that all the charge transported is left on the electrodes. The charge transported is the time integral of the current over the time of polarization. Because the analytical form of the current tim.e behavior is not known, the integral above is approximated by assuming that the current drops from its initial value to zero in the time interval considered. This assumption can be justified by reference to Figure 10. The conductivity-time and current-time curves have the same shape because of Ohm's law; hence Figure 10 can be looked upon as a current-time plot. The value of the current transported during the polarization depicted in the 180°C curve of Figure 10 calculated in the manner described above yields a surface charge of 1.33 • 10"^ coulombs/cm^ .

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51 Comparing this value with the value calculated from the model above, it is evident that the agreement is good considering the approximations made. There are, at least, two possible reasons for the differences. The electrodes have been assumed to be totally blocking to the charge carrier but it is possible that this is not the case and part of the charge transported is not on the surface of the dielectric. Another possible source of the difference is the approximation of the shape of the current-time curve. The approximation of the curve shape in Figure 10 yields a value of the current which is too high. These two factors are sufficient to explain the observed difference and the experimental measurement. The results above indicate that the problem of polarization can be approached by two methods. The first and most direct consists of measuring the polarization behavior of each sample and analyzing the conductivities directly from the polarization results. The second and faster method is to calculate the conductivity required to cause polarization in the measurement time used. This would allov; prediction of cases in which polarization difficulties will arise. Clearly both techniques have some inherent difficulties but either one or both methods should be used to eliminate or reveal polarization problems in DC measurements.

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52 The DC measurements reported herein were made by measuring representative polarization curves and taking the conductivity values from the time independent portions of those curves.

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CHAPTER IV EXPERIMENTAL RESULTS X-Ray The objective of the x-ray studies was to detect and identify the crystals appearing in the initial stages of crystallization. The initial x-ray studies were conducted to determine the lower limits of detection of crystalline phases in a glassy matrix. Examination of a series of lithium disilicate glasses mixed with 5.0, 1.0, 0.5 and 0.1 weight % crystalline lithium metasilicate showed that the crystals in a 0.1 weight % standard could be observed. A similar experiment with crystalline lithium disilicate in lithium disilicate glass gave the same lower lim.it of detection. Results of the x-ray examination of the 30 mole % lithia-silica glass following various heat treatments at 500°C are presented in Table 2. The actual patterns are not presented because the lines of interest are very weak even after 50 hour x-ray exposures and photographic reproduction is difficult. The phases present change with thermal treatment at 500°C from the glassy material in the as cast form, to a glassy material plus a phase tentatively identified as crystalline lithium metasilicate 53 -

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54 TABLE 2 SUMMARY OF X-RAY LATTICE SPACINGS AND RELATIVE INTENSITIES FOR 3 MOLE % LITHIA GLASSES Standard Li20*2Si02 Standard Li20'Si02 30 mole % Li20 Heat Treated 5 hrs at 500°C 30 mole % Li20 Heat Treated 50 hrs at 500°C J (A) Intensity d (A) Intensity d (A) Intensity d (A) Intensity 7.4 5.45 1.70 1.18 J. 75 J. 65 J. 58 J. 30 1.95 >.90 2.70 >.39 >.35 1.21 2 10 KM.S.) 1 10 10 10 1 (M . S .) 1 1 (M.S.) 5 5 2 4.70 3.30 2.70 2.35 2.08 10 10 10 10 4.70 4.35 4.18 3.30 2.91 2.81 2.70 1 1 1 1 1 5.5 3.75 3.65 3.58 1 1 1

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55 with a weight fraction of approximately 0.1 weight %. This 5 hour x-ray pattern shows seven lines, three of which correspond to the stronger lithium metasilicate crystal lines. The other four lines are weaker and, at present, not conclusively identified. It appears that the unidentified lines correspond to a transition phase of lower s^Tnmetry than the orthorhombic metasilicate. Indexing is difficult, if not impossible, because of the small number of lines observed. Examination of the 10 and 20 hour samples show only a diffuse peak characteristic of the glassy state with no evidence of crystalline phases. The 50 hour sample exhibits the four strongest lines corresponding to the equilibrium lithium disilicate crystal. X-ray results for the 33 mole % lithia (lithium disilicate) glass are shown in Table 3. The reaction sequence observed in this glass is that of the glassy material transforming to the equilibrium lithium disilicate crystal with thermal treatment. The equilibrium lithium disilicate precipitate is observed after 50 hours at 500°C. X-ray examination of the 26.4 mole % lithia glass did not reveal the presence of any crystalline phases with thermal treatments up to 50 hours at 500°C.

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56 TABLE 3 SUMMARY OF X-RAY LATTICE SPACINGS 7VND RELATIVE INTENSITIES FOR 33 TIOLE % LITKIA-SILICA GLASSES Standard Crvstal i, 3 3 mole % LiaO Li20-2Si02 1 50 hrs at 500°C

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57 The results of x-ray examination of the 33 mole % soda (sodium disilicate) glass are presented in Table 4. These results show the appearance of the equilibrium sodium disilicate crystal after 10 hours at 550°C. Further treatments at this temperature only increase the amount of crystalline sodium disilicate present in the glass . The x-ray results of the 25 mole % soda glass do not show evidence of crystallization with treatments of 50 hours at 550°C. Eleotvon Miarosaopy The objective of the electron microscopy was to identify and determine the morphology of the phases appearing during the initial stages of crystallization. Replica Techniques Replica electron micrographs of the 30 mole % lithia glass are presented in Figures 13 to 17. The micrograph of the as cast glass (Figure 13) exhibits a droplike structure with separated regions in the size range 0.1 to 0.5 micron. It is known (Vogel and Byhan 1964) that silica rich glasses are attacked by hydrofluoric acid much less rapidly than lithia rich silicate glasses. From this information and a knowledge of the shadowing direction, it is evident that the drop regions in Figure 13 are richer in silicon than the surrounding matrix.

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58 TABLE 4 SUMMARY OF X-RAY LATTICE SPACINGS AND RELATIVE INTENSITIES FOR 33 MOLE % SODA-SILICA GLASSES Standard Crystal Na20'2Si02

PAGE 74

59 ..^Ba Figure 13. --Replica electron micrograph of 3 mole lithia-silica glass (30,000x) as cast.

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60 SSi^agife^ Figure 14. --Replica electron micrograph of 30 mole % lithia-silica glass (26,000x) after 5 hours at 500°C.

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61 Figure 15. — Replica electron micrograph of 30 mole % lithia-silica glass (31,000x) after 10 hours at 500°C.

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62 Figure 16. --Replica electron micrograph of 30 mole lithia-silica glass (31,000x) after 20 hours at 500°C.

PAGE 78

63 Figure 17 . ---Replica electron micrograph of 30 mole % lithia-silica glass (37,000x) after 50 hours at 500°C.

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54 The micrograph of the 30 mole % sample heat treated 5 hours at 500°C is shown in Figure 14. This micrograph exhibits the separated drop regions characteristic of the as cast sample and further shows a secondary drop-like separation in the matrix surrounding the primarydrops. This secondary separation in the matrix is the separation of the lithium metasilicate in the matrix. Figure 15 corresponds to the 30 mole % glass heat treated at 500°C for 10 hours. This micrograph shows some evidence of the' initial droplets but they are not as pronounced as in the earlier thermal treatments. The matrix exhibits a fine scale separated drop morphology similar to that in the 5 hour treatment. Thus, the drop regions are beginning to dissolve with thermal treatment. The micrograph of the 3 mole % glass heat treated 20 hours at 500°C is shown in Figure 16. This micrograph exhibits an overall uniformity of structure with only traces of the primary drop separation. The micrograph of a sample heat treated 50 hours at 500°C is shown in Figure 17. This micrograph shows remnants of the structure present at 5 hours but it is generally more homogeneous than all the other glasses examined in the above series. The raised regions are the equilibrium lithium disilicate precipitate.

PAGE 80

65 Transmission Electron Miarosaopy Diffraction Patterns The electron diffraction patterns taken during a thermal treatment at 4S0°C are presented in Figures 18 to 23. The patterns initially show no crystallinity followed by the appearance of lithium metasilicate and lithium disilicate crystals. The interplanar spacings obtained from the diffraction patterns and tabulated in Table 5 indicate that during the early stages of crystallization the crystalline lithium metasilicate predominates, but in the latter stages the lithium disilicate predominates. The lattice spacings taken from the ASTM card file (Smith 1967) are those corresponding to room temperature, while those discussed above correspond to 480°C. The results of Glaser (1967) with high temperature x-ray diffraction on the lithium disilicate indicate that the expected expansion in the {130} lattice spacing for example is less than one percent. This is within the expected experimental error in the electron diffraction measurements so the comparison with room temperature lattice spacings is a valid one. Transmission Electron Micrographs Figure 24 is a micrograph of the 30 mole % glass before thermal treatment. This micrograph exhibits a netv;ork structure which is undoubtedly a result of the

PAGE 81

66 Figure 18 . --Electron diffraction pattern of as cast 30 mole % lithia-silica glass at room temperature.

PAGE 82

67 Figure 19 . --Electron diffraction pattern of 30 mole lithia-silica glass at 4S0°C after 5 minutes at 480°C.

PAGE 83

68 Figure 20. — Electron diffraction pattern of 30 mole lithia-silica glass at 480°C after 14 minutes at 480°C.

PAGE 84

69 Figure 21. — Electron diffraction pattern of 30 mole % lithia-silica glass at 480°C after 26 minutes at 480°C.

PAGE 85

70 Figure 22. — Electron diffraction pattern of 30 mole % lithia-silica qlass at 480°C after 2.5 hours at 480°C, Figure 23. — Electron diffraction pattern of 30 mole % lithia-silica glass at 480°C after 3.75 hours at 480°C,

PAGE 86

71 LA < Eh U o ^— . O O CO CM &4 u o o CO •H c> 0) H O •H ^ O 03 4J CO iH fC -P X o rH CO rs • !h O -P CO H fd -p X H CO CO CO CO en CO ^ "^ XrH *r-* »r~( 1 4-) •H W 4-1 H > r\ S S £ > > > CO o o O") 1^ LO rLD on on CM CO en CO CO CO S a S S S CO > o CO > > Ln in > o CO 00 CO CM > en 5 CO > CO [2 g: S 12 & S > > > > -P •H W S -P c H 5 (N 00 o CM O CM 00 o CM
PAGE 87

72 K m < U o O CO .— . '^ on CN] <2J (U •H r^ ^— en U o . — . o oj CO CN in CN U o

PAGE 88

73 Figure 24 . --Transmission electron micrograph 30 mole % lithia-sdlica glass at room temperature as cast (41,000x) .

PAGE 89

74 carbon deposited on the glass surface. The carbon layer partially masks the glass structure but the diffraction pattern leaves no doubt that some of the structure present is crystalline metasilicate . The carbon surface film is discussed further in the following section. Figure 25 is a micrograph taken at 4S0°C after 11 minutes at 480°C. It shows the appearance of black spots which are apparently crystals in a glassy matrix. The diffraction pattern at this point in time shows that the crystals present are primarily lithium metasilicate. Bright and dark field micrographs taken after 35 minutes at 430°C are presented in Figure 26. The dark field micrograph was taken from the {170} diffraction ring of the lithium disilicate so that the bright crystals in the dark field image are lithium disilicate crystals. Micrographs taken after 2.3 hours at 480°C are shown in Figure 27. The dark field image corresponds to the {170} diffraction ring in the lithium disilicate pattern hence all the bright areas in the dark field micrograph are lithium disilicate crystals. The disilicate crystal size has increased in the time elapsed between Figure 26b and Figure 27b, indicating the growth of the lithium disilicate crystals in the glassy matrix. The last micrograph in this series at 480°C, shown in Figure 28, was taken after 4.8 hours. This micrograph shows the development of an elongated morphology from the previously equiaxed morphology

PAGE 90

75 Figure 25, --Transmission electron rr.icrograph of 30'inole % lithia-silica glass at 480°C after 11 minutes at 480°C (42,000x).

PAGE 91

76 rff-'-i; • 'rfj . L Figure 26. --Transmission electron micrograph of 30 mole % lithia-silica glass at 480°C after 35 minutes at 480°C. (a) Bright field (b) Dark field (46,000x).

PAGE 92

11 ?'^\^. ."•"' V .'jl-*-.'%? Figure 27. — Transmission electron micrograph of 30 mole % lithia-silica glass at 480°C after 2.3 hours at 480°C. (a) Bright field (b) Dark field (31,000x) .

PAGE 93

78 Figure 28. — Transmission electron micrograph of 30 mole % lithia-silica glass at 480°C after 4.8 hours at 480°C (23,000x).

PAGE 94

79 The effect of the carbon film on the surface of these glasses was investigated by evaporating a carbon film similar to that used on the glasses on a thin mica crystal and examining the evolution of the structure with thermal treatment. The carbon film on mica exhibited an initial structure similar to that shown in Figure 24. Keating of the mica-carbon film to 480°C and following the microstructural changes revealed a lower volume fraction of essentially the same black spots characteristic of the glass micrographs in Figures 25 and 26a. It is thus concluded that some of the black spots must be carbon. It was not possible to obtain a dark field image from the weak and diffuse lithium metasilicate rings during the sequence. This difficulty indicates that the observed black spots may be crystalline carbon but no carbon diffraction rings were observed. This leaves the identity of the black spots open to question but it appears that some of the black spots must be lithium metasilicate. It is possible that the carbon masks the lithium metasilicate crystals and hence they are not observed except in the diffraction pattern. Figure 29a and 29b are micrographs taken at 500°C after 1 hour at 500°C. The dark field image is taken from the area circled on the diffraction pattern in

PAGE 95

80 / y i ^P^v Figure 29. --Transmission electron micrograph of 30 mole % lithia-silica glass at 500°C after 1 hour at 500°C. (a) Bright field (b) Dark field (31,000x) .

PAGE 96

81 Figure 30. The important feature of these micrographs is the extremely long crystals growing in the glass matrix. The growth of the whiskers out of the glass matrix is also an interesting feature. A room temperature transmission micrograph of a glass heat treated in the bulk form for 5 hours at 500°C is shown in Figure 31. There is a pronounced similarity between this micrograph and the replica micrograph of a similar sample shown in Figure 14. AC Results The objective of the AC measurements was to monitor the initial stages of crystallization and characterize the structure and structural changes occurring during this period. 20 Mole % Lithia-Silioa Glasses The results of the AC measurements are presented in the form of the AC loss angle (tan 6^^) as a function of frequency and temperature. The tan 5 ^ values were calculated from equation (4-1) following the method of Charles (1963). f Total . p , 4.-,^ r _ 1 ^^^ ^otal = — C = ^^^ ^AC + ^^^ ^DC -^ 1 + 1 ^AC ^DC (4-1)

PAGE 97

82 Figure 30. — Electron diffraction pattern of 30 mole lithia-silica glass at SOO'^C after 1 hour at 500°C.

PAGE 98

83 Figure 31 .--Transmission electron micrograph of 30 mole % lithia-silica glass at room temperature after 5 hours at 500°C in bulk form (31,000x)

PAGE 99

84 The tangent of the loss angle is shown as a function of log frequency for the various heat treatments at 500°C in Figures 32-3 6. It can be noted from Figure 32 that the as cast glass is free of loss peaks, but after a 5 ± 1/2 hour heat treatment (Figure 33) at 500°C, large loss peaks have appeared. The tolerance on the time of appearance for the loss peaks v;as established by heat treating a sample in 1/2 hour increments and mieasuring the AC properties following each heat treatm.ent. Further heat treatments at the same temperature for times up to 50 hours (Figures 34, 35 and 36) cause the magnitude of tan 5^_, to decrease and the peak location to shift to higher frequencies. The tem.perature dependence of the frequency maxima of the tan 6^^ curves is shown in Figure 37. The frequency maxima curves have the Arrhenius form typical of a thermally activated process. The activation energy of the frequency maxim.a-reciprocal temperature curves is 14.7 kilocalories mole in all cases shown in Figure 37. The equivalence of the loss process activation energy and the DC conductivity activation energy indicates that the loss process is the result of ionic motion. The loss behavior results are summarized in Figure 38. 23 Mole % Lity.i aSilica Glasses Results for the 33 m.ole % lithia glass are summarized in Figure 39. The loss spectrum of this glass shows the appearance of the loss peak between 2 and 5 hours heat

PAGE 100

85 NO HEAT TKEATK: »»8 122 — o 92 — O 24 .-.o...-_ ta LOG„ FREQLENCY Figure 32. — Tan 6 versus logio frequency for the 30 mole % lithia-silica glass as cast.

PAGE 101

-sera u < z < t5 5 HOURS SOO'C. ise

PAGE 102

87 8 • lOHOUFS 500*C. It* -o 122 --0-C2 — 24 --01 4 LOGio FREQUENCY Figure 34. — Tan 5 versus logio frequency for the 30 mole % lithia-silica glass after 10 hours at 500°C.

PAGE 103

88 18Z < 20 HOURS 500'C. Its -o 122 — o 2 — o 14 .—o...... LOG,o FREQUENCY Figure 35. — Tan 5 versus logio frequency for the 30 mole % lithia-silica glass after 20 hours at 500°C.

PAGE 104

89 50 HOURS 500° C. S •o loGk, frequency Figure 36. — Tan 5 versus logio frequency for the 30 mole % lithia-silica glass after 50 hours at 500°C.

PAGE 105

90 50 HOURS ,-20 HOURS r-iO HOURS 2.4 ze 28 iO i2 3.4 Figure 37. — Log of frequency maxima versus reciprocal temperature for various thermal treatments at 500°C.

PAGE 106

91 AS CAST 30 mole % L12O LOGo FteQUENCY Figure 38. — Tan 5 versus logio frequency for the 30 mole % lithia-silica glass for various thermal treatments measured at 80°C.

PAGE 107

92 35 L.,0-2S.O^ Tf»otm«ot MeOSureTienf Terpp, eo "C 84 'C O OS OBT vy 5 nrs at SOCC O 10 ITJ ot 5O0«C — •— lOOhrsot 500*C 88 -C 79 'C LOG,„ FREQUENCY Figure 39. — Tan 6 versus logio frequency for the 33 mole % lithia-silica glass for various thermal treatments.

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93 treatment at 500°C followed by its growth between 5 and 10 hours. Between 10 and 50 hours the loss peak disappears and with further thermal treatment the loss background is depressed. 26.4 Mole % Lithia-Silioa Glasses AC loss behavior for the 26.4 mole % lithia glasses is presented in Figure 40. The loss spectra of this glass are essentially unchanged by thermal treatment up to 20 hours at 500°C but between 20 and 30 hours a loss peak appears. Further thermal treatment up to 50 hours decreases the magnitude of the loss peak. 2 3 Mole % Soda-Silica Glasses AC loss spectra of the 33 mole % soda glasses are presented in Figure 41. The as cast glass is free of losses but with 5 hours thermal treatment at 550°C a loss peak appears. Further thermal treatment increases the m.agnitude of the loss peak (i.e. 10 hour curve) and the loss peak eventually declines in magnitude and is not observed after 50 hours at 500°C. 25 Mole % Soda-Silioa Glasses The results of the AC measurements in 25 mole % soda glasses are summarized in Figure 42. The behavior of the loss spectra of this glass with thermal treatment is unlike the other glasses examined. The low frequency

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94 It cast 84'C i bourses 10 hou r s 84 JOhoort IS SOhours IJ SOhoars 13 J 4 } lOe,Q FREQUENCY (Hii Figure 40. — Tan 5 versus logio frequency for the 26.4 mole % lithia-silica glass for various thermal treatments.

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95 150 125 lOO_i < o c o I75 5025 /-NOjO2SiOj Treofmont D 03 cost a— 5 hrs. Of 550 °C G-10 hra m 550°C 50 hra ot SSO'C Meosurcment Temp. 84 "C 79 "C 8 2 "C 775 "C LOG„ FREQUENCY 'o Figure 41. — Tan 5 versus logio frequency for the 33 mole % soda-silica glass for various thermal treatments.

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96 CO \l

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97 rise to large values is observed in all the thermal treatments examined. The frequency of this low frequency rise changes in a systematic manner in the course of thermal treatments. DC Conductivity Results The objective of the DC measurements presented here was to analyze the morphological changes occurring during the early stages from the DC results. All the DC results reported here were taken in the time independent region of the conductivity-time curves prior to the appearance of electrode polarization. The activation energies for the conduction process shown on the respective conductivity plots were calculated assuming an Arrhenius type temperature dependence. Lithia-Silica Glasses The DC results for the 30 mole % lithia glass are presented in Figure 43 in the form of the logarithm of conductivity versus the reciprocal of absolute temperature. The data points corresponding to the 5 and 20 hour heat treatments were omitted because they are identical to the as cast and 10 hour samples respectively. The magnitude of the conductivity is decreasing while the activation energy is increasing. The magnitude of the activation energy

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98 IS) Kcji/M NO HEAT TREATICNT 2.0 22 2.4 ^6 2B Figure 43. — Log DC conductivity versus reciprocal temperature for the 3 mole % lithia-silica glass.

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99 change is near the anticipated experimental error, thus the change in the activation energy may not be real. Conductivity results for the 33 mole % lithia glass are shown in Figure 44. The DC conductivity of the 33 mole % lithia glass after 5, 10 and 50 hours at 500°C was measured but it was identical to the as cast glass and was omitted from Figure 44. The DC conductivity of the 33 mole % glass is unchanged by thermal treatments for times up to 50 hours at 500°C but the magnitude of the conductivity drops between 50 and 100 hours. The activation energy after the conductivity drop is unchanged within the limits of experimental error. The conductivity-temperature behavior of the 26.4 mole % lithia glass is presented in Figure 45. The behavior of the conductivity of this glass with thermal treatment is similar to that for the 3 and 33 mole % glasses except that the conductivity drop occurs between 20 and 5 hours. The activation energy after the conductivity drop is also unchanged in this glass. Soda-Siliaa Glasses The DC conductivity behavior of the 33 mole % soda glass is shown in Figure 46. No change in the conductivity behavior was observed in this glass with thermal treatments up to 50 hours at 550°C, hence no results other than the as cast glass are presented.

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100 UjO as'O, yks OBT 9^ onnoaivd 14 9 Kcal/M Figure 44. — Log DC conductivity versus reciprocal temperature for the 33 mole % lithia-silica glass.

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101 eA» cast XOO hour* St 500'C tr.i Kci!/mol« faiL. J. A .17 :!. ^. »!<"""" f -' u u u u u k'/t Figure 45. — Log DC conductivity versus reciprocal temperature for the 26.4 mole % lithia-silica glass.

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102 > 10 lO E « o :6 lO 1 21 Nq,0 2SiO, As ccEl and annealed 12 9 Kcal/M 23 25 27 1000 29 31 33 (°K ') Figure 46. --Log DC conductivity versus reciprocal temperature for the 33 mole % soda-silica glass.

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103 Conductivity behavior of the 25 mole % soda glass is shown in Figure 47. Data points are included for the as cast glass and the glass after heat treatment for 50 hours at 550°C but both sets of data are within expected experimental error. The conductivity behavior of the 25 mole % soda glass is thus unchanged during thermal treatments up to 50 hours at 550°C. Thermal Expansion The lower transformation point (LTP) or softening points were measured to assess the relative 'refractoriness' of the different glasses to allow comparison of the kinetics of the processes occurring in the initial stages of crystallization. The thermal expansion coefficients obtained as a by product of the LTP measurements will be compared to the values in the literature to establish the validity of the measurements. A thermal expansion curve typical of the alkalisilicate glasses investigated is shown in Figure 48. The line drawn through the curve serves to define the coefficient of linear expansion and the lower transition point. The lower transition point, as shown on the curve, is defined as the temperature at which a sudden increase in the expansion coefficient is observed (Morey 1954) .

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104 bAi catt *IOOhour« •! fSO'C ,» /— 1«.0 Kcal/mal* I I I I I I 1 ] 1 1 1 1 L_ U 2J U U 3^ 3-' 3'i >' 3-0 >-' 3-3 3.3 ].< IS U W'/T Figure 47. — Log DC conductivity versus reciprocal temperature for the 25 mole soda-silica glass.

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105 -

PAGE 121

106 Values of the linear coefficient of expansion and the LTP for the glasses exar:\ineG. are tablulated in Table 6. The esti~.ated error in the expansion coefficient measurement is 10%, thus all of the values for the lithia glasses are the same within experimental error. For comparison purposes, values for the glasses obtained from Morey (1954) and Leko (1967) are presented in the table. Values for the lithia and soda glasses are largely in agreement with those previously reported. The errors in the LTP as estimated by comparison of different measurements on identical samples is r 10°C. The values of the LTP for the lithia glasses are identical within experimental error and do not greatly differ from those in the literature. Values of the LTP for the 33 mole % soda glasses are within the experimental error of those reported by Morey (19 54) but the values for the 25 m.ole % soda glass are approximately 100°C higher than Morey ' s value . The thermal expansion coefficient and the LTP is essentially the same for all the lithia glasses examined v/hile the expansion coefficient and LTP for the soda glasses are markedly dependent upon composition.

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107

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CHAPTER V DISCUSSION OF RESULTS 20 Mole % Lithia-Silioa Glasses The sequence of phases observed in the 3 mole % lithia glasses by x-ray techniques during thermal treatment at 500°C is: (1) a glassy phase, (2) a glassy phase plus a srr.all (0.1 v;eight %) fraction of crystalline lithium metasilicate , (3) a glassy phase and (4) the equilibrium lithium disilicate precipitate. The term glassy phase as applied to the structure present in step 3 above indicates only that any crystals present are present in a quantity less than the lov/er limit of detection of the x-ray technique . Eleots-on Microscopy The replica electron micrographs of the as cast glass shov/n in Figure 13 exhibit a primary drop like phase separation of a silica rich phase in a lithia rich matrix. Heat treatment of this glass for 5 hours at 500°C causes a secondary phase separation to occur in the matrix as shown in Figure 14. The x-ray data in Table 2 show that the metasilicate appears concurrently with the secondary separation . 108 -

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109 The metasilicate is also observed in the electron diffraction patterns although the time of appearance is much shorter than indicated by the x-ray results. This difference is undoubtedly due to the difference in kinetics between the bulk and thin samples. The precise measurement of the temperature during thermal treatments in the electron microscope is complicated by the poor thermal conductivity of the glass and the unavoidable electron beam heating. The observed difference in the kinetics of appearance of the lithium metasilicate in bulk as opposed to thin samples is thus due to the surface effects or the unavoidable uncertainty in the sample temperature. It is evident from examination of Figures 27 and 29 that the morphology of the lithium disilicate crystal is extremely sensitive to the thermal treatment temperature. The 20°C difference in the two treatment temperatures yields a very obvious change in the axial ratio of the crystals. The growth of a precipitate out of the matrix as shown in Figure 29 could be the result of several effects peculiar to these experimental conditions. One possible reason for the growth of the crystals out of the matrix is the thermal gradient in the thin edge of the sample. This type of growth could also result from the recession of the viscous glass from around the crystalline particles. The thermal treatment temperature

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110 in this case is considerably above the softening point of the glass and the thickness of the edge under examination is in the neighbrohood of 1,000 A so that the surface driving force to cause recession of the edge is quite high. This argument is further reinforced by the observation that the sample actually sagged during examination in the microscope to the extent that the sag v/as visible to the unaided eye. The transmission electron microscopy results indicate that the structure of the glass during thermal treatment can be monitored but the uncertainty as to the surface effect and sample temperature makes a direct comparison with bulk samples difficult. The masking of the structure by the carbon film necessary to conduct av/ay charge also makes resolution of the fine detail in the early stages difficult if not impossible. It is, however, possible to follow the sequence of reactions occurring in the thinned sample. The sequence of reactions observed in the bulk samples is the same as observed in the thin samples but the morphology is apparently different. Structural Analysis of AC Properties The AC electrical properties of the 30 mole % lithia glass shown in Figures 32 and 33 show the appearance of a relaxation loss peak after 5 hours thermal treatment

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Ill at 500°C. The x-ray and electron microscopy results indicate that the lithium metasilicate precipitate appears concurrently with the appearance of this loss peak. The electron microscopy indicates that the metasilicate is dispersed in the glass matrix but the morphology of the precipitate is not clearly resolved by either the transmission or replica electron microscopy. Calculations of the peak location from the heterogeneous dielectric theory presented in Chapter I require experimental values of the conductivity, dielectric constant and volume fraction of each of the two phases involved. The value of the mietasilicate conductivity was obtained from Mazurin (1965) while the dielectric constant was measured in the present work. The values of the conductivity and dielectric constant for the glass matrix were taken to be those actually measured on this glass in the present work. Consider the loss peak located at 500 Hz shown in Figure 33 for a 30 mole % glass measured at 80°C after a 5 hour thermal treatment at 500°C. If a volume fraction of 0.01 is assumed the calculations using equation (1-7) for dispersed spheres yields a frequency of 1.8 Mhz for the peak location. This value is clearly in poor agreement with the observed value of 50 Hz. The frequency value can be modified by changing the assumed morphology or changing the values of the constants used in the equation.

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112 If the sphere model calculation above is modified by increasing the conductivity of the dispersed phase by a factor of 1.2, then the location of the peak is shifted from 1.8 Mhz to 2.2 Mhz . Examination of the effect of increasing the volume fraction from 0.01 to 0.1 yields a change in frequency from 1.8 Mhz to 1.9 Mhz. Clearly both of these alternatives shift the location of the peak in the wrong direction to help the agreement in the case above but they will be useful in evaluating the effect of thermal treatments . The alternative of inserting the experimental value of the frequency in the equation and solving for the axial ratio of the dispersed phase is difficult because the solution involves a transcendental equation in the axial ratio. This difficulty was resolved by a successive approximation technique that consisted of assuming various axial ratios and calculating the corresponding relaxation frequencies. Substituting axial ratios of 10, 10^ and 10^ into equation (1-12) gives frequency values of 7 • 10**, 3 • 10^ and 4 • 10^ respectively. Thus the AC electrical data indicate that the axial ratio of the dispersed metasilicate particles lies between 10^ and 10^. The calculations above will now be extended by analogy to the subsequent thermal treatments of the same glass in order to examine the structural implications of

PAGE 128

113 the AC rr.easurements . This requires the assumption that the AC properties of the latter thermal treatments reveal the presence of the metasilicate although the amount of metasilicate is too small to be observed by x-ray techniques . The location of the frequency maxima measured at 3 0°C for the subsequent thermal treatments shifts to higher frequencies (e.g. Figure 38). From the above discussion it can be seen that the frequency shift implies either that the axial ratio of the metasilicate is decreasing or the conductivity is increasing. As the m>etasilicate dissolves, the diffusion of lithium ions out of the metasilicate regions must in turn reduce the conductivity of the metasilicate. This would lead to a frequency shift to lower frequencies and clearly is not occurring during heat treatment. Thus, the change of conductivity during thermal treatment is dismissed as improbable. The other possibility of increasing the volume fraction of lithium metasilicate during thermal treatment causes the predicted frequency to shift in the right direction but an increase in the volume fraction of the metasilicate is not observed in the x-ray results. Therefore, it is concluded that the axial ratio changes during thermal treatment.

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114 DC Conduotivity The DC conductivity results for the 3 mole % lithia glass shown in Figure 43 indicates that the DC conductivity is unchanged until after the AC loss peak appears. The DC conductivity of the glass is, to a first approximation, governed by the composition of the matrix phase if the dispersed phases have a small volume fraction, The evidence presented above indicates that the metasilicate occupies a small volume fraction so the above approximation should be valid. The formation of the metasilicate necessarily depletes the matrix of lithia but the volume fraction is so small that the depletion is not observed in the DC conductivity. This is the reason that the loss peak appears but no change in the DC conductivity is observed. The precipitous drop in conductivity after the loss peak appears indicates that the equilibrium phase is beginning to appear and the matrix material is being seriously depleted in lithia. Thermal Expansion and Transformation Points The thermal expansion and lower transformation points for the 30 mole % lithia glass were observed (Table 6) to be unchanged by thermal treatments up to 50 hours at 500°C. This indicates that the matrix phase

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115 which governs these properties is unchanged by thermal treatments up to 50 hours at 500°C insofar as the thermal expansion and softening points are concerned. Relationship of Observed Morphology to Phase and Free Energy Diagrams The explanation of the sequence of phases and morphologies observed is based on the phase diagram (Figure 49) from Kracek (1939), Holmquist (1961) and Glasser (1967) and a free energy composition diagram (Figure 50) . The free energy diagram was constructed from free energy values for the stoichiometric compounds obtained by assuming a linear relationship to the corresponding soda-silica compounds whose free energies are reported in the literature. The free energy of formation of lithium metasilicate, sodium disilicate and sodium metasilicate was obtained from Richardson, Jeffes and Withers (1950) . The direction of change of free energy for the nonstoichiometric compounds was inferred from first principles, while the relative values for the liquids were inferred from the phase diagram. The phase diagram includes one feature, the metastable miscibility gap, which is not generally included on an equilibrium diagram. The "S" shaped liquidus curve suggests that a metastable miscibility

PAGE 131

116 LI,0-SI0, After Kracek. Holmquist, and Glasser. > Quartz P Crislobalite T Tridymlte I LljO-2SlO, 11:21 « U,OSiO, 11:11 SIO, U,0 2S1O2 UJQ ImolaM Li]0 SiO, Figure 49. — Phase diagram for the lithiasilica system.

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117 ^ 10 cs u >.-10 « e -20 -zzv liquid 47 Metasilicsle Oiiilicate 25 LijO (Molc?^o) 50 Figure 50. — Free energy-composition diagram for the lithia-silica system at 500°C.

PAGE 133

118 gap will occur in the system and the work of Charles (1966, 1967) and others has confirmed the appearance of this metastable reaction. The metastable miscibility gap is represented by a dotted line on the equilibrium diagram for the purpose of simplifying the explanation of the evolution of the glass structure. Consider first the 30 mole % lithia glass above the liquidus temperature. As the glass is cooled below the liquidus temperature it encounters the metastable miscibility gap and separates into two immiscible liquids of lithia rich and lithia poor composition. The position of the gross composition with respect to the location of the miscibility gap indicates that the lithia rich phase will have the larger volume fraction. The evidence discussed earlier indicates that the lithia rich phase is the matrix and the separated phase of lower volume fraction is the silica rich phase. The cooling rate will affect the size of this initial phase separation but the structure after an ordinary casting procedure is that shown in Figure 13. The effect of a thermal treatment on this glass at 500°C according to the phase diagram is to produce a homogeneous lithium disilicate crystal. If, however, we examine the free energy composition diagram (Figure 50)

PAGE 134

119 for the heat treatment temperature it is evident that several intermediate steps can occur. The free energy of the glass with the heterogenities discussed above is represented by a tie line connecting points 1 and 2 on the free energy diagram in Figure 50. As an intermediate step in the crystallization of lithium disilicate, precipitation of a metasilicate crystal from the matrix of composition 2 will reduce the free energy of the m.atrix from 2 to 7 on the diagram. This lowers the overall free energy of the 'glass' from 3 to 6 on the diagram. Further crystallization of the equilibrium disilicate crystal leads to an overall reduction in free energy from 6 to 8 on the diagram. . . Thus the appearance and disappearance of the lithium metasilicate and the appearance of the equilibrium lithium disilicate is rationalized by reference to the phase diagram and the free energy-composition diagram. S3 Mole % LithiaSilica Glasses The AC loss spectra of the 33 mole % lithia glass shown in Figure 3 9 is similar to that for the 3 mole % lithia glass shown in Figure 38. The initial absence of a loss peak in the as cast glass followed by the appearance of a loss peak with thermal treatment at 500 °C and its subsequent reduction in height indicates that the same

PAGE 135

120 processes are occurring in this glass as were shown by more detailed investigation to occur in the 30 mole % lithia glass . The x-ray analysis does not detect the metasilicate phase as the loss peak appears in this glass but it appears from the AC relaxation behavior that the metasilicate is present. The DC results from the 33 mole % lithia glass also show the same general behavior as the 30 mole % lithia glass. It is thus concluded that the series of phases observed in the 30 mole % lithia glasses also appears in the 33 mole % lithia glasses. 26.4 Mole % Lithia-Silica Glasses The AC behavior of the 26.4 mole % lithia glass is also quite similar to the 30 mole % lithia glass but the time required to develop the loss peak is longer than in the 30 mole % glass. The metasilicate phase was not detected by x-ray techniques but once again this does not preclude its presence in small quantity. It is concluded from the AC behavior that the metasilicate is present but its volume fraction is less than that required for detection by x-ray techniques. The DC conductivity behavior of this glass with thermal treatment is quite similar to the 30 mole % lithia

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121 but the precipitous drop in conductivity observed between 5 and 10 hours in the 30 mole % glass is not observed until between 50 and 100 hours in the 26.4 mole % lithia glass . Summary of Lithia Glasses The experimental results for the 30 mole % lithia glass are summarized in Figure 51. The bar graphs representing each of the experimental results indicate that the sequence of phases observed is: (1) glass, (2) glass and metasilicate and (3) glass and disilicate. The kinetics of appearance of the metasilicate shown in Figure 52 are fastest in the least refractory (33 mole % lithia glass and slowest in the 26.4 mole % glass. The crystallization incubation time at 595°C reported by Freiman and Hench (1968) is 50 minutes for the 26.4 mole % glass and 15 minutes for the 33 mole % glass. This supports the conclusion that the metasilicate which nucleates crystallization requires a longer time to form in the more refractory glass. Freiman and Hench 's growth kinetics data for the 33 mole % glass after nucleation for 3 hours at 475°C and 24 hours at 475°C show that the incubation period after the longer nucleation time at 475°C is reduced and the activation energy for crystallization is reduced from 92

PAGE 137

122 CliSS XcusS'KS y^ CUM ci»ss t «n*siiic»i( v^ci*>s I Hi > OS GilSS • OS *iS / Cl*ss . HS bktss ft 0
PAGE 138

123 o I 0) E o CO o CM 35 Li20 (r^olo%) Figure 52. — Thermal treatment time required to develop loss peak versus composition for lithiasilica glasses.

PAGE 139

124 to 52 kilocalories/mole by the longer nucleation treatment. The above discussions lead to the conclusion that the nucleating agent in the crystallization of lithia glasses is the metastable raetasilicate crystal. S3 Mole % Soda-Siliaa Glasses The AC loss spectra of the 33 mole % soda glass shown in Figure 41 is similar in appearance to the loss spectra of the lithia glasses discussed above. The principal difference in the behavior of the two types of glasses lies in the appearance and growth of the loss peak in the soda glasses while in the lithia glasses the loss peak appeared suddenly and declined in magnitude with further heat treatment. The x-ray results for this glass do not detect any phases other than the equilibrium sodium disilicate. It is the opinion of the author that the development of the loss peak in this glass is the result of the appearance of the sodium metasilicate precipitate. The phase diagram of the soda-silica system from Kracek (1939) and Holmquist (1961) is shown in Figure 52. The equilibrium phase that should develop in the 33 mole % glass is the sodium disilicate. The free energy-composition diagram in Figure 54 was constructed from the data of

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125 NOjO-SiOj After Kracek and Holmquist a. Quartz fi Cristobclite V Tridymite « Na,OSiO, (I 60 NajQSiQj NajO(mole %; Figure 53. --Phase diagram for the sodasilica system.

PAGE 141

126 20

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127 Richardson, Jeffes and Withers (1950) in the same manner as the lithia-silica free energy diagram discussed previously. Examination of the free energy composition diagram shows that the decomposition of the homogeneous as cast glass into a silica rich glass and a sodium metasilicate crystal reduces the free energy of the system from 1 to 4 as shown on Figure 53 . This step is analagous to the separation of lithium metasilicate in the lithiasilica glasses and is an intermediate step in the precipitation of the sodium disilicate crystal. The analysis of the AC loss spectra could be carried out in a manner analagous to that for the 30 mole % lithia glass but the AC loss spectra will be discussed by analogy to the calculations above rather than by a detailed calculation. The appearance of the loss peak after 5 hours at 550°C indicates that the sodium metasilicate phase has appeared. The shift of the loss peak to higher frequencies with heat treatment indicates that the axial ratio of the precipitate is decreasing or the conductivity of the precipitate is decreasing. The DC conductivity of the 33 mole % soda glass does not change during thermal treatments of 5 hours at 550°C and as a result no information as to the structural changes occurring can be obtained from the DC data.

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128 Therefore, it is concluded that the sodium metasilicate appears as a precursor to the equilibrium phase in the crystallization of a 33 mole % soda-silica glass. 25 Mole % Soda-Siliaa Glasses The loss spectra of the 25 mole % soda glass are markedly different from the other glasses examined. The loss peak observed in the other glasses is not observed in this glass. A metastable phase was not detected in this glass thus, no evidence was obtained to indicate that the metastable precipitate forms in the 25 mole % soda glass . The absence of the metasilicate as a transition phase in the crystallization of the 25 mole % glass can be explained by examination of Figure 53. The magnitude of the free energy reduction when the metasilicate precipitates in the 33 mole % glass is represented by the line from 1 to 4 . As the composition of the glass is moved to the left toward the 25 mole % glass the length is reduced and the tendency for this intermediate step to occur is reduced. Summary of Soda Glasses The discussion above indicates that the metasilicate precipitate forms in the 33 mole % glass but not in the

PAGE 144

129 25 mole % glass. This is rationalized on the basis of free energy considerations. The discussion above indicates that in some soda glasses the metastable sodium metasilicate precipitate is a precursor to equilibrium crystallization .

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CHAPTER VI SUMMARY The conditions which lead to polarization effects in electrical measurements in an ionically conducting m.aterial have been shown to be a critical combination of high conductivity of the material and long polarization times. A model consisting of a uniform negative volume charge with a uniform positive surface charge has been developed to explain the polarization behavior. The model yielded theoretical predictions of the polarization charge which agreed within 40% of the observed values. The experimental results for thermally treated lithia-silica glasses revealed the appearance of metastable lithium metasilicate crystals prior to the appearance of the equilibrium lithium disilicate precipitate. The precipitation and dissolution of the metastable lithium metasilicate was observed by transmission electron microscopy, x-ray and AC electrical measurement techniques. The appearance and resorption of this metastable precipitate was explained by employing a free energy-composition diagram. The 33 mole % soda-silica glass examined exhibited the appearance and dissolution of a metastable precipitate 130 -

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131 prior to the equilibrium precipitation in a manner similar to that observed in the lithia-silica glasses. The 25 mole % soda glass examined did not show evidence of the metastable precipitate. This behavior was shown to be consistent with a free energy-composition diagram in a manner similar to that used to explain the lithia-silica behavior. The discussion below presents evidence that the above reaction sequence explains the bulk crystallization kinetics of glasses. It was observed in the lithia-silica glasses that the time required to form the metastable precipitate increased with the refractoriness of the glass. The lithium metasilicate develops most rapidly in the Last refractory 33 mole % lithia glass and slowest in the most refractory 26.4 mole % lithia glass. A recent investigation (Freiman and Hench 196 8) has shown that the incubation time for bulk crystallization of lithia-silica glasses also increases with the refractoriness of the glasses. Consequently it can be concluded that the relative rate of the development of the lithium metasilicate crystals during the nucleation step is responsible for the compositional effects in the bulk crystallization of the lithia-silica glasses.

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132 The duration of the nucleation treatment has also been shown to change the rate of bulk crystallization (Freiinan and Hench 1968). They reported that an increase of nucleation time from 3 to 24 hours at 475°C increased the crystallization rate by a factor of 3. The present study shows that the concentration of the lithium metasilicate crystals increases with heat treatment time. A critical time in excess of 5 hours at 500 °C is necessary for the metasilicate to appear. Therefore, the 3 hour nucleation treatment above was insuf f icienct to develop the metasilicate v/hile the 24 hour treatment developed the metasilicate nuclei. The subsequent crystallization behavior of the glass with and without the metasilicate nuclei present is a clear example of crystallization in a glass with and without pre-existing nucleation sites. Freiman and Hench also reported that the activation energy of bulk crystallization rate in a sample nucleated 24 hours at 475°C was identical to the activation energy for crystal growth. The activation energy for bulk crystal growth after a 3 hour nucleation treatment at 475°C was a factor of 2 larger than in the sample nucleated 24 hours at 475°C. The present work indicates that this difference is the result of the presence of the metasilicate as a nucleating agent in the nucleated glass. Thus the appearance of the metasilicate as a

PAGE 148

133 nucleating agent controls the activation energy of . nucleation and thereby the temperature dependence of bulk crystallization. The microstructure control presently used in commercial glass-ceramics involves the empirical addition of a nucleating agent such as titania or zirconia. The nucleation is thus artifically introduced by an intentional additive. The nucleation treatments of the commercial glasses are generally conducted at higher temperatures than those used in the present work. The results of the present investigation indicate that a nucleating agent need not be intentionally added to these glasses but that the fine grain size desired can be developed by nucleation of metastable phases followed by crystal growth at an elevated temperature.

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CHAPTER VII CONCLUSIONS 1. The lithium metasilicate phase has been observed in three lithia-silica glasses as a transition phase in the development of the lithium disilicate precipitate . 2. There is evidence of the sodium metasilicate phase in one of the two soda glasses examined and it appears that it is a precursor to the crystallization of the equilibrium phase in this glass. 3. The metastable phases observed in alkalisilicate glasses appear to be the nucleating agent for the equilibrium precipitation in these glasses. 4. The AC property measurements have been shown to be a valuable tool for examination of the structure of glasses where the relative conductivities of the phases involved is suitable to obtain Maxwell-Wagner-Sillars losses. 5. The hot stage transmission electron microscopy has shown the sequence of phases occurring in the 30 mole % lithia-silica glass and this sequence of phases has been confirmed by other techniques. 134 -

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135 6. The polarization results establish a well defined method for avoiding electrode polarization effects in DC conductivity measurements.

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CHAPTER VIII RECOMI'IENDATIONS TO FUTURE INVESTIGATORS The hot stage transmission electron microscope is an ideal tool for continuing the investigation of the structural details of the metastable phase in the very early stages of crystallization. The problem of carbon masking the structure encountered in the present work might be solved by using silicon monoxide or a metal layer to conduct away the electrical charge. It is also recommended that lower temperatures or thermal cycling be utilized to allow careful study during the process. It is recomniended the x-ray small angle scattering technique be employed as an independent source of structural parameters to allow a check of the electrical property interpretation and the transmission electron microscopy. The temperature dependence of the loss peak appearance is an obvious area for further research which would yield activation energies for the metastable precipitation and allow a more detailed structural model for the crystallization process to be developed. The activation energy for the development of the peak in the 30 mole % lithia glass would allow analysis of the species 136 -

PAGE 152

137 diffusing during the formation of the metasilicate and thus provide the basis for an atomistic model. Experiments conducted in a high temperature Guinier-DeVJolf f x-ray camera might allow the metastable phases to be detected in the glasses which post facto experiments in this work did not detect. The high temperature cair^era would have the advantage that the crystal present would be examined at the time when it is present in its largest quantity. This work might also be expanded to examine the effects of ternary additions and allow application of the above results to the practical problem of optimizing the grain size in glass-ceramics as discussed in the introduction.

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BIBLIOGRAPHY Aver'yanov, V. I. and Porai-Koshits , E. A., 1966, Proa. Seventh Alt-Un-ion Conferenae on the Glassy State ^ Leningrad, 1966. {The Structure of Glass ^ 6, 98, Consultants Bureau, New York, 1966). Barton, M. , 1965, 67th Annual American Ceramic Society Meeting, Unpublished. Cahn, J. W., 1968, TAIME , 242, 166. Cahn, J. W. and Charles, R. J., 1965, Phys. and Chem. of Glasses , 6 , 181 . Charles, R. J., 1963, J. Am. Ceram. Soc, 46, 235. Charles, R. J., 1967, J. Am. Ceram. Soc, 50, 631. Cohen, J., 1957, J. Appl. Phys., 28, 795. Debye, P., 1929, Polar Molecules (Chemical Rubber). DeVJolff, P. M. , 1947, Appl. Sci. Res., Bl, 119. Foex, M. , 1944, Compt. rend.^ 218, 196. Fousserau, G., 1883, Compt. rend., 96, 785. Freiman, S. V7. and Kench, L. L., 1968, J. Am. Ceram. Soc., 51, 382. Fricke, H. , 1953, J. Phys. Chem., 57, 934. Frohlich, H. , 1958, Theory of Dielectrics (Oxford University Press) . Fulda, M. , 1957, Sprechsal, 60, 769. Glasser, F. P., 1967, Phys. and Chem. of Glasses, 8, 224. Guinier, A., 1956, Theory et Technique de la Radiocry stallographie (Dunod) . Haller, V?., 1965, Jnl. Chem. Physics, 42, 686. 138 -

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139 BIBLIOGRAPHY — Continued Keroux, L. , 1958, J. Appl. Phy s . , 29, 1639. Holmquist, S. B,, 1961, J. Am. Ceram. Soc, 44, 82. Kopkinson, J., 1876, Phil. Trans.:, 166, 489. Hughes, K. and Isard, J. 0., 1968, Phy s . and Chem. of Glasses J 9 , 37 . Isard, J. 0., 1962, Proa. Inst. Elect. Engrs. (London) , 109B, 440. Jovner, B. and Bell, W. , 1953, J. Am. Ceram. Soc, 36, 263. Kirby, P. L., 1950, Soi. Progr., 38, 257. Kracek, F. C, 1939, J. Am. Chem. Soc, 61, 2870. Krans, C. A. and Darby, E. E. , 1922, J. Am. Chem. Soc, 44, 2783. Kohlrausch, R. , 1847, Ann. Physik. Lpz., 12, 393. Lebedev, A. A., 1940, Bull. Acad. Soi. USSR (Physical Series ) , 4 , 137 . Le Blanc, M. , and Kerschbaum, F. , 1910, Z. Physik. Chem., 72, 468. Leko, V. K. , 1967, Bull. Acad. Sci. USSR (Inorganic Matls. Series), 3, 1224. Littleton, J. T. and Morey, G. W. , 1933, The Electrical Properties of Glass (Wiley) . Littleton, J. T. and Wetmore, W. L., 1936, J. Am. Ceram. Soc, 19, 243. Mackenzie, J. D., 1960, Modern Aspects of the Vitreous State (Butterworth) . Maxwell, J. C., 1892, Electricity and Magnetism (Clarendon)

PAGE 155

140 BIBLIOGRAPHY — Continued Kazurin, 0. V., 1965, Proc. Fifth All-Union Conference on z'ne Glassy State, Leningrad, 1965. {The Structure of Glass, 4, 5, Consultants Bureau, New York, 1965). McCurrie, R. A. and Douglas, R. W. , 1967, Phys. and Chem. of Glasses J 8, 132. McDowell, L. S. and Begemann, H. L. , 1929, Phys. Rev., 33 , 55 . Moore, H. , and DeSilva, R. L., 1952, J. Soc. Glass Tech. ^ 36, 5t. Morey, G. W. , 1954, The Properties of Glass (Reinhold) . Mulligan, M. J., Ferguson, J. B. and Rebbeck, J. W. , 1925 J. Am. Ceram. Soc., 8, 329. Ov;en, A. E., 1961, Phys. and Chem. of Glasses, 2, 87. Ov;en, A. E. , 1963, Progress in Ceramics Science, 3, edited by J. E. Burke (MacMillan) . Page, L. and Adans , N., 1958, Principles of Electricity (van Nostrand) . Poole, H. H., 1921, Nature, 107, 584. Prebus, A. E. and Michener, J. W. , 1954, Ind. Eng . Chem., 16, 147. Proctor, T. M. and Sutton, P. M. , 1959, J. Chem. Phys., 30, 212. Proctor, T. M. and Sutton, P. M. , 1960, J. Am. Ceram. Soc, 43, 173. Prod'homine, L., 1960, Verres et Refractories, 14, 24. Reaumur, M. , 1739, Memoires de l^Academie des Sciences, 370. Richardson, F. D., Jeffes, J. and Withers, G. , 1950, Jnl. Iron and Steel Inst., 166, 213.

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141 BIBLIOGRAPHY — Continued Robinson, D. M. , 1932, Physics, 2, 52. Seward, T. P., Uhlraann, D. R. , Turnbull, D. and Pierce, G. R. , 1967, J. Am. Ceram. Soa., 50, 25, Shaw, R. R. and Uhlraann, D. R. , 1968, J. Am. Ceram. Soa . , 51, 377. Sillars, R. W. , 1937, Proc. Inst. Elec. Engrs. (London)^ 80, 378. Slayter, G., 1952, Bull. Am. Ceram. Soc._, 31, 276. Smith, J. v., 1967, Powder Diffraction File (ASTM) . Smyth, C. P., 1955, Dielectric Behavior and Structure (McGraw-Hill) . Stevels, J. M. , 1946, J. Soc. Glass Tech., 30, 31t. Stevels, J. M. , 1950, J. Soc. Glass Tech., 34, 80t. Stookey, S. D., 1962, Symposia on Nucleation and Crystallization in Glasses and Metals (American Ceramic Soc. ) . Strutt, xM. J., 1931, Arch. Elektvotech, 25, 715. Taylor, H. E. , 1957, J. Soc. Glass Tech., 41, 350t. Tegetmeier, F. , 1890, Ann. Physik. Chem. , 41, 18. Tran, T. L., 1965, Glass Tech., 6, 161. van Beek, L. K. H. , 1967, Progress in Dielectrics, 7, edited by J. B. Birks (Chemical Rubber), p. 69. Vogel, W. and Byhan, H. G., 1964, Silikattechnik, 15, 212 Wagner, K. W. , 1914, Archiv fur Elektuotechnik, 2, 371. Warburg, E. , 1884, Ann. Physik., 21, 622. Warren, B. E. , 1937, J. Appl. Phys., 8, 645.

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142 BIBLIOGRAPHY — Continued Williamson, J., and Glasser, F. P., 1966, Phy s . and Chem. of Glasses J 7, 127. Zachariasen, W. H. , 1932, J. Am. Chem. Soa.^ 54, 3841.

PAGE 158

BIOGRAPHICAL SKETCH Donald L. Kinser was born September 28, 1941, in Loudon, Tennessee. He received his early schooling in Alabama and Florida and was graduated from DeLand High School, DeLand, Florida in 1959. In September, 1959 he entered the University of Florida, receiving a Bachelor of Metallurgical Engineering degree in April, 1964, and a Doctor of Philosophy in 1968. The author is married to the former Barbara E. Lange and is the father of two children, Elizabeth and Cynthia. He is a member of American Institute of Mining, Metallurgical and Petroleum Engineers, American Society for Metals, American Ceramic Society, Society of Glass Technology, Electrochemical Society, AZM and ZT. 143

PAGE 159

This dissertation was prepared under the direction of the chairman of the candidate's supervisory committee and has been approved by all members of that committee. It was submitted to the Dean of the College of Engineering and to the Graduate Council, and was approved as partial fulfillment of the requirements for the degree of Doctor of Philosophy. December, 1968 , College of Engineeri/g Dean, Graduate School Supervisory Committee: Cfiaidrmaii 7<.'7rA C. n,^,0, 0^^^

PAGE 160

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