THE EFFECT OF P205 ON THE DEVITRIFICATION! '
0F LEAD SILICATE GLASSES

Dissertation for the Degree of Ph. D.
MICHIGAN STATE UNIVERSITY
SANDRA HAARALA CARR
1977

 

 

Tm T T T A Mi ii iii L i
University

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

This is to certify that the
thesis entitled
THE EFFECT OF P205 ON THE DEVITRIFICATION
OF LEAD SILICATE GLASSES

presented by

Sandra Haarala Carr

has been accepted towards fulfillment
of the requirements for

Jh D degree in We ience

 

- \Y- Mama/uh cw

Major professor

 

Date February 25 1977

0-7 639

 

'll‘ .

 

 

ABSTRACT
THE EFFECT OF P205 ON THE DEVITRIFICATION
OF LEAD SILICATE GLASSES
By

Sandra Haarala C arr

The effect of P205 on the devitrification of binary lead silicate
glasses containing 64 mole percent and 59 mole percent PbO was
studied. Glasses containing 64 mole percent PbO were prepared with
P205 concentrations of 0. 5 and l. 0 mole percent, and glasses con-
taining 59 mole percent PbO were prepared with P205 concentrations
of 0. l and O. 5 mole percent. These glasses were devitrified at
temperatures of 400°C, 450°C, 500°C, and 5500c. The crystallization
products were determined by x-ray analysis, and the crystal growth
rates and microstructures were determined by electron and optical
microscopy.

The major product of crystallization for all compositions of
glasses studied was a polymorph of 3PbO'ZSiOZ. The 64 mole percent
PbO glasses contained a low temperature polymorph of ZPbO- SiO2 as
an additional phase, and the 59 mole percent PbO glasses contained
a polymorph of PbO' SiOZ.

Concentrations of 0. 5 and 1. 0 mole percent P205 promoted in-
ternal crystallization in the form of spherulites. The maximum nucle-
ation rates for Spherulitic crystallization occurred at 400°C. They
were determined to be 3. 29 x 106 spherulites-cm- 3-min-1 and
Z. 10 x 104 spherulites-cm- 3-min-1 for the 64 mole percent PbO glass
containing 1. O and O. 5 mole percent P205, respectively; and about

3 x 105 spherulites-cm- 3--min-1 for the 59 mole percent PbO glass con-

taining 0. 5 mole percent P205. Spherulitic growth rates were constant

with time “7
tion energ.V
ite d8V010P-
characteris‘
morphology.
textures at :
litic growth

well for the

Sandra Haarala Carr

with time for isothermal heat treatments. The experimental activa-
tion energy for Spherulitic growth was about 84 kcal/mole. Spheru-
lite development followed the sheaf—to-spherulite sequence that is
characteristic of spherulites in other materials. The spherulite
morphology, which varied with temperature, displayed coarse, open
textures at 550°C and fine textures at 400°C. The model of spheru-
litic growth developed by Keith and Padden45 accounts reasonably

well for the Spherulitic growth mode.

TH?

 

 

THE EFFECT OF P205 ON THE DEVITRIFICATION
OF LEAD SILICATE GLASSES

By

Sandra Haarala C arr

A DISSERTATION

Submitted to
Michigan State University
in partial fulfillment of the requirements

for the degree of
DOCTOR OF PHILOSOPHY
Department of Metallurgy, Mechanics, and Materials Science

1977

to the memory of my father

ii

 

 

 

I we
are respo
rm .
Lfigness tr
the 1
Donald J,
and Suppo
Dr.
Who go out
students;
my F
faith and 8
my h
neSS and a
and,

and encOur

AC KNOWL EDGM ENT S

I would like to give special thanks to the following people, who
are responsible in one way or another for the completion of this work:

my advisor, Dr. K. N. Subramanian, for his guidance and will-
ingness to help, regardless of the time of day (or night);

the faculty and staff of the MMM Department, especially Dr.
Donald J. Montgomery and Dr. Robert Summitt, for their friendship
and support;

Dr. Gary Hooper and his associates at the Electron Optics Lab,
who go out of their way to assist and co-operate with campus graduate
students;

my parents, William and Henrietta Haarala, for their unfailing
faith and support;

my husband, Dale, for his resourceful "handy-man" innovative-
ness and amazing ability to live with a graduate student wife;

and, finally, my fellow graduate students for their friendship,

and encouragement, over the years.

iii

 

 

——-_—_—_——————
—

A—fi— ._ —_ _.

LIST OF TABII
LIST OF FlGl’Ii
Chapter
I. INTRO
II. THE( )1‘

III. THE

TABLE OF CONTENTS

LIST OF TABLES .
LIST OF FIGURES .

Chapter
I.
II.

III.

VI.

INTRODUCTION
THEORY .

2.1 Classical Nucleation Theory
2.2 Growth Processes.

2.3 Crystallization in Glass- -Forming Systems .

THE LEAD SILICATE SYSTEM .

3. 1 Compound Formation
3. 2 Rate of Crystallization

EXP ER IM ENT AL PROC EDUR E

4.1 Glass Preparation .
4. 2. Devitrification of Glasses .
4. 3 Methods of Observation and Analysis

RESULTS AND DISCUSSION

Crystallization Products
Microstructure of Glasses .
Spherulitic Growth Rate .
Spherulitic Morphology
Role of P205

SUMMARY .

5
5
5.
5
5

Uiuwan-a

APPENDIX

A.

HISTOGRAMS OF SPHERULITE POPULATIONS
IN DEVITRIFIED GLASSES

LIST OF REFERENCES

iv

Page

vi

10
12

16

16
23

26

2.6
28
29

33

33
43
58
66
75

83

85
108

II{IIIIII[II|‘I"IIIll{llll Ten III‘ .11

 

Table

LIST OF TABLES

The d- spacings and relative intensities of the

x-ray spectra for the major crystalline phase . .

The d-spacings and relative intensities of the
x-ray spectra for the low temperature phase
in devitrified 70- 30 glass .

The d- spacings and relative intensities of the
x-ray spectra for the secondary crystalline
phase in devitrified 60-40 and 60-40-0. 5 glass

Crystal growth rate as a function of
temperature and glass composition

Page

36

37

44

64

 

‘lI‘Illlll-l Iitilll

LIST OF FIGURES

Figure

1.

10.

ll.

12.

Rates of nucleation and growth for viscous

liquids as a function of the degree of undercooling .

Variation of Gibbs Free Energy with radius for
condensation of a spherical droplet in a
supersaturated vapor .

Reaction paths between the initial
glass and the final glass-ceramic

Phase diagram of the lead
silicate system by Geller et a1.

 

Phase diagram of the lead
silicate system by Ott and McLaren

Phase diagram of the lead
silicate system by Smart and Glasser

Microtome set-up for sectioning glass

X-ray diffraction spectra of 70-30 glass

devitrified for (a) 7 days at 400°C, (b) 6 days
at 450°C, (c) 3 days at 500°C and (d) 3 days
at550°C

X-ray diffraction spectra of 70-30- 1. 0
glass devitrified for (a) 10 days at 400°C
and (b) 3 days at 550°C . . . .

X-ray diffraction spectra of (a) 70- 30 glass
devitrified for 3 days at 550°C and (b) 70-30-1. 0
glass devitrified for 3 days at 550°C

X-ray diffraction spectra of 60- 40 glass
devitrified for (a) 18 days at 400°C, (b) 6 days
at 450°C, (c) 3 days at 500°C, and (d) 3 days
at550°C
X-ray diffraction spectra of 60- 40-0. 5 glass
devitrified for (a) 18 days at 400°C and (b)
3daysat550°C

vi

Page

14

17

22

24
31

34

4O

41

42

45

Figure Page

13. Electron micrographs of replicas of quenched
(a) 70-30 glass, (b) 70- 30-0. 5 glass,
and (c) 70-30-1.0 glass . . . . . . . . . . . . . . . 46

14. Electron micrograph of radial crystal growth
pattern in quenched 70- 30-0. 5 glass . . . . . . . . . 48

15. Electron micrograph of phase separated region
between crystalline growth front and matrix
in quenched 70-30-0. 5 glass . . . . . . . . . . . . . 49

16. Micrographs of surface crystallites in
(a) 70-30 glass, (b) 70-30-0. 5 glass, and
(c) 70- 30-1. 0 glass after devitrification
at400°Cfor42hours................ 50

17. Micrograph of surface crystals in 70- 30-0. 5
glass after devitrification for 1 hour at 400°C
followed by 12 hours at 450°C . . . . . . . . . . . . 52

18. Micrographs of (a) 60-40 glass devitrified at
550°C for 1 hour and (b) 70-30 glass devitrified
at 550°C for 72 hours . . . . . . . . . . . . . . . . 53

19. Microstructure of crystalline surface layer in
(a) 60-40-0. 1 glass devitrified for 1 hour at
550°C and (b) 60-40-0. 5 glass devitrified
for72hoursat550°C................ 55

20. Microstructures after devitrification at 550°C
for 72 hours of (a) 60-40-0. 5 glass and

(b) 70-30-1. 0 glass . . . . . . . . . . . . . . . . . 56
21. Microstructure of 60-40-0. 5 glass after
devitrification at 500°C for 72 hours . . . . . . . . . 57

22. Micrographs of spherulites in (a) 60-40-0. 5
glass devitrified for 14 days at 400°C, (b) 70-30-0. 5
glass devitrified for 8 days at 400°C, and (c) 70-30-1. 0

glass devitrified for 4 days at 400°C . . . . . . . . . 59
23. Growth curves for spherulites in 70- 30-0. 5

glass at (a) 400°C, (b) 450°C, (c) 500°C, and (d) 550°C. . 61
24. Growth curves for spherulites in 70- 30- l. 0

glass at (a) 400°C, (b) 450°C, (c) 500°C, and (d) 550°C. . 62
25. Growth curves for spherulites in 60-40-0. 5

glass at (3) 400°C, (b) 450°C, (e) 500°C, and ((1) 550°C, . 63

26. Graph of ln(crystal growth rate) versus the
reciprocal of the absolute temperature . . . . . . . . 65

vii

Figure

27.

28.

29.

30.

31.

32.

33.
34.

35.

36.

Spherulites grown at 550°C showing typical
(a) sheaf, (b) bundle, and (c) hexagonal
morphologies . . . . . . . . .

Diagram of spherulite orientation .

Micrograph of a spherulite in 70-30-0. 5
glass after growth at 450°C

Electron micrograph of fibrous growth morphology
in spherulites grown at 400°C in 70-30-0. 5 glass .

Micrograph of an early growth stage in a
spherulite grown at 550°C

Electron micrograph of branching fibers
in a spherulite grown at 550°C

Micrograph of a mature spherulite grown at 55 0°C

Network of inter- Spherulitic boundaries in a
completely devitrified 70-30-0. 5 glass

Free energy of an interface versus
occupied fraction of surface sites

Electron micrograph of faceted crystal growth in a
60-40-0. 1 glass after devitrification at 450°C
for 1 hour

Histograms of spherulites grown in 70-30-0. 5
glass at 400°C for (a) 2 days, (b) 3}; days, (c) 4 days,
(d) 5 days, (e) 6 days, and (f) 14 days . . . .

Histograms of spherulites grown in 70-30-1. 0
glass at 400°C for (a) 2 days, (b) 3} days, and
(c)4days................

Histograms of spherulites grown in 60-40-0. 5
glass at 400°C for (a) 14 days, (b) 18 days,
(0) 25 days, and (d) 32 days . . .

Histograms of spherulites grown in 70-30-0. 5
glass at 450°C for (a) 200 min. , (b) 240 min. ,
(c) 360 min. , (d) 480 min. , and (e) 600 min. ,
after a nucleation heat treatment of 1 hour at 400°C .

Histograms of spherulites grown in 70-30-1. 0
glass at 450°C for (a) 200 min. , (b) 275 min. ,
(c) 480 min. , (d) 600 min. , and (e) 720 min.

viii

Page

67
69
7O
71

72

73
74

76

78

81

87

89

90

92

94

 

1°11, Iii III||IIII|I [III

II I Ilal'lli'lllf"

 

Figure

A. 6.

A. 10.

A. 11.

A. 12.

Histograms of spherulites grown in 60-40-0. 5
glass at 450°C for (a) 720 min. , (b) 1080 min. , and
(c) 1405 min. after a nucleation heat treatment of

1 hour at 400°C; and for (d) 1440 min. , (e) 1680 min.,

and (f) 1920 min. after a nucleation heat treatment
of 6 hours at 400°C .

Histograms of spherulites grown in 70-30-0. 5
glass at 500°C for (a) 10 min. , (b) 15 min. ,

(c) 22 min. , (d) 30 min. , and (e) 37 min. , after
a nucleation heat treatment of 1 hour at 400°C .

Histograms of spherulites grown in ‘70- 30-1. 0

lass at 500°C for (a) 10 min. ., (b) 15 min. ,
fc) 22 min. , (d) 30 min. , and (e) 37 min. ,

Histograms of spherulites grown in 60-40-0. 5

glass at 500°C for (a) 20 min. , (b) 40 min. ,

(c) 60 min. , (d) 80 min. , (e) 100 min. , (f) 120 min. ,
and (g) 140 min. , after a nucleation heat treatment
of 6 hours at 400°C .

Histograms of spherulites grown in '70- 30-0. 5
glass at 550°C for (a) 3 min. and (b) 5 min. , after a
nucleation heat treatment of 1 hour at 400°C

Histograms of spherulites grown in 70-30-1. 0
glass at 550°C for (a) 3 min. , (b) 5 min. , and
(0)10min................

Histograms of spherulites grown in 60-40-0. 5 glass

at 550°C for (a) 15 min. , (b) 22 min. , and (c) 30 min. ,
after a nucleation heat treatment of 6 hours at 400°C .

ix

Page

96

98

100

102

105

106

107

C HAPT ER I
INTRODUCTION

The PbO-SiO2 combination forms the basis for many industrially
important polycomponent systems, including flint glasses, pottery
glazes, enamels, and solder glasses for metal-to-glass seals. Much
of the usefulness of this system is directly dependent upon its devit-
rification properties. Solder glasses, for example, require high flu-
idity and rapid devitrification at low to moderate temperatures. The
lead silicate system, especially those compositions high in lead oxide
content, has proven extremely useful for these applications.

In the lead silicate system, devitrification is rapid for composi-
tions of high lead oxide content and usually starts at the glass surface
and proceeds towards the interior. Devitrification via surface crystal-
lization is often accompanied by sample deformation and cracking, and
usually results in weak materials with degraded mechanical properties.
If crystallization can be made to proceed by internal (or bulk) crystal-
lization, however, the degradation of the mechanical properties can
often be avoided. Recent work on the devitrification of glasses has
shown that internal crystallization can sometimes be promoted if small
amounts of nucleating agents are incorporated into the glass.

Pavlushkin et a1. 1' z, 3 have investigated the effect of various nucle-

 

ating agents, e. g. F, T102, and P205, on the devitrification of lead
silicate glasses with high lead oxide content, and in 1966 reported that
P205 was the most effective catalyst for promoting internal crystal-
lization in this system. According to their observations, crystallization
proceeds through the following sequence: during cooling, phase separa-
tion occurs followed by the formation of crystalline nuclei which grow

during subsequent heat treatment. In support of this mechanism is the

work of Vogel “ who reported that glasses containing 60 mole percent
PbO and 40 mole percent 8102 phase separate. In 1969, however,

Shaw and Uhlmann 5

predicted, on the basis of their work on the
density versus composition curves, that the most probable compo-
sition range for immiscibility in the lead silicate system was in the
region from 0 to 50 mole percent PbO. The glasses studied by both
Pavlushkin 1:2: 3 and Vogel “ were outside this composition range.

The purpose of this research was to determine the effect of
P205 on the devitrification of lead silicate glasses of high PbO con-
tent. Two glass compositions were studied, one containing 59 mole
percent PhD and the other containing 64 mole percent PbO. Inves-
tigations were aimed at determining the effect of P205 on

(1) the rate of crystal growth,

(2) the resultant crystal phases,

(3) the microstructure of the devitrified glasses, and

(4) the crystal morphology.

C HA PT ER II
THEORY

The glassy state is a solid state characterized by the absence
of molecular arrays with long range order. With many materials,
it is difficult to achieve high enough cooling rates to prevent molecules
from rearranging themselves into regular arrays with long range
order. These materials crystallize upon solidification. With other
materials, however, the time required for molecular rearrangement
is long enough that it is possible to cool the system from the liquid
state without crystallization occurring. Although many materials
can be formed as glasses with rapid enough cooling rates, the phrase
glass-forming materials is generally reserved for those materials
which form glasses in bulk form when cooled at moderate rates from
the liquid state. Thus, the tendency of a melt to form a glass upon
cooling depends upon the interplay between (1) the rate at which it is
cooled and (2) the kinetics of the crystallization processes involved.
Materials with slow crystallization rates tend to form glasses easily.

The vitreous state is a metastable, not an equilibrium, state
from which the materials would like to escape, via devitrification.
If a sufficient amount of energy is available, the vitreous material
can devitrify and transform into a crystalline material. The crystal-
line products formed are usually the equilibrium phases for the sys-
tem, as determined by the phase diagram for the system.

Crystallization of a supercooled liquid usually starts with the
formation of nuclei. It is, therefore, customary to view crystal-

lization as a two-step process involving nucleation and growth.

 

Formation of nuclei occurs during the nucleation stage. If the nuclei

form spontaneously from the transforming phase, unaided by the

3

4

presence of foreign particles, then the process is termed homogeneous

 

nucleation. If, on the other hand, nuclei formation is initiated by the

 

presence of foreign particles, then the process is termed hetero-

geneous nucleation. The foreign particles may be anything present

 

in the parent phase that causes it to be non-homogeneous, e. g. im-
purities such as dust or dirt, air bubbles, crystals of another phase,
or even crystals of the same phase if they 'are added to the parent
phase (instead of being nucleated from it). Crystal growth proceeds
from these nuclei during the growth stage. The rate at which crys-
tallization occurs is a function of both the nucleation and growth rates
and is limited by the slower of the two processes. Either a low
nucleation rate or a low growth rate can prevent crystallization from
occurring on a reasonable time scale.

In studying the solubilities of materials, Ostwald 6 observed
that some degree of undercooling from the equilibrium melt temper-
ature always preceeded crystallization. Homogeneous crystallization
seemed to be dependent upon the liquid's reaching some degree of
supersaturation before crystallization could take place. From his
observations, Ostwald postulated the existence of a metastable zone
of supersaturation within which crystal nuclei could not form spon-
taneously. Experimental studies since then have been carried out for
a wide variety of materials, and results indicate that the "lowest tem-
perature to which liquid droplets can be cooled without crystallizing,
and 0. 85 T , " 7 where

E E

TE is the equilibrium melt temperature. Tammann, in investigating

crystallization in undercooled organic liquids and inorganic glasses,

for nearly all materials lies between 0. 75 T

found that "melts which increase rapidly in viscosity during cooling
and attain a glassy state, have maxima in their nucleation and growth

" 6 as shown in Figure 1. Both the nucleation rate and the

rates,
growth rate are limited at each extreme of undercooling by the exis-
tence of metastable zones. ”At small degrees of undercooling, there

is a metastable zone in which nuclei do not form at any measurable

Metastable zone High-viscosity
metastable zone‘

‘\

\\\t\\?\\§\
\s
\\ \
\

\\

I Crystal growth rate

\

/ Spontaneous
nucleation
rate

\

Rate ——~>
\\

\ .-\.\-\
\\\. \‘
\\;\ \

 

, \‘.
.\\‘\\ \ \
~‘4'\\ \

Degree of undercooling I

Melting
temperature of viscous liquid

Rates of nucleation and growth for viscous liquids as a

Figure 1.
function of the degree of undercooling. "

6

6 At larger

rate, but in which crystals, , once nucleated, can grow. "
degrees of undercooling there is another metastable zone, in which
high viscosity prevents crystallization. The maximum in the growth
rate curve was found to occur at smaller undercoolings than the max-
imum in the nucleation rate curve. The implication is that upon cool-
ing, glass-forming melts tend not to crystallize since, owing to the
relative positions of the nucleation and growth rate curves, an appre-
ciable number of nuclei are not formed until after the melt has cooled
through the region of rapid growth. In this case, crystallization would
seem to be limited by the lack of nuclei. And indeed, it has been ex-
perimentally verified that many glass-forming materials that fail to
crystallize when cooled directly from the melt temperature to a growth
temperature, Tg' will crystallize extensively provided that the melt

is first cooled to the glassy state and then reheated to Tg’ Such ob-
servations lend support to the belief that it is not a slow growth rate,
but rather a lack of nucleation sites that limits the crystallization of
glass-forming melts during cooling.) Further support is found in the
observation that crystallization in glasses is often found to nucleate
from external surfaces, bubbles, and other sources of impurities.
Some investigators. ' however, believe that there are always suffi-
cient nuclei present to initiate crystallization and attribute glass-form-
ing ability to slow growth rates. This dispute has not been satisfac-
torily resolved, and indeed may persist for some time. The major
source of uncertainty is that when crystallization occurs in glasses, it
is difficult to determine whether it has occurred through homogeneous
nucleation processes, or whether some impurity has nucleated the

process.

2. 1 Classical Nucleation Theory
Mathematical equations describing the rate of homogeneous nucle-
ation in glass-forming systems have evolved as extensions of classical

nucleation theory; Developed to describe condensation in a

7

9' 10 is built upon

supersaturated vapor, classical nucleation theory
the premise that Gibbs free energy is the driving force for conden-
sation. The problem is formulated by considering the effect that
droplet condensation has on the free energy of the system. Conden-
sation of a spherical droplet of radius r leads to a change in the sys-
tem's free energy, AF, with

4 3 2
3- TYr f + 4TTr fs , (1)

AF =
where fV is the difference in free energy per unit volume between the
vapor and the liquid states, and fs is the free energy per unit surface
area (associated with the formation of the vapor-liquid interface).

For a supersaturated vapor, fV is negative and the variation of

with r will be as shown in Figure 2. As r increases, AF first in-
creases, and then goes through a maximum and decreases. The max-
imum value of AF, denoted AFC, occurs when the droplet radius

reaches a critical size rc, where

 

-2 fS
rc = . (2)
f
v
so that
16 in:
AF = “—3 . (3)
° 3f
v

Droplets with radii smaller than rc are termed "embryos" and tend to
evaporate, whereas those with radii larger than rc are termed "nuclei"
and tend to grow. Droplets of critical size are unstable, and may grow
or shrink, since either process leads to a decrease in AF. The rate

of nucleation, I, is given by

I = Kl exp(-AFc/kT) , (4)

AF

 

 

I
.’
'c \ r

Figure 2. Variation of Gibbs Free Energy with radius for condensation
of a spherical droplet in a supersaturated vapor.

where k is Boltzmann's constant, T is absolute temperature, and
K1 is proportional to the probability of a vapor atom's colliding with
a critical-size nucleus. Several different expressions have been
derived for K1, which vary depending upon the assumptions made re-
garding the exact form of the distribution function for embryos in the
system. The value of I is dominated by the exponential term, however,
so that K1 is often approximated as the collision frequency of vapor
atoms in the system.

The rate of homogeneous nucleation in condensed systems was

shown by Becker 1° to be governed by an equation of the form
I = K2 exp(- AFC/kT) , (5)

provided that K2 is interpreted as the encounter rate of atoms and
AFc is interpreted as the free energy associated with a solid embryo
of critical size. In a condensed system, K2 is given by

K2 = A exp(-Q/k'r) , (6)

where A is proportional to the atomic vibrational frequency and Q is
the activation energy for diffusion, which in glass-forming systems
is often approximated as the activation energy for viscous flow. Upon

substituting for K2, equation 5 can be rewritten as
I = A expE-(AFc + Q)/kT] . (7)

In glass-forming systems, both AFc and Q are temperature dependent;
however, AFc decreases with decreasing temperature, whereas Q
increases with decreasing temperature. Thus, the maximum in the
nucleation rate, observed experimentally by Tammann for inorganic
glass-forming materials, arises in equation 7 from the interplay
between AFc and Q with temperature change. Equation 7 may also be
used for heterogeneous nucleation, where the presence of impurities

effectively lowers the interfacial energy between the nuclei and the

 

 

Ill}

 

i

 

10

parent phase, provided that a factor is included in equations 1 and 3

to reduce the interfacial energy term fs.

2. 2 Growth Processes

 

Once formed, nuclei grow at the expense of the parent phase
(matrix), securing material of the appropriate composition and ar-
ranging this material to fit into the crystalline structure. The rate
at which crystalline nuclei grow can be limited by (l) the rate at which
material can be transported to the crystal-matrix interface, termed

diffusion controlled growth, and/or (2) the rate at which rearrangement

 

at the interface occurs, termed interface controlled growth.

 

When growth is interface controlled, e. g. in the solidification
of a pure melt, a linear dimension of the growing crystal is prOpor-
tional to time, thus making the growth rate time independent. The
growth rate, G, is a function of undercooling, the exact form of which
depends upon the assumed mechanism of crystal growth. Essentially
three mechanisms for crystal growth have been proposed.

(1) Growth by two- dimensional nucleation9-- developed by Gibbs,
Becker and Doring, and Frenkel. Crystal growth is assumed to occur
through the repeated nucleation and subsequent lateral growth of
"island" clusters on close—packed surfaces. Once an island nucleus
is formed, it grows laterally by the addition of atoms at the preferen-
tial sites, called "steps", which bound the island until the entire crys-
tal surface is covered with another close-packed layer. Further
growth cannot occur until another stable island nucleus is formed.
Thus, the crystal growth rate is a function of the rate at which island

nuclei are formed, and as a result should be proportional to

exp (-B/ AT)
Tl

where n is the viscosity, AT is the degree of undercooling, and B is a

constant.

ll

(2) Growth at repeatable steps 9 -- proposed by Frank. This
mechanism assumes that crystal growth occurs through the addition
of atoms to sites of crystal imperfections, repeatable steps, which
are self-perpetuating, e. g. screw dislocations. Since sites are
always available for crystals containing these imperfections, close-
packed layers are not formed, and the formation of island nuclei,
required for the two-dimensional mechanism, is unnecessary. Under

these conditions, the crystal growth rate will be proportional to

(AT)
Tl

 

(3) Growth without steps 9 -~ prOposed by Cahn. This mecha-
nism proposes that under certain circumstances (e. g. at high under-
coolings) atoms may attach themselves to any position on the crys—
tal interface so that the interface advances uniformly, without steps,

at a rate prOportional to

AT

TI

In diffusion controlled growth the crystal growth rate is limited
by the rate at which material can be transported to and from the

growth front. The growth rate, G, is given by
G = a D/y , (8)

where D is the diffusion coefficient, y is the effective distance over
which atoms migrate, and a is a term dependent upon the compositions
of the crystal, the matrix, and the interface. In most cases of dif-
fusion controlled growth, a concentration gradient builds up ahead of
the growth interface with time, resulting in an increase in the effec-

tive distance over which atoms must migrate, so that

12

(1301/2 . (9)

y

where t is time. Upon substituting for y, equation 8 becomes

G = 01(D/t)l/2 . (10)

Under some circumstances, e. g. the growth of a platelet in its own
plane or the lengthening of a needle- shaped crystal, crystal growth
can proceed without an increase in y. For these cases, a linear
dimension of the growing crystal will be prOportional to time, and

the growth rate will have no time- dependence.

2. 3 Crystallization in Glass-Forming Systems

 

Investigations of crystal growth rates as a function of under-
cooling in several glass-forming systems suggest that the mechanism
of crystal growth may well vary from one system to another. H. R.
Swift 11 showed that the growth rates varied linearly with AT/TI in
several soda-lime glasses. A similar dependence of crystal-growth
rate on undercooling has since been observed for several simple
oxide glasses, ‘2 e. g. 8102, GeOZ, and P205. In the sodium-silicate
system, however, experiments by Scott and Pask ‘3 indicated that
the growth rate varied linearly with ATL 75/1], and in the lithium-
silicate system, experiments by Morley “ showed that with the appro-
priate choice of constants the experimental growth rate versus tem-
perature curves could be fitted by either a ATL 75l“ or an exponen-
tial relationship.

The tendency of a glass towards devitrification can often be mod-
ified by nucleating agents. Such agents can change the rate of devit-
rification, the microstructure of the crystalline product, and to some
extent, the composition of the crystalline phases. Numerous additives,
including noble metals, oxides, fluorides, sulfides, transition-group

elements, and halogens, have been found to promote internal crystal-

lization in glasses. The effectiveness of a particular nucleant depends

13

upon the glass system being nucleated and the concentration in which
the nucleant is added. Once a nucleating agent is added to a glass,
crystallization is usually brought about by a two- stage heat treatment
process in which the glass is first heated at a low temperature to
induce nuclei formation, and then at a higher temperature to induce
crystal growth.

The mechanism through which a nucleating agent acts varies
with both the nucleating agent and the glass system. Some of the
reaction paths along which crystallization may proceed in nucleated

glasses are shown in Figure 3. ‘5

Only two of the paths involve crys-
tallization of the nucleant prior to crystallization of the major phase.
Each of the four reaction paths is discussed in more detail below.

When devitrification occurs via path 6 - 5, crystallization of
the major phase is catalyzed by the crystallization of the nucleating
agent. Most of the experimental evidence for this sequence is found
with noble metal nucleants. Stookey ‘6 found that the addition of
small amounts of a noble metal to some lithia aluminosilicate glasses
produced photosensitive glasses. Photonucleation resulted in the
growth of metallic crystals which served as nucleation sites for the
crystallization of lithium metasilicate. This reaction path has not
been commonly observed for other nucleating agents.

Crystallization occurring along path 2 - 3 has been observed and
is the source of much debate. Glasses following this crystallization
sequence exhibit glass-in-glass phase separation prior to crystal-
lization of the major phase. Amorphous phase separation describes
the process whereby a glass separates into two amorphous phases,
with droplets of one phase dispersed in a matrix of the other phase.
This process is caused by the immiscibility of the two phases, and
hence is dependent upon composition and temperature. Only certain
glass systems undergo phase separation, and then only over a specific

range of composition and temperature. Nucleating agents have been

found to influence the ability of a glass to phase separate ‘2 --

 

 

i (e. n C» a

 

 

14

m—

.oHEwaoonmmmHm Re: 65 new mmflm 13:5 65 come-«on mfima cofiomom .m oasmrm

 

:mmém BZ<HAUDZ= .mO
mJ<HmWflU A<QHOAJOU

a

 

ZOHB<m<nHmnm

1' magma @0332 .mO
ZOHE<NHAJ<Em>m0

a

 

mm<AU

 

mm<mm mDOEnHmOSSwAP

mDOflZMUOEOfi

15

sometimes suppressing phase separation, e. g. A1203 in sodium
borosilicate glasses, and sometimes enhancing phase separation,
e. g. P205 in sodium silicate glass. Whether phase separation is
always a precursor to crystallization has been debated. It seems
likely that phase separation may often be associated with, but not
required for, subsequent crystallization.

Glasses crystallizing along path 2 - 4 - 5 undergo phase sepa-
ration followed by crystallization of the nucleating agent, which in
turn is followed by crystallization of the glass. Such a sequence
was observed for a magnesium aluminosilicate glass containing
titania ,16 where phase separation was followed first by crystallization
of magnesium titanate and then by crystallization of the major phase.
Again, no cause- effect relationship has been established between
phase separation and the subsequent crystallization of the major
phase. Results of one study suggested "that the colloidal crystals
of the nucleant phase which occur after the amorphous phase sepa-
ration and before crystallization of the major phase are not neces-
sarily an integral part of the nucleation of the major phase". 17

Reaction path 1, in which a homogeneous glass crystallizes
directly into its major phases without phase separation or crystal-
lization of a nucleant phase, has not been observed as a common
' occurrence, although such a course was followed by TiOz nucleated

lithia aluminosilicate glass, according to Barry et a1. . ‘3: ‘9

 

CHAPTER III
THE LEAD SILICAT E SYSTEM

There are two naturally occurring lead silicate compounds - the
minerals Alamosite (PbO-SiOZ) and Barysilite (3PbOoZSiOZ). In addi-
tion, several lead silicate compounds have been synthesized by the
devitrification or by the sintering of PbO and SiOZ. For many years
the accepted phase diagram of the lead silicate system was that of
Geller, Creamer, and Bunting, 2° which is shown in Figure 4. Three
stable lead silicate compounds with PbO:SiO2 ratios of 4:1, 2:1, and
1:1 are shown. The 4PbO'SiO2 compound has three polymorphic mod-
ifications, " and the PbO:SiO2 compound corresponds to the mineral
Alamosite. The work of other investigators, however, has indicated
the existence of additional compounds, some of which may be meta-
stable. Since devitrification in the lead silicate system may be com-
plicated by the existence of metastable phases, this chapter will first
review the various investigations that have been made of phase rela-
tions in the lead silicate system. Then, the results of previous

kinetic studies of lead silicate compounds will be presented.

3. 1 Compound Formation

 

In 1909 Cooper, Shaw, and Loomis 3‘ investigated the thermal
properties of the lead silicate system. Heating curves were deter-
mined for a number of crystalline samples, which had been prepared
from slowly cooled melts whose compositions ranged from 43 to 100
mole percent PbO. Melting points of the various compositions were
determined from breaks in their heating curves and used to construct,
melting point curves. On the basis of the melting point data, the
orthosilicate and metasilicate compositions, 2PbOvSiOz and PbO-SiOZ

l6

l7

 

   

 

 

 

 

 

 

 

 

 

 

1 F T I f T
.- Tndymilo
0L .'
900 1—
Liquid
(L)
800 2:: 1:: Quart—
. imsso, a!“
«1+1. {area-550. ' IL
' . , ‘1 l "‘
'3‘) A?” (31"- :, v -
700 -—
PhD M ' ZPDOSIO' PbO;SiO'
P004510] shawl, homo,
I
q
1 :::bopboglsmr I (T! l v 1 I J
30
“Peight % $182 Si0.""

Figure 4. Phase diagram of the lead silicate system by Geller et al.z.°
The actual composition of the 70-30 base glass studied in this inves-
tigation is indicated by the solid triangle along the base line, that of
the 60-40 base glass with the open triangle.

18

respectively, were established as lead silicate compounds. Well
defined maxima in the curve corresponded to the melting points of
these two compounds. Two eutectic points were found over the com-
position range studied: one occurred between the lead oxide and the
orthosilicate compositions, while the other occurred between the
orthosilicate and the metasilicate compositions.

The following year Hilpert and Nacken 7'2 reported that their
work indicated the existence of the orthosilicate and metasilicate com-
pounds and the possible existence of compounds with compositions
3PbO-SiOZ and 3PbO-ZSiOZ.

Investigations of the lead silicate system begun by Cooper, Shaw,
and Loomis 7" were continued by Cooper, Kraus, and Klein 7‘3 and ex-
tended to cover optical, as well as thermal, properties. Their study
included observations of both natural minerals Alamosite and Bary-
silite and synthetic lead silicate crystals prepared by the slow cool-
ing of lead oxide-silica melts. Results of optical and thermal analy-
ses led the authors to conclude that the compositions PbO-SiOz,
ZPbO-SiOz, and 3PbO-ZSiOz existed as compounds in the lead silicate
system. Properties of the PbO-SiO2 and the 3PbO-ZSiOZ compounds
prepared in the laboratory were found to correspond to those of natu-
rally occurring Alamosite and Barysilite. The existence of a com-
pound with the composition 3PbO-SiOZ was deemed probable on the
basis of its sharp melting point and its somewhat anomalous optical
characteristics.

In 1931 Krakau and Vakhromeev 2“ reported results of their
thermal and Optical investigations of the lead silicate system, and in
basic agreement with previous studies, they concluded that four com-
pounds existed: 3PbOoSiOz, ZPbO-SiOZ, 3PbOoZSiOZ, and PbO-SiOZ.
In addition, a polymorphic transformation at 62 0°C was found in the
orthosilicate compound.

A few years later Geller et a1. 2° conducted their investigation

 

which indicated the existence of three stable compounds in the lead

l9

silicate system, namely 4PbO-SiOZ, 2PbO-SiOZ, and PbO-SiOz.
Identification of the compounds with compositions 2PbO-SiOZ and
PbO-SiOZ was in close agreement with earlier reports, as were the
values determined for their melting points. The 4PbO°SiOZ compound,
however, had not previously been reported. More important was the
omission of compounds with compositions of 3PbO-SiO2 and

3PbO-ZSiOZ from the phase diagram presented by Geller et al. .

 

Attempts by Geller _e_t_a_l. to synthesize either compound were unsuc-
cessful. Mixtures containing 20 mole percent SiO2 crystallized as
a single phase, which formed PbO and glass upon heating to 725°C -
735°C. The authors suggested that their 4PbO SiOz compound was the
3PbO'SiO‘2 compound reported by Krakau and Vakhrameev. Mixtures
of the 3PbO'ZSiO2 composition were found to crystallize into a com-
bination of orthosilicate and metasilicate after a heat treatment of
20 hours at 705°C; while "specimens quenched after holding 3 hours
or less at the same temperature consisted almost wholly of a phase
resembling Barysilite in color and index of refraction, but which was
either all glass or glass containing extremely minute crystals. " 2°

The existence of the compounds 2PbO'SiOz and PbO'SiO2 was
verified with x-ray diffraction in 1935 by Valenkov and Porai-Koshits.Z4
They concluded that the compounds 3PbO-SiO2 and 3PbO-ZSiOZ, as
well as the 4PbO'SiOz compound reported by Geller gt_a_l_. .20 did not
exist as stable phases in the lead silicate system. The 4PbO'SiOz
composition yielded mixtures of BPbO and 2PbO'SiOz. ‘5

In 1939 an x-ray diffraction investigation was conducted by
McMurdie and Bunting 7'5 to verify the existence of the compounds
that had been reported earlier by Geller £31. . 2° Their results in-
dicated that the compounds 4PbO-SiOz, 2PbO-SiOZ, and PbO-SiOZ did
exist in the lead silicate system. Further, they reported that the
diffraction patterns obtained from the 4PbO-SiO2 compound could not
have arisen from a mixture of BPbO and ZPbO'SiOZ, as had been sug-

gested by Valenkov and Porai-Koshits.“4

20

The existence of 2PbO-SiO2 and PbO-SiO
24

2 was verified again in

1956 by Milent'eva and Solov'eva with thermographic and micro-

scopic analyses. In 1960 Argyle and Hummel 7"

made an x-ray and
dilatometric study of synthetic lead silicate compounds. They, too,
were able to verify the existence of the 2PbO°SiOz and PbO'SiOZ com-
pounds. No polymorphic transformation was reported for the 2PbO-
SiOz compound in as much as this phase exhibited no discontinuities
in the thermal- expansion curve. This result was in agreement with
the findings of Geller and co-workers. Argyle and Hummel also veri-
fied the existence of the 4PbO'SiO2 compound, including the trans-
formation of y-4PbO-SiOZ to B-4PbO:SiOZ. The B toa inversion was
not observed; the authors felt, however, that this inversion would
have been difficult to detect by thermal-expansion methods. Attempts
to prepare the 3PbO-SiOz and the 3PbO-2SiO2 compounds by high-tem-
perature sintering and devitrification techniques failed, as mixtures
of 4PbOoSiO2 and 2PbO»SiOZ were obtained for the 3PbO-SiOz compo-
sition, and mixtures of 2PbO°SiOZ and PbO°SiOZ were obtained for the
3PbO-ZSiOZ composition. These investigators concluded that these
two compositions probably did not exist under normal conditions of
temperature and pressure.

Work on the lead silicate system was taken up again a few years
later by Berezkina and Chizhikov 2‘ with an x-ray diffraction study
aimed at determining the compositions of the compounds present in the
system. Crystalline samples were obtained by devitrification of melts
of lead oxide and silica. Their work confirmed the existence of 4PbO -
SiOz, 2PbO'SiOz, and PbO°SiOZ. Two polymorphic modifications of

lead orthosilicate were found. Furthermore, 3PbO-ZSiO was report-

ed to occur as a distinct compound which existed, howevcir, only over

a narrow temperature range near the melting point. At lower temper-
atures, the compound became unstable and decomposed into a mixture
of 2PbO 'SiOz and PbO'SiOZ. The 3PbO'SiOZ composition was found to

crystallize as a mixture of 4PbO-SiO2 and 2PbO'SiOZ.

21

Smirnova,27' 28' 29 in the late 1960's, made several studies of
crystallization in lead silicate glasses utilizing infrared spectroscopy,
electron microscopy, and x-ray diffraction techniques. Results in-
dicated that there were at least two polymorphic modifications for
both the 2PbO'SiOZ and the PbO'SiO2 compounds. The transition for
the orthosilicate occurred between 450°C and 620°C. The modifica-
tions for the PbO'SiOZ compound arose from differences between sur-
face crystals and internal crystals.

In 1967 Pavlushkin and Lisovskaya 1 reported that crystalliza-
tion of lead silicate glass of composition close to 60 mole percent
PbO resulted in mostly ZPbO'SiOZ. A year later they reported 7'
that crystallization of a glass containing 60 mole percent PbO at
600°C produced crystals of the following compounds: 2PbO'SiOZ,
4PbO-SiOZ, and 3PbOoZSiOZ. A more recent study, 3 published in
1971 concluded that the same glass crystallized into a combination of
2PbO'SiO2 and 4PbO-Si02 when heat treated at temperatures of 450°C
and 550°C.

In the various studies made on the lead silicate system, only the
existence of the compounds ZPbO'SiOZ and PbO'SiO2 was agreed upon.
Controversy surrounded the existence of compounds with compositions
of 4PbO'SiOz, 3PbO'SiOZ, and 3PbO'ZSiOz. In addition there was no
agreement as to the existence of polymorphic modifications for the
2PbO'SiOZ and PbO'SiO2 compounds. In an effort to resolve this con-
troversy, Ott and McLaren 3° conducted a DTA and x-ray diffraction
investigation of the system. The resulting phase diagram appeared in
1970, and is given in Figure 5. The region above 700°C was based on
the phase diagram of Geller 33in" The authors cautioned, however,
that since crystallization occurred very slowly, their proposed phase
diagram might not be a true equilibrium diagram due to the possible
inclusion of metastable phases. X-ray diffraction analysis indicated
the existence of compounds 4PbO'SiOz, 3PbO'SiOZ, 2PbO'SiOZ, 3PbO'
ZSiOZ, and PbO'SiOZ. The three modifications of 4PbO°SiOZ, reported

 

22

 

 

 

 

 

 

 

 

 

 

 

.— a- a- N a-

v ("I a .7. .'.'.

| I I l I
800' '20 .

um
700 O O O O . . O l O
""“M l=|+H§iOz
":0 My .
600 l' M12” 0 H2” 0 0 0 o l 0
H3” ””2”” 1:1-+1550:
(l 3:2 .
500 ° 0.. 0 0 8 TI c- 0 -
M4:l+12:l"
i ll 3:2 ‘9' 0.5502
t12:1
3,”__... 12:! f
400 M4” 3+” 3.2
a A a A 1 . +
to Is 20 as so 35
wucur PEI cem smca
Figure 5. Phase diagram of the lead silicate system by Ott and
McLaren. 3° The solid triangle along the base line indicates the

actual composition of the 70- 30 base glass studied in this inves-
tigation, the open triangle the actual composition of the 60- 40
base glass.

23

by Geller et a1. ,7‘° were confirmed, as were the two modifications of

ZPbO'SiOZ, as reported earlier by Krakau and Vakhrameev, 7‘4

Berezkina and Chizhidov, 7'4 and Smirnova et al. .27: 2°: 7'9 The 3PbO'

 

SiOZ and 3PbO-ZSiOz compounds were found only at low temperatures
and were determined to be unstable above 430°C and 585°C respec-
tively. In addition, Ott and McLaren found that the PbO'SiOZ com-
pound was unstable below 525°C, decomposing into a mixture of 3PbO'
ZSiOZ and 8102. Ott and McLaren's work somewhat clarified the con-
troversy surrounding the lead silicate system.

A study of compound formation in the lead silicate system by

Smart and Glasser 3‘

also helped to resolve some of the apparent con-
tradictions of earlier investigations. The resulting phase diagram,
which is presented in Figure 6, is quite similar to that of Geller

et al. . 2° One major difference between the two diagrams is the in-

clusion by Smart and Glasser of a stability field for the 3PbO'SiOZ com-

 

pound. Another difference was the inclusion of SPbO'BSiO2 as a pos-
sible stable compound. In all, fifteen lead silicate compounds were
prepared by Smart and Glasser, including two modifications of the
3PbO°ZSiO2 compound. Many of these compounds, however, were
determined to be metastable phases and hence were excluded from the
proposed phase diagram. The authors suggested that the metastable
phases arise as intermediate crystallization products, which even-

tually convert into stable crystalline phases.

3. 2 Rate of Crystallization

 

The crystallization velocity of lead silicates has been studied by
Hilpert and Nacken 22 and by Rita and Bergeron. 37‘ In 1910, Hilpert
and Nacken ‘2 reported results of their study on the crystallization of
2PbO'SiOZ and PbO°SiOZ. Both compounds reached maxima in their
growth rates at approximately 690°C. The maximum growth rate of
120 mm/ hr observed for the orthosilicate compound, ZPbO'SiOZ, was

substantially higher than the 20 mm/ hr maximum observed for the

 

 

 

 

 

 

90

KJ

Figur e (
GlaSSer
glaSSes

triangle

24

 

I

    
    

LIQUID

 

 

 

 

 

 

 

 

 

800 "
739:2‘
ALAMOSITE . HIGH ouaarz
700 d t . ":2. . )— — -1 — 3.790.— -------
4:1ouz:1 "‘ ”190‘ |
.5. ‘L . | HIGH QUARTZ .
._°_‘.‘£=__( .. "955'. 0,,
H 96,510. 0" '
. f7 AMOSITE A 573"“
96,90, 4:1 2 ”- go. If
(71 I
o I” II
a 5 2 l LOW ouamz .
500 "‘ 2 . a g ' ”550.02.
a j I
RED no 8 < :
PD‘SIO. |
L Q I! v A
r 1 :1 1‘ g. I 510
P130 20 4O 60 80 2

MOLE % SIC2

Figure 6. Phase diagram of the lead silicate system by Smart and
Glasser. 3‘ The actual compositions of the 70-30 and 60-40 base
glasses studied in this investigation are indicated by solid and open
triangles, respectively, that lie along the base line.

25

metasilicate compound, PbO'SiOZ. Both compounds also exhibited
regions of spontaneous crystallization -- ranging from 630°C to 400°C
for ZPbO’SiOZ and from 670°C to 470°C for PbO-SiOZ -- over which
rapid formation of nuclei interfered with growth rate measurements.
Below this region the growth rates of both compounds dropped off
significantly. For the orthosilicate compound the crystal growth rate
was reduced to about 1 mm/hr, and that of the metasilicate compound
was too small to measure.

Rita and Bergeron 37‘ published results of their study of the
crystallization of 2PbO'SiOz in 1976. Crystal growth was measured
over the temperature range of 500°C to 740°C with a maximum growth
rate of 28 um/sec occurring at approximately 690°C. According to
the authors, the growth rate data did not appear to fit any of the stan-

dard crystal growth models.

CHAPTER IV
EXPERIMENTAL PROCEDURE

4. 1 Glass Preparation

 

All glasses in this study were prepared from analytical reagent
grade PbO and 30 um Min-U-Sil, a high-purity form of silica. Phos-
phorus pentoxide was added in the form of ammonium dihydrogen
phosphate. Two nominal compositions of base glasses were prepared --
one containing 70 mole percent PbO and 30 mole percent SiOZ, and the
other containing 60 mole percent PbO and 40 mole percent 8102, here-
after referred to as 70-30 and 60-40 respectively. Ammonium di-
hydrogen phosphate was added to the 60-40 base glass to give glasses
with 0.1 and 0. 5 mole percent P205, designated as 60-40-0.l and
60-40-0. 5 respectively, and to the 70- 30 base glass to give glasses
with 0. 5 and l. 0 mole percent P205, designated as 70- 30-0. 5 and
70-30-1. 0.

After careful consideration, 900°C was selected as the melting
temperature for glass formation. This temperature was considered
high enough to insure complete melting of the batch materials, since
it is above the melting points of both PbO and the glasses being formed,
and yet low enough to avoid drastically altering the final composition
of the glasses through excessive volatilization.

The high lead oxide content of the glasses made the melting pro-
cedure a bit difficult. The melts were extremely corrosive, especially
at the high melt temperature used, and readily attacked the crucibles
in which they were being melted. Several batches were lost when the
melts dissolved the crucibles. Alumina, fused silica, and Vycor
crucibles were tried with varying degrees of success. Alumina cru-
cibles held up the best. The thick walled Norton alumina crucibles
(70318) were capable of holding the melt up to 4 or 5 hours. The Vycor

26

27

crucibles usually held melts for l or 2 hours, and the fused silica
crucibles dissolved rapidly, usually in less than 1 hour. The alumina
crucibles were, in spite of their superior performance, not ultimately
used for glass formation since any alumina dissolved from the crucible
by the melt would go into solution and remain as an impurity in the
glass. Since the melts were so corrosive, there was a strong pos-
sibility that the amount of alumina picked up by the glass during the
melting operation would be even larger than the amount of P205 being
added as a nucleating agent. This was later confirmed by microprobe
analysis which showed that glasses melted in alumina crucibles con-
tained as much as 2 percent aluminum. Vycor, on the other hand, is
made from sodium borosilicate glass from which the sodium borate
phase has been leached out, leaving behind a network of 96% $102.
Batch material melted in Vycor crucibles would probably pick up
some additional 8102; however, no significant amount of impurity
would be added to the glasses being formed. For this reason, Vycor
crucibles were selected for all batch melting.

The corrosive nature of the melts made it necessary to keep the
time required for adding batch material to the crucible as short as
possible. Therefore, all glasses were made in a Lindberg Hevi- Duty
furnace which was capable of reaching the melt temperature very
rapidly, usually within 30 minutes. During this time, frequent addi-
tions of batch material were made to the crucible so that when the fur-
nace reached 900°C, all batch material was added and molten. The
melt was then held at 900°C for 1 hour. This procedure worked rea-
sonably well, and only occasionally were melts lost due to crucible
failure. The complete procedure followed in making the glass samples
is listed below.

1. PhD and SiO were weighed out and thoroughly mixed.
Glasses were made in 150 gm batches.

2. Vycor crucibles were placed in the furnace and the batch
materials were added as the furnace was heating.

28

3. Melts were held at 900°C for 1 hour.

4. At the end of 1 hour, the melts were quenched in water,
yielding fine, glassy pellets of the base compositions.

5. The glass pellets were crushed and sifted through a 400
mesh sieve.

6. The sifted glass powder was weighed out in 150 gm batches.

7. Ammonium dihydrogen phosphate was weighed out and added
to the powdered base glass to form the P205 nucleated

glasses.

8. All batches were remelted in new crucibles according to
steps 2 and 3.

9. Finally, the melts were quenched between steel plates, pro-
ducing glass plates, approximately 1/8 in. thick.

4. 2 Devitrification of Glasses

 

The crystallization of each glass was studied at temperatures of
400°C, 450°C, 500°C, and 550°C. Samples were placed in graphite
crucibles and inserted into thermally equilibrated furnaces. The heat
treatment time was measured from the time of the sample's insertion
to the time of its removal from the furnace. During heat treatment,
the furnace temperature was monitored with a calibrated thermo-
couple to determine: (1) the length of time required for the furnace
to re- equilibrate after insertion of the sample, and (2) the maximum
fluctuation in temperature that occurred during the heat treatment.
The furnace usually regained crystallization temperature within 3 to
4 minutes. Heat treatment times ranged from a few hours for crystal-
lization at 450°C to several days for crystallization at 400°C. For
both of these cases, the time required for the furnace to regain equi-
librium after sample insertion was considered negligible with respect
to the long heat treatment times involved. Crystallization heat treat-
ment times at 500°C ranged from 10 minutes to 140 minutes, and
the furnace recovery time was again neglected. For crystallization
at 550°C, however, heat treatment times were so short, ranging from

3 to 10 minutes for the 70-30-1. O composition, that a way had to be

29

devised for inserting the glass samples without lowering the tempera-
ture of the furnace. This was accomplished by using a furnace with a
small front window and equilibrating it with the window open. The
open front window had no noticeable effect on the stability of the fur-
nace temperature. Glass samples to be heat treated were placed in

a small basket that had been formed at the end of a long metal rod.
With the basket, it was easy to insert and remove the sample through
the window without disrupting the equilibrium of the furnace. The
maximum temperature fluctuation observed during heat treatment at
any temperature, after the furnace re-equilibrated, was less than

8°C.

4. 3 Methods of Observation and Analysis

 

The size and morphology of crystals in devitrified samples were
studied by light microscopy, utilizing both the reflected and transmitted
light modes. Thin sections of samples were prepared for study under
transmitted light, while polished faces, which had been lightly etched
with 1% HNO3, were used for reflected light microscopy. The diam-
eters of the crystals in each sample were measured from micrographs
taken at random intervals over the sample. Using this data, histo-
grams showing the size distribution of the crystals in each sample
were constructed. Then the mean radius, and standard deviation,
of each crystal population were calculated and plotted versus growth
time. The crystal growth rate was determined as the slope of the
best fit straight line through these points. Data points were subjected
to t- analysis and were rejected only if they lay outside of 95% con-
fidence limits.

Transmission electron microscopy was used to examine the
ultrastructure of both the quenched and the devitrified glasses. Sam-
ples~ were studied by replication and by direct observation of thin
sections. The thin sections were obtained by microtoming samples

on a Porter-Bluhm II microtome, equipped with a diamond knife.

30

Figure 7 shows the instrument set-up. The sample is drawn over the
stationary knife edge at a constant, predetermined velocity. After
each pass, the microtome automatically retracts the sample, swings
it back up to the top of the cycle, and advances the sample a set
distance, which becomes the thickness of the section cut on the next
pass. Cut sections float off the knife edge and into the boat where
they are later picked up onto grids. The samples were rather dif-
ficult to section, and often only shattered fragments of glass were
obtained. When intact thin sections were obtained, however, they
were uniformly thin and of much higher quality than thin foils pre-
pared by other methods, 9. g. chemical or mechanical thinning.
Samples were also studied by the replication method. Replicas
were made of (l) freshly fractured, (2) freshly fractured and then
etched, and (3) polished and then etched surfaces. Usually, however,
either fractured surfaces or fractured and then etched surfaces were
used as these seemed to produce the most consistent results. Most
replicas were prepared by the two- stage process in which the surface
features of the sample are impressed into replicating tape, which is
subsequently shadowed with a heavy metal and then backed with a layer
of carbon. The plastic tape is then dissolved, leaving a metal- shad-
owed carbon replica of the surface. Some difficulty was encountered
in dissolving the tape from the replica. Removing the plastic by
placing the plastic-backed replicas on grids which are then either
(I) placed on filter paper which is kept saturated with acetone until
the plastic is dissolved or (2) placed on fine mesh screens in con-
tainers which are kept filled with acetone to the level of the screen
until the plastic is dissolved, often resulted in badly torn, sometimes
completely destroyed, replicas due to the swelling of the tape which
occurred during the dissolving process. An alternative method, 33: 3‘
in which the replicas are washed in an atmosphere of refluxing acetone,
was found to produce cleaner, virtually intact, replicas and was adopt-

ed as the normal procedure for removing the plastic backing.

31

 

Figure 7. Microtome set—up for sectioning glass.

32

Occasionally single- stage metal- shadowed carbon replicas
were made. The heavy-metal shadowing and subsequent carbon back-
ing were deposited directly onto the glass surface, and then the replica
was removed by scoring the deposited film and floating pieces off in
a solution of nitric acid.

Electron and x-ray diffraction techniques were used to check
for crystallinity and to determine d- spacings for the crystalline com-
pounds formed during devitrification. Electron diffraction patterns
were obtained from thin sections and from extracted crystal fragments
which often adhered to the replicas, while x-ray diffraction spectra

were obtained from powdered, devitrified samples.

CHAPTER V
RESULTS AND DISCUSSION

5. 1 Crystallization Products

 

For all compositions studied, the melting procedure yielded
clear, yellow glasses. Chemical analysis indicated that the glasses
picked up silica during the melting operation. The 70- 30 base compo-
sition produced a glass containing 64 mole percent PbO and the 60-40
base composition produced a glass containing 59 mole percent PhD.
The actual compositions of both glasses are shown on the phase dia-
grams in Figures 4, 5, and 6. The 70-30 glass is close to the 3PbO-
2SiOZ eutectic, as shown in Figure 5, and is of the same composition

that was studied by Pavlushkin et al. ." 2'3 Figures 4 and 6 indicate

 

that the crystallization products in both the 70-30 and the 60-40 glasses
would be the same, namely 2PbO°SiO2 and PbO'SiOZ. However,
Figure 5 indicates that over the temperature range 400°C to 550°C, the
crystallization products of the 70- 30 glass would be 3PbO-ZSiO2 and
ZPbO-SiOZ, and those of the 60- 40 glass would be 3PbO°ZSiOz and
PbO-SiOZ or L-SiOZ.

X-ray diffraction spectra for the 70- 30 composition, after pro-
longed heat treatment at 400°C, 450°C, 500°C, and 550°C are shown
in Figure 8. Curve (a), corresponding to a heat treatment of 7 days
at 400°C, consists of only one broad peak, located between 25°<26< 33°
indicating that crystallization has not proceeded very far. Upon com-
paring the other three spectra for heat treatments at 450°C, 500°C, and
550°C, two trends can be seen.

(1) The peaks marked with solid arrows on curve (d) constitute

the major peaks for specimens heat treated at 550°C and 500°C and are

also found in the spectra of specimens devitrified at 450°C. These

33

34

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35

peaks occur with approximately the same relative intensities in all
three spectra and show a gradual sharpening and intensification with
increasing crystallization temperature. These peaks are probably
associated with the major crystalline phase, which apparently is
stable over the temperature range 450°C to 550°C.

(2) The peaks marked with open arrows on curve (b) decrease
in intensity as the heat treatment temperature increases, disappear-
ing completely at 500°C. X-ray spectra of glasses heat treated for
shorter periods of time at 500°C and 550°C also did not reveal these
peaks. These peaks most likely belong to a low-temperature, second-
ary phase which becomes unstable about 500°C.

The d-spacings and relative intensities for the major crystalline
phase are given in Table 1, while those of the low-temperature, sec-
ondary phase are presented in Table 2. The major crystalline phase
can be identified on the basis of its x- ray diffraction spectra as a poly-
morph of 3PbO'ZSiOZ. Although there are some inconsistencies in
the published x-ray diffraction data of several lead silicate compounds,
there is good agreement by three investigators 30:31:35

of this 3PbO-ZSiOZ modification. Both Smart and Glasser 3‘ and

on the spectra

Billhardt 35 identified a low-temperature modification of Barysilite
which formed readily upon devitrification of lead silicate glasses con-

taining 60 mole percent PbO. The x-ray diffraction data obtained by

31

Smart and Glasser for this compound, while not given, was report-

ed to agree with that of Billhardt.” The x-ray data of Ott and
McLaren,3o

with that of Billhardt.35 In all three studies, the 3PbO°ZSiOZ modifi-

for their 3PbO:ZSiOz compound, was also in agreement

cation was obtained only through devitrification and was found to be
unstable at temperatures above 585°C - 650°C. Attempts by Smart
and Glasser 3‘ and by Ott and McLaren 3° to synthesize this com—
pound from sintered mixtures of other crystalline phases failed, lead-
ing Ott and McLaren 3° to question the compound's existence as an

1

equilibrium phase and causing Smart and Glasser 3 to conclude that

36

 

 

 

 

  

 

    
  

 

 

 

 

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37

Table 2. The d-spacings and relative intensities of the x-ray spectra
for the low temperature phase in devitrified 70-30 glass.

 

dIA) I/Ig
3.90 33
3.41 50
3.18 83
3.14 50
2.99 100
2.93 67
2.05 25

 

 

 

 

38

the compound was metastable. Although there is some doubt as to
whether or not this compound is stable in the strict thermodynamic
sense, these investigators recognized that below 585°C - 650°C this
modification of 3PbO-ZSiOz is a very persistent phase. In fact,

I found that this phase only partially decomposed

Smart and Glasser 3
after being heated at 580°C for 21 days. In the course of this study,
decomposition of the major phase was never observed.

The d- spacings and relative intensities attributed to the low-
temperature, secondary phase in devitrified 70-30 glass agree quite
well with those listed by Ott and McLaren 3° for a low-temperature
polymorph of 2PbO'SiOZ. In addition, the presence of this phase in
the x-ray spectra for glass devitrified only at 450°C is in accord with
the polymorphic transition reported by Ott and McLaren 30 at 460°C:I:
15°C for 2PbO-SiOZ. Smart and Glasser 3‘ and Billhardt 3‘ also
observed the low 2PbO'SiOZ phase. Three metastable and one stable
polymorphs of 2PbO'SiOZ were reported by Smart and Glasser,31
designated L, M, M', and H respectively. The interphase relation-

ships were seen as:

0 i 0 0 i 0
Metastable: L2 PbO'SiOZMMZPbo-SiOZMM'ZPbOSiOZ

Stable: HZPbO‘ SiO‘2

Although x-ray data was not given by Smart and Glasser 3’ for L2PbO'
8102, their L2PbO'SiO2 phase was identified with the low-ZPbO'SiOZ
phase obtained by Ott and McLaren. 3° Thus, the low-temperature,
secondary phase present at 450°C in devitrified 70- 30 glass is probably
low-ZPbO-SiOZ. The absence of any 2PbO-SiOZ compound (e. g. M,

M', or H) at temperatures above 450°C is puzzling since the 72 hour
heat treatments used in this study at 500°C and 550°C to devitrify
samples for x-ray analysis are comparable with the 72 hour heat treat-
ments used by Ott and McLaren 3° and the 24 and 96 hour heat treat-

ments utilized by Smart and Glasser.in

39

The addition of P205 to the 70-30 base glass did not alter the
crystalline products of devitrification. X-ray diffraction spectra for
the 70-30 glass containing 1. 0 mole percent P205, after extended heat
treatment at 400°C and 550°C, are presented in Figure 9. Major peaks
can be observed for specimens heat treated for 10 days at 400°C, as
shown in curve (a). The positions of these initial peaks agree with
the positions of the major peaks of the crystallized 70-30 base glass.
The x-ray spectra of the 70-30 base glass and the 70-30-1. 0 glass,
after devitrification at 550°C, are presented in Figure 10 for compar-
ison. The major peak positions are identical, corresponding to the
3PbO°ZSiO2 polymorph. There are, however, a few differences in the
two spectra. The resolution of the three main peaks, with d- spacings
of 2. 87A (31. 1°), 2. 861 (31. 3°), and 2. 811 (31. 8°) is not as good in the
devitrified 70-30-1. 0 composition as it is in the 70- 30 composition.
Moreover, the peak at 1. 81A (50. 5°) and the doublet at 1. 87A (48. 7°),
observable in the spectra of the 70- 30 glass devitrified at 550°C, are
missing from the spectra of the 70- 30-1. 0 composition. These anom-
alies, however, are also present to some extent in the spectra of the
70- 30 composition, since the 1. 813. peak and the 1. 87A doublet became
pronounced only in specimens heat treated at 550°C, and are probably
due to differences in the extent to which crystallization has progressed.
The d- spacings and relative intensities for the major peaks of the
70-30-1. 0 composition after devitrification at 550°C are given in Table
l. The results of x-ray diffraction studies of devitrified 70- 30-0. 5
glass were consistent with those of the devitrified 70-30-1. 0 compo-
sition.

The x-ray diffraction spectra for the 60-40 base glass are pre-
sented in Figure 11. The crystallization process is slower in this
glass than in the 70-30 composition, as evidenced by the broad peaks
for the 450°C heat treatment, curve (b), compared to the sharp, clear—
ly defined peaks of the 70-30 glass after the same heat treatment, as

seen in Figure 8, curve (b). The diffraction spectra of specimens

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41

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42

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43

crystallized at 550°C show that the major crystalline phase is the same
3PbO’ZSiO2 polymorph that was observed for the 70- 30 composition.
The peaks attributed to this phase are marked by solid arrows on
curve ((1), and the d- spacings and relative intensities are given in
Table 1. In addition, the devitrified 60-40 glass seems to contain a
second crystalline phase, as seen by the additional peaks, marked
with open arrows on curve ((1). The positions of the peaks suggest

that this phase is different from the low-2PbO-SiOZ modification found
in the devitrified 70- 30 glass. Since the 60-40 composition lies to the

right of the 3PbO-ZSiO2

sonable to expect crystallization of a PbO-SiO2 compound. Three crys-

eutectic on the phase diagrams, it seems rea-

talline forms of composition PbO'SiOz were reported by Smart and
Glasser:31 Alamosite, low-PbO'SiOZ, and "hexagonal" PbO'SiOZ.
Only Alamosite was considered thermodynamically stable. The x-ray
data obtained in this study for the secondary crystalline phase in de-
vitrified 60-40 glass closely resembles that presented by Smart and
Glasser31 for "hexagonal" PbO'SiOZ. The d- spacings and relative
intensities for this phase are listed in Table 3.

The crystallization products of 60-40 glasses were not changed
by the presence of P205. X-ray diffraction spectra for the 60-40-0. 5
composition, after heat treatment at 400°C and 550°C are shown in
Figure 12, while the corresponding d- spacings and relative intensities
are listed in Tables 1 and 3. The major crystalline phase is 3PbO°
28102 (the low temperature modification), while the secondary phase

is probably the ”hexagonal" PbO'SiOZ reported by Smart and Glasser.31

5. 2 Microstructure of Glasses

 

Electron microscope examination of the quenched glasses indi-
cated that all glass compositions contained submicroscopic crystals.
Representative micrographs taken from replicas of the quenched 70-30
based glasses are shown in Figure 13. Crystals apparently precipitated

from the glass during the quenching operation, so that each glass was

44

Table 3. The d-spacings and relative intensities of the x-ray spectra
for the secondary crystalline phase in devitrified 60-40 and 60-40-0, 5
glass.

 

 

 

 

I71Q
d(ii) 60-40(550°C) 60-40-0. 5(550°C)
3.46 100 55
3.41 80 45
3.19 90 82
3.10 60 55
3.02 60 36
2.91 100 90
2. 69 50 .-
2.01 90 100

 

 

 

 

 

 

m A
V M I O O H m o o O
0 NH com

 

4s
a
( {cos meaitqm) Kitsuaiul

 

 

 

 

Figure 13. Electron micrographs of replicas of quenched (a) 70-30
glass, (b) 70-30- 0. 5 glass, and (c) 70-30-1. 0 glass.

47

actually composed of a fine dispersion of crystals amidst a glassy
matrix. In the 70-30 and 70-30-0. 5 glasses, the crystals occurred
in small clusters, whereas in the 70-30-1. 0 glass they appeared as
lath-like single crystals. In addition to the fine dispersion of crys-
tals, each glass also contained larger regions which were densely
crystallized. One such region is shown in Figure 14, for the 70- 30-
0. 5 composition, where a distinct radial growth pattern is evident.
Microstructures of the quenched 60-40 based glasses were similar
to those of the 70-30 and 70-30-0. 5 glasses.

No evidence of widespread phase separation was found for any
of the glasses studied. However, some phase separated regions were
observed in front of crystalline growth fronts, as seen for example
in Figure 15. A narrow phase separated zone separates a heavily
crystallized region from the surrounding matrix. Since P205 is apt
to be rejected by growing crystals, the phase separated zone is likely
to be of higher P20 content than the surrounding matrix. If this is

5
the case, then higher concentrations of P205, e. g. the 2 to 3 mole

 

percent concentrations used by Pavlushkin et al. f! 7" 3 could very
well promote phase separation. This would not, however, explain the
presence of phase separation in the binary lead silicate glass studied
by Vogel.4 The absence of phase separation noted in this study is in
agreement with the expectations reported by Shaw and Uhlmann 5 for
lead silicate glasses containing more than 50 mole percent PbO.
Continuation of the crystallization process, through isothermal
heat treatment, led to surface crystallization and/or internal crys-
tallization, depending upon the glass composition and the heat treat-
ment temperature. Some degree of surface crystallization was pre-
sent in all devitrified glasses. Development of the crystalline surface
layer was slow at 400°C and began with the formation of discrete
crystallites. Micrographs of surface crystals in the 70-30, 70-30-
0. 5, and 70-30-1. 0 glasses after 42 hours of growth at 400°C are pre-

sented in Figure 16. As can be seen, the number of crystallites per

48

 

Figure 14. Electron micrograph of radial crystal growth pattern in
quenched 70-30-0. 5 glass.

49

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0.”...- ‘r’ n 1. l

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a, me.

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Figure 16. Micrographs of surface crystallites in (a) 70- 30 glass,
(10) 70-30-0. 5 glass, and (c) 70-30-1. 0 glass after devitrification at
400°C for 42 hours.

51

unit area increased significantly with the P205 content of the glass.
Occasionally the crystallites formed in clusters or strung end-to-end
as if scratches or other surface irregularities had acted as preferen—
tial nucleation sites. In both the 70-30 and 70-30-0. 5 glasses, sur-
face crystallites had a distinct radial fiber structure. No such struc-
ture was apparent in the surface crystallites of the 70-30-1. 0 compo-
sition. Possibly the structure is of the same type only on a much
‘ finer scale. The 70-30-1. 0 composition also contained a number of
single crystals, dispersed among the crystallites. These appear as
short, dark fibers in Figure 16(c). In all compositions, the surface
crystallites grew in both size and number as crystallization continued,
and in the 70- 30- 1. 0 composition an increasing number of single crys-
tal fibers appeared. Eventually the glass surface was covered with
a crystalline layer. Thereafter, the surface layer grew inward as a
unified front. At higher temperatures, the surface crystals grew as
lath-like single crystals, as shown in Figure 17. This morphology
change was gradual and probably resulted from an aging process.

Such a process was reported by Burnett and Douglas 37

for crystal-
lization in the soda-baria- silica system, where a Spherulitic struc-
ture was observed to disintegrate into detached fibers upon extended
heat treatment. Development of the crystalline surface layer in the
60-40 based compositions was similar to that in the 70- 30 and 70- 30-
1.0 glasses.

Although P205 did not affect the crystallization products of the
base glasses, it did have a marked effect on the microstructure of the
devitrified glasses. The 70-30 and 60-40 glasses were dominated by
surface crystallization at all temperatures studied. These composi-
tions crystallized into fine-grained materials as shown in Figure 18.
The faint boundaries commonly seen in these samples, e. g. along AB
in Figure 18(a), seemed to result from the intersection of different
growth fronts. Occasionally, these compositions contained coarser,

more fibrous crystals. At high temperatures, crystallization was

52

 

H.011 11”“ ;, I

Figure 17. Micrograph of surface crystals in 70-30-0. 5 glass after
devitrification for 1 hour at 400°C followed by 12 hours at 450°C.

fl“). ,. .

53

 

 

Figure 18. Micrographs of (a) 60-40 glass devitrified at 550°C for
1 hour and (b) 70-30 glass devitrified at 550°C for 72 hours.

r..4.-W...ml|¥‘

54

rapid and accompanied by extensive deformation and cracking. Voids
were frequently observed in these samples, and the devitrified samples
seemed to be quite mechanically weak.

With the addition of 0. 1 mole percent P205, the crystalline sur-
face layer became coarser and more fibrous; however, no internal

crystallization was observed. The surface layer of all glasses con-
taining P205 was composed of rather coarse, branched fibers which
were directed into the glass interior and which were separated from
one another by the same type of finely crystallized material as found
in the devitrified base glasses. Figure 19(a), a micrograph of devit-
rified 60-40-0. 1 glass taken with crossed polarizers, shows the
branching nature of the fibers, while Figure 19(b), a micrograph of
devitrified 60- 40-0. 5 glass, shows the fine grained material which
separated the fibers. The 60-40-0. 1 glass was dominated by surface
crystallization at all temperatures studied.

In concentrations of O. 5 and l. 0 mole percent, P205 promoted
internal crystallization in the form of spherulites. The number of
Spherulites nucleated per unit volume was extremely sensitive to both

the P 0 concentration of the glass and the crystallization temperature.

2 5
At 550°C, glasses containing 0. 5 mole percent P205 crystallized com-

pletely from the surface, as shown in Figure 20(a) for the devitrified
60-40-0. 5 composition. In the 70- 30- 1. 0 composition, however, the
advance of the surface layer was limited by the presence of a number
01' spherulites which grew in the interior of the glass, as shown in
Figure 20(b). Heat treatment at 500°C again led predominately to sur-
face crystallization in the 60-40-0. 5 and 70- 30-0. 5 compositions;
however, a few spherulites were observed in the interiors of these
glasses, as shown in Figure 21 for the 60-40-0. 5 composition. In the
70— 30-1. 0 composition the lowered crystallization temperature re-
sulted in an increased number of spherulites, and hence an even nar-

rower surface layer than was seen for specimens devitrified at 550°C.

The number of spherulites per unit volume continued to increase as

55

' V x. _
. "v o ' ,_ ‘
{Aw-1w,

Jr- .,:,

 

K‘s-(r;

Figure 19. Microstructure of crystalline surface layer in
(a) 60-40-0. 1 glass devitrified for 1 hour at 550°C and (b) 60-40-0. 5
glass devitrified for 72 hours at 550°C.

56

 

Figure 20. Microstructures after devitrification at 550°C for 72
hours of (a) 60-40-0. 5 glass and (b) 70-30-1. 0 glass.

I III- II. .I III-Ii

57

 

Figure 21. Microstructure of 60-40-0. 5 glass after devitrification
at 500°C for 72 hours.

58

the crystallization temperature was lowered, so that at 400°C all of
the glasses containing P205 (except the 60-40-0. 1 composition) crys-
tallized with a high concentration of spherulites, as shown in Figure
22. Devitrification at 400°C was significantly more effective at ini-
tiating Spherulitic growth than devitrification at any of the higher

temperatures. The nucleation rates for spherulites at 400°C was

determined to be:

 

 

Glass Composition Nucleation Rate(number-cm-3-min-l)
70-30-1.0 3.29 x 106
70-30-0. 5 2.10 x 104
60-40-0. 5 "' 3 x 105

At 450°C,the nucleation rate in the 70- 30- 1. 0 glass dropped to 5. 5 x

1 Oz spherulites— cm-3-min- 1, while that in the 70- 30-0. 5 and 60-40-0. 5
compositions became too low to measure. Nucleation rates were too
low for measurement in all compositions at 500°C and 550°C. Since
preliminary investigations indicated that lowering the crystallization
temperature to 350°C or 300°C would not further increase the concen-
tr ation of spherulites, 400°C was taken to be the nucleation tempera-

ture for Spherulitic growth in the glasses containing P205.

5 . 3 Spherulitic Growth Rate
Since the 60-40-0. 5 and 70- 30-0. 5 glasses were dominated by

Surface crystallization at all temperatures except 400°C, a two- step
heat treatment was used for these compositions to study Spherulitic
growth rates at 450°C, 500°C, and 550°C. In the first step, a heat
treatment consisting of 6 hours at 400°C for the 60-40-0. 5 composition
and 1 hour for the 70-30-0. S composition was used to induce spheru-
litic nuclei formation. Then, in the second step, the ”nucleated”
glasses were heat treated at one of the higher temperatures to pro-
duce spherulitic growth. The 70-30-1. 0 composition was not given

any nucleation treatment, since it displayed some internal

I

1;

~-t

1 ’1‘ '1
e .‘11

I, .‘

3%;

if
‘. '13

'..Q

 

Figure 22. Micrographs of spherulites in (a) 60-40-0. 5 glass
devitrified for 14 days at 400°C, (b) 70-30-0. 5 glass devitrified for 8
days at 400°C, and (c) 70-30-1. 0 glass devitrified for 4 days at 400°C.

60

crystallization at all temperatures; and the 60-40—0. 1 composition was
not used for Spherulitic growth measurements, since even long nucle-
ation treatments at 400°C did not initiate widespread Spherulitic
growth.

Under isothermal growth conditions, spherulite radii increased

linearly with time. Growth curves for crystallization at 400°C, 450°C,

500°C, and 550°C, plotted as the mean spherulite radius versus
growth time, are given in Figures 23, 24, and 25 for the 70-30-0. 5,
70-30-1. 0, and 60-40-0. 5 compositions respectively. The nucleation
heat treatments given to the 70-30-0. 5 and 60-40-0. 5 glasses had no
significant effect on the determined growth rates, as seen in Figure

25, curve (b), where samples nucleated for 1 hour at 400°C show the

same growth rate as those nucleated for 6 hours. Spherulitic growth

rates, as determined from the slope of each curve, are given in

Table 4. The growth rates were extremely temperature sensitive.

Plots of ln(growth rate) versus (absolute temperature)"1 are presented

in Figure 26, and the activation energy for Spherulitic growth, as
calculated from the slopes of these plots, is about 80 kcal/mole for

all compositions. From viscosity data of Bair,3° the activation

energy for viscous flow in a lead silicate glass containing 60 mole
percent PbO can be estimated to be about 84 kcal/mole for tempera-
tures up to and including 400°C. No data was available for the tem-
Perature range 400°C to 550°C. Thus, the activation energy for spher-
ulitic growth is close to that required for viscous flow in binary lead
Silicate glass.

Among the glasses containing P205, growth was much more
rapid in the 70-30 based glasses than it was in the 60-40 based glasses.
P205, itself, markedly effected the Spherulitic growth rate. Increased
concentrations of P205 slowed spherulite growth substantially, as seen
by comparing the growth rates for the 70-30-0. 5 and 70-30-1. 0 com-
POSitions. Although no quantitative data was obtained for the base com-

poSitions, observations of surface layer growth, indicated that

61

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.,

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In GROWTH RATE

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Figure 26. Graph of ln(crystal growth rate) versus the reciprocal
of the absolute temperature. (Data points for the 70-30-0. 5, 70-
30-1. 0, and 60-40-0. 5 compositions are plotted as squares, circles,
and triangles, respectively)

66

crystal growth was even faster in the base glasses than in the P205
glasses.

Histograms, showing the size distribution of spherulite popula-
tions after isothermal crystallization, are presented in Appendix A.
Comparison of these histograms shows that for each composition:

(1) the upper boundary (the largest size) of the spherulite pOpu-
lation increased with isothermal growth time, while

(2) the lower boundary (the smallest size) remained approxi-
mately constant with time for growth at 400°C and began to increase
more and more with time as the growth temperature was raised.
The inability of each composition to keep the lower size intervals
filled at higher temperatures as crystallization progresses substan-
tiates earlier observations that the nucleation process for spherulitic

growth is inhibited at temperatures above 400°C.

5. 4 Spherulitic Morphology

 

Spherulites observed in this study were composed of arrays of
crystalline fibers, arranged radially outward from a common center.
Growth occurred through a lengthening of the fibers and was accom-
panied by fiber branching, which tended to "fill- in" the interior of the
spherulite and trap matrix material between the fibers. The width of
the fibers and the degree to which they branched varied with the growth
temperature.

At 550°C, spherulites grew with coarse, Open textures. Three
basic morphologies could be distinguished, as shown in the micro-
graphs presented in Figure 27. The spherulite in Figure 27(a) is char-
acterized by an elongated major axis from which fibers fan out on
either side, creating a sheaf-like appearance. In Figure 27(b), the
spherulite is seen as a bundle of more-or-less parallel fibers, while
in Figure 27(c), it takes on a distinctly hexagonal shape. These dif-
ferent morphologies were found to arise not from three different

growth patterns, but rather from three different views of the same

67

 

Figure 27. Spherulites grown at 550°C showing typical (a) sheaf,
(b) bundle, and (c) hexagonal morphologies.

68

growth pattern, so that when a Spherulite is oriented as shown in
Figure 28, the three morphologies shown in Figure 27(a), (b), and
(c) correspond to views along the x, y, and z axis respectively.

As the growth temperature was lowered, the spherulites became
more spherical in shape, and the fibers became much finer. Figure
29 is a micrograph of spherulites in the 70-30-0. 5 composition after
growth at 450°C. The distinct sheaf-like shape observed in spherulites
grown at 550°C is not seen in spherulites grown at 450°C; however,

a slight sheaf-like arrangement of fibers can still be detected. In
spherulites grown at 400°C, the fiber structure was no longer resolv-
able with light microscopy. However, an electron micrograph show-
ing the fibrous growth morphology of spherulites grown at 400°C is
presented in Figure 30. Spherulite texture was also effected by the

P 0 content of the glass, with increased concentrations of P205 pro-

2 5
ducing finer spherulitic textures.

Several growth stages were observed in spherulites crystallized
at 550°C or 500°C. In the earliest stages of growth, spherulites con-
sisted of just a few fibers, or platelets, which spread out slightly from
center, as shown in Figure 31. Development of these spherulites con-
tinued with repeated branching, which caused the fibers to fan out more
from center, forming fuller sheaves. Figure 32, an electron micro-
graph of the outer tips of a spherulite grown at 550°C in 60-40-0. 5
glass, shows the branching nature of the fibers. As growth continued,
branching fibers gradually encircled regions on either side of center,
forming eye-like regions made up of trapped matrix, as shown in
Figure 33. At lower temperatures, spherulite morphology did not
change significantly with growth, as even the smallest spherulites
observed had a fully developed, spherically symmetric appearance,
indicating that the development stage was completed while the spheru-
lites were still of submicrosc0pic size. At all temperatures, a
Spherulite's radial growth continued until halted by its impingement

with neighboring spherulites. In this way a polyhedral network of

69

 

 

Figure 28. Diagram of spherulite orientation.

7O

 

Figure 29. Micrograph of a spherulite in 70-30-0. 5 glass after
growth at 450°C.

71

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y%3a

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72

 

Figure 31. Micrograph of an early growth stage in a spherulite
grown at 550°C.

73

 

Electron micrograph of branching fibers in a spherulite

grown at 550°C.

Figure 32.

74

 

Figure 33. Micrograph of a mature spherulite grown at 550°C.

75

inter- spherulitic boundaries was formed in highly crystallized samples,
as shown in Figure 34 for the 70-30-0. 5 composition. Matrix trap-
ped inside the spherulite gradually crystallized, producing the fine-

grained texture.

5. 5 Role of P205

 

Our investigations show that, as reported by Pavlushkin et al.,l'z:3

 

P205 is effective in promoting bulk crystallization in lead silicate
glasses containing between 50 and 60 mole percent PbO. The pre-
sence of P205 in sufficiently high concentrations (0. 5 mole percent

for the two glass compositions studied) leads to the creation of internal
nucleation sites from which spherulitic growth occurs. The exact

role played by P205 in promoting Spherulitic growth is, however,
unclear. Since the addition of O. 1 mole percent P205 led to a more
fibrous growth morphology, but not to spherulitic nucleation, the role
of P205 may be two-fold, namely (1) to nucleate internal crystals of
the major crystalline phase, 3PbO-.?.Si0‘Z and (Z) to promote a fibrous
growth morphology. Pavlushkin _e_’£_a_l_.3 proposed that the presence of
P205 led to phase separation followed by the formation of crystalline
nuclei. Since no evidence of widespread phase separation was found

in this study, it seems unlikely that phase separation (even though it
may occur with higher P205 concentrations) is a pre-requisite for
internal nucleation in these glasses. Furthermore, the matrix crys-
tals, present in the quenched glasses, do not seem to promote inter-
nal crystallization of the major crystalline phase. The fact that no
phosphorus compounds were observed during devitrification, coupled
with the observation the P205 led to greatly reduced spherulitic growth
rates, suggests that P205 may be acting solely as an impurity in these
glasses, promoting internal nucleation by reducing the interfacial

energy between the glass and the crystal. Such a role was suggested

by James and Keown 39 to explain the increased nucleation rate of

76

 

Figure 34. Network of inter-spherulitic boundaries in a completely
devitrified 70- 30-0. 5 glass. The arrows indicate the boundary
network.

77

spherulites in lithium silicate glasses containing 1 mole percent
P205.

Recently, crystal growth morphologies have been examined in
relation to the entropy of fusion, and the importance of this parameter

in predicting crystal-liquid interface structure and crystal morphology

d.4°: 4‘: 42: 43 4°, 4‘: *2 calculated the

has been demonstrate Jackson
free-energy change associated with the addition of atoms to available
sites on growth interfaces of crystals. According to his model, the
surface free-energy change can be expressed in terms of a parameter
a, which can be written as

AS

a=__£'lg

R

where ASfm is the entropy of fusion, R is the universal gas constant,
and C is the fraction of total bonding energy contributed by the addition
of atoms at adjacent sites on the surface. C , always less than unity,
depends upon the structure of the crystal growth face, being largest
for the most closely-packed planes. The surface free energy was
found to vary with a as shown in Figure 35. The position of the minima
depends upon the value of a. For a<2. O, the minimum free- energy
change occurs for half of the available surface sites being filled,

which is representative of a growth surface which is rough on an atomic
scale. For 0! > 2. 0, however, minima in the surface free-energy
change occur for surfaces in which there are either a very small or
very large fraction of filled sites, which corresponds to a growth plane

which is smooth on an atomic scale. Since,

the oz<2 situation occurs for

78

 

AF
3
NkT E

 

 

 

RELATIVE FREE ENERGY

 

 

 

 

-O.

 

0.1 0.2 0.3' 0.4 0.5 0.6 0.7 0.8 0.9 1.0

OCCUPIED FRACTION OF SURFACE SITES

Figure 35. Free energy of an interface versus occupied fraction
of surface sites. ‘2

79

ASfm
_— < 2
R g 3
OI‘,
As < .23.
fm g

Since §< 1, materials with AS m< 2R fall into the oz< Z category, and

f
are expected to crystallize with "atomically" rough surfaces and ex-
hibit low growth rate anisotropy and non-faceted morphologies. The
a > 2 situation occurs for

2R

ASfm > T

Since g is E 0. 5 for close-packed planes, materials having ASfm > 4R
are expected to crystallize with ”atomically" smooth surfaces for close-
packed planes (which corresponds to the o: > 2 situation) and rougher
surfaces along less closely packed planes (where :1 falls below 2 due
to the decreasing c). For these materials, a high growth-rate anisot-
ropy and a faceted morphology along close-packed planes is also
expected.

Materials with high entropies of fusion may also, according to
Uhlmann,"’3 be expected to have growth rates not well described by
the standard models for crystal growth. These materials often show
a tendency towards spherulitic growth at large undercoolings, which
become even more pronounced in the presence of impurities. The
growth morphologies observed in this study closely match those pre-
dicted by Jackson and Uhlmann for materials with high entropies of

fusion. Based on Shartsis and Newman's 4“

study of energy relations
in the PbO-SiO2 system, the entropy of fusion for 3PbO'ZSiOZ can be

estimated to be approximately 4. 3R. In the base glasses and the

80

60-40-0. l composition, where spherulitic growth was not observed,
a faceted growth morphology was apparent, as seen in Figure 36.
As the P205 concentration was increased, the tendency towards
spherulitic growth became more pronounced, finally becoming the
preferred mode of crystallization with the addition of l. 0 mole per-
cent P205.

Certain morphological features and growth patterns have been
found to be characteristic of spherulitic crystallization, irrespective
of the system in which crystallization occurs. Typically, spherulites
are composed of fibers which radiate outward from a common center.
The fibers are of approximately constant thickness and have a prefer-
red crystal axis along the radial direction. Fiber thickness increases
with increased growth temperature and decreases with increased
impurity content. The fibers are separated from one another by un-
crystallized melt. Spherulitic development proceeds from bundles of
parallel fibers which fan out to form intermediate structures called
sheaves. As growth continues a spherical shape gradually develops.
Spherulitic growth rates, characteristically, are constant with time
under isothermal conditions.

A theory to account for spherulitic crystallization has been pre-
sented by Keith and Padden. ‘5 Two properties, high viscosity and
the presence of impurities, are seen as fundamental to spherulite
formation. The presence of impurities in melts with high viscosities,
and low growth rates, leads to the build-up of a narrow impurity rich
layer ahead of the growing crystal, which causes the interface to be-
come unstable with respect to pertubations. Fibrous growth results,
with the fiber thickness determined by 6, the thickness of the interface.
6 is approximated as D/ G, where D is the diffusion coefficient for the
impurity in the melt and G is the growth rate. The variation in fiber
thickness with growth temperature, characteristic of spherulites, is
explained in terms of the change of 6 with temperature. At large under-

coolings, 6 is small, leading to relatively narrow fibers. As the

81

 

Figure 36. Electron micrograph of faceted crystal growth in a
60-40-0. 1 glass after devitrification at 450°C for 1 hour.

82

crystallization temperature is raised 6 increases, and the fiber width
broadens. Once fibrous growth is established the spherulite morphol-
ogy can result, through the sheaf-to- spherulite development sequence,
provided that a source of branching is present. Branching is seen as
resulting from the encounter of a growing fiber with any singularity or
region of disorder whose size is comparable with 6. As 6 decreases,
the frequency of branching is assumed to increase due to the greater
probability of encounters with singularities large enough to cause
branching. Thus, the size at which the sheaf-to- spherulite transfor-
mation occurs is expected to decrease as the spherulite's texture be-
comes finer, with the more frequent branching, expected for narrower
fibers, leading to a more rapid spherulite development.

The spherulite morphologies and growth patterns observed in
this study were consistent with those characteristic of spherulites in
general. The theory of spherulitic growth proposed by Keith and
Padden ‘5 can account reasonably well for the spherulitic growth ob-
served in this study, and the implication is again that P205 functions
as an impurity in these glasses. Therefore, it seems likely that the
crystallization behavior observed in the glasses containing P205 re-
sults from its role as an impurity -- increasing the internal nucle-
ation rate by reducing the interfacial energy between the glass and the
crystal and promoting a spherulitic morphology by enhancing fibrous

growth.

CHAPTER VI
SUMMARY

Concentrations as low as 0. 5 mole percent P205 promoted internal
crystallization in the lead silicate glasses studied. Internal crys-
tallization initiated from discrete centers within the glass and ex-
hibited a spherulitic morphology.

No evidence of widespread phase separation was found in any of the
glasses. This was unexpected on the basis of the work of Pavlushkin

et a1.3 and Vogel. 4 All quenched glasses were, however, partially

 

crystallized, containing dispersions of submicrosc0pic size crys-
tals.

In the low concentrations studied, P205 had no detectable effect on
the phases which occurred as crystallization products. The major
crystalline phase in all devitrified glasses was a low temperature
polymorph of 3Pb0- ZSiOZ. Secondary products of crystallization
were tenatively identified as low-ZPbO' SiO2 in the 70-30 based
glasses and "hexagonal" PbO- SiO2 (a polymorph of PbO. SiO2 re-
ported by Smart and Glasser”) in the 60-40 based glasses.

400°C was found to be the most effective temperature for nucleating
spherulitic growth. Devitrification at this temperature led to high
concentrations of Spherulites in all glasses containing at least 0. 5
mole percent P205. Concentrations of l. 0 mole percent P205 were
needed to produce detectable levels of spherulitic nucleation at
temperatures above 400°C.

Spherulitic radii increased linearly with time under isothermal

conditions. Spherulitic growth rates increased with increased

crystallization temperature, showing an experimental activation

83

84

energy of about 84 kcal/mole. Increased concentrations of P205
led to decreased spherulitic growth rates.

Spherulites observed in this study exhibited morphological fea-
tures and growth patterns characteristic of spherulites found in
polymers and other materials. Spherulite morphology deveIOped
through the sheaf-to-spherulite sequence, and spherulite texture
coarsened and became more open as the growth temperature was
increased. The observed features were consistent with Keith and
Padden's45 theory of spherulitic growth.

The entropy of fusion of base glasses studied was estimated to be
about 4. 3R, where R is the universal gas constant. The spheru-
litic growth habit was in accord with the prediction of J ackson“Z
and Uhlmann43 for materials with entropies of fusion greater than
2R.

The crystallization behavior observed in this study can be explained
on the basis of P205 acting as an impurity since it increased the

internal nucleation rate and promoted spherulitic crystal growth.

APPENDIX

APPENDIX A

HISTOGRAMS OF SPHERULITE
POPULATIONS IN DEVITRIFIED GLASSES

The sizes of spherulites in devitrified glass samples were stud-
ied by light microscopy. The diameters of spherulites in each sample
were measured from micrographs taken at random intervals over the
sample. This data was used to construct histograms showing the total
number of spherulites observed in each size in the devitrified sample.
Since most of the data was obtained from photographs taken at a mag-
nification of 75X, the size intervals on the histograms are presented
as 75X the actual spherulite diameter. Histograms of spherulite
populations occurring in 70-30-0. 5, 70-30-1. 0, and 60-40-0. 5 glasses
after extended heat treatments at 400°C are presented in Figures
A. l, A. Z, and A. 3 respectively. As the heat treatment time was ex-
tended at 400°C, the spherulites in each glass grew in size, occupy-
ing higher size intervals. However, the lower size intervals also
remained filled. Histograms of spherulite populations in 70-30- 0. 5,
70-30-1. O, and 60-40-0. 5 glasses after growth at 450°C are presented
in Figures A. 4, A. 5, and A. 6 respectively. Those of spherulite
populations in 70-30-0. 5, 70-30-1. O, and 60-40-0. 5 glasses after
growth at 500°C are given in Figures A. 7, A. 8, and A. 9 respectively:
while those of spherulite populations in 70-30-0. 5, 70-30-1. 0, and
60-40-0. 5 glasses after growth at 550°C are presented in Figures
A. 10, A. 11, and A. 12. As the growth temperature was increased
from 400°C, there was a tendency for the lower size intervals to
empty, as the higher intervals filled. It is clear from the histograms
of spherulite populations in these glasses at 450°C, 500°C, and 550°C
(Figures A. 4 through A. 12) that no significant amount of spherulite

85

86

nucleation occurred during growth at these temperatures.

87

 

 

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75 x Spherulite Diameter (in)

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x Spherulite Diameter (in)

Number of Spherulites

 

 

 

N \0 oo
0 O O O
c; o' o' c;
75 x Spherulite Diameter (in)

Figure A. l. Histograms of spherulites grown in 70-30—0. 5 glass at
400°C for (a) 2 days, (b) 3% days, (c) 4 days, (d) 5 days, (e) 6 days,
and (f) 14 days.

88

 

Spherulite Diameter (in)

 

 

 

 

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cont.

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89

 

Number of Spherulites

 

 

0.

75 x Spherulite Diameter (in)

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Number of Spherulites

 

75 x Spherulite Diameter (in)

 

Number of Spherulites

 

 

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75 x Spherulite Diameter (in)

Figure A. Z. Histograms of spherulites grown in 70-30-1. 0 glass
at 400°C for (a) 2 days, (b) 3} days, and (c) 4 days.

90

Number of Spherulites

 

 

75 x Spherulite Diameter (in)
b)

80"

60"

 

Number of Spherulites

 

75 x Spherulite Diameter (in)

Number of Spherulites

 

 

N v
C O
0' c5
75 x Spherulite Diameter (in)

Figure A. 3. Histograms of spherulites grown in 60—40-0. 5 glass at
400°C for (a) 14 days, (b) 18 days, (0) 25 days, and (d) 32 days.

91

d)
100 r-

00
O

0‘
O

.b
O

 

Number of Spherulites
N
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75 x Spherulite Diameter (in)

Figure A. 3. cont.

92

a)

Number of Spherulites

 

 

 

75 x Spherulite Diameter (in)

Number of Spherulites

 

 

 

75 x Spherulite Diameter (in)

C)
80

60

40

20

Number of Spherulites

    

 

o o v
F‘ N N
o' o‘ o“ 60' o' o‘ 60' o'
75 x Spherulite Diameter (in)
Figure A. 4. Histograms of spherulites grown in 70-30-0, 5 glass at

450°C for (a) 200 min., (b) 240 min., (c) 360 min., (d) 480 min., and
(e) 600 min. , after a nucleation heat treatment of 1 hour at 400°C.

40-

30

20

10

Number of Spherulites

 

Number of Spherulites

 

Figure A. 4.

93

    

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Figure A. 5. Histograms of spherulites grown in 70-30-1. 0 glass at
450°C for (a) 200 min. , (b) 275 min. , (c) 480 min., (d) 600 min. , and
(e) 720 min. .

95

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75 x Spherulite Diameter (in)

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75 x Spherulite Diameter (in)

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Number of Spherulites

 

 

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75 x Spherulite Diameter (in)

b)

Number of Spherulites

 

 

 

75 x Spherulite Diameter (in)

C)

30"

Number of Spherulites

 

 

 

75 x Spherulite Diameter (in)

Figure A. 6. Histograms of spherulites grown in 60-40-0. 5 glass at
450°C for (a) 720 min. , (b) 1080 min. , and (c) 1405 min. after a
nucleation heat treatment of 1 hour at 400°C; and for (d) 1440 min. ,
(e) 1680 min. , and (f) 1920 min. after a nucleation heat treatment of
6 hours at 400°C.

97

 

 

 

75 x Spherulite Diameter (in)

75 x Spherulite Diameter (in)
cont.

75 x Spherulite Diameter (in)

    
 

 

 

 

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p . r p - _) b . b . b r . A.

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Number of Spherulites

 

75 x Spherulite Diameter (in)

 

Number of Spherulites

 

75 x Spherulite Diameter (in)

 

Number of Spherulites

 

75 x Spherulite Diameter (in)

 

Figure A. 7. Histograms of spherulites grown in 70-30-0. 5 glass at
500°C for (a) 10 min. , (b) 15 min. , (c) 22 min. , (d) 30 min. , and
(e) 37 min. , after a nucleation heat treatment of 1 hour at 400°C.

 

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75 x Spherulite Diameter (in)

 

cont.

Figure A. 7.

100

20_

10'-

Number of Spherulites

 

 

 

 

75 x Spherulite Diameter (in)

40"
30"
20"

10-

 

Number of Spherulites

 

 

 

\oov
F‘NN
0:50

0.08;
012

oo
N
o'

75 x Spherulite Diameter (in)
40 "

30-

20'-

10-

   

Number of Spherulites

 

 

~oo<rooN~o
HNNNMM

 

0‘ c5 0' o' c5 c5
75 x Spherulite Diameter (in)

Figure A. 8. Histograms of spherulites grown in 70-30-1. 0 glass at
500°C for (a) 10 min. , (b) 15 min. , (c) 22 min. , (d) 30 min. , and
(e) 37 min. .

d)
40

30
20

10

Number of Spherulites

e)
40

30

20

Number of Spherulites

Figure A. 8.

 

101

 

 

 

 

6 6 6 6
75 x Spherulite Diameter (in)

  

Ncovwmoovw
mmvvvmmxoxoxo
6666666666

75 x Spherulite Diameter (in)

102

 

 

 

34.3 a)
'3
’53
.2: 90-
o.
m
“5 60-
L.
o
.o
E
:3
Z
O O
75 x Spherulite Diameter (in)
b)
P
90 >—

Number of Spherulites

 

 

 

0 02
0 - 04 :Eiga::;::::;g:gsg
0 O 6 ..,;;_ ’-;.;:.,,;.¢;

75 x Spherulite Diameter (in)

 

 

 

 
   

 

 

.33 C)

0H

g 120

L.

g .......

o. 90

m

“6 60 -

L. . ,

g

E 30
w \o ado
o o o -:
c;o 6<5

75 x Spherulite Diameter (in)

Figure A. 9. Histograms of spherulites grown in 60-40-0. 5 glass at
500°C for (a) 20 min., (b) 40 min., (c) 60 min., (d) 80 min. , (e) 100
min. , (f) 120 min. , and (g)l40 min., after a nucleation heat treatment
of 6 hours at 400°C.

103

 

 

 

_ p . b
x) O O 0
d 9 6 3

mosfisuonom .8 969522

0 N .
_. . NH .
0H .

O O O O O O
75 x Spherulite Diameter (in)

 

 

 

 

b p p _
e 01 6 3

$223.82an mo ponEsZ

75 x Spherulite Diameter (in)

MN .o
_. ,.h..,._. ON .o
,_ 2 .o
2 .o

 

 

«l 6
NH 6

 

 

— p p p
0 O O
\f) 9 6 3

mofifiiozam mo honfisz

75 x Spherulite Diameter (in)

cont.

Figure A. 9.

104

gmL
'5‘
$4
.2
a, 90r
(I)
‘0—0
0 60'
$4
.8
E 30’
:3
z

 

 

 

75 x Spherulite Diameter (in)
Figure A. 9. cont.

33a)
:1
:1
I-c
330
O.
U)
'0-1
020
I-a
2
E10
5
2
3‘2")
;::8
:3
:3
E6
(I)
‘84
$4
8
E2
:3
z

105

 

 

 

 

. --< "/
,., y \

V'LnOt‘wO‘O HN

 

6 6 6 6 6 6 .—2 .4 ...:
75 x Spherulite Diameter (in)

2‘

 

 

 

CDO‘Ol-‘NMV‘IDQ boo
664.44.44.54 .—:—:

75 x Spherulite Diameter (in)

Figure A. 10. Histograms of spherulites grown in 70-30-0. 5 glass at
550°C for (a) 3 min. and (b) 5 min. , after a nucleation heat treatment
of 1 hour at 400°C.

106

a)

Number of Spherulites

 

 

 

75 x Spherulite Diameter (in)

       

 

 

 

 

b>
IF:
:3
L.
0
ti 30*
m
‘O-l
O 20'
h
3
E 10"
f") V“ In \0 1‘ co 0‘ O H
6 6 6 6 6 6 6 .J .4
75 x Spherulite Diameter (in)

8 c)
.t'.
H
:3
s.
52’
Q 30
m
'04
0 20
s.
0
E 10
:1
Z , r

7 ' o~o Mmmvm Mm

Jain'w'm'u'w' «iv;

75 x Spherulite Diameter (in)

Figure A. ll. Histograms of spherulites grown in 70-30-1. 0 glass at
550°C for (a) 3 min., (b) 5 min., and (c) 10 min. .

107

30”
20"

1" J).

M V‘ LO \0
6 o' 6 o'
75 x Spherulite Diameter (in)

Number of Spherulites

 

b)

 

Number of Spherulites

 

 

6 6 6 6 6
75 x Spherulite Diameter (in)

act

10'-

    

Number of Spherulites

 

1‘ CD 0‘ O H

6 6 6 .4 .4
75 x Spherulite Diameter (in)
Figure A. 12. Histograms of spherulites grown in 60-40-0, 5 glass at
550°C for (a) 15 min. , (b) 22 min. , and (c) 30 min. , after a nucleation
heat treatment of 6 hours at 400°C.

10.

ll.

12.

l3.

14.

15.

LIST OF REFERENCES

N. M. Pavlushkin and G. P. Lisovskaya, Tr. Mosk. Khim. -
Tekhnol. Inst. No. 55, 81(1967).

N. M. Pavlushkin and G. P. Lisovakaya, Steklo, _3_, 114(1968).
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Lisovskaya, Eksp. Issled. Mineraloobrazov., Mater. Vses.

Soveshch. Eksp. Tekh. Petrogr., 8th 1968, 169(1971).

Werner Vogel, Structure and Crystallization of Glasses, Pergamon
Press, 1971.

 

R. R. Shaw and D. R. Uhlmann, J. Noncrystalline Solids, _1_,
474(1969).

S. D. Stookey, in Ceramic Fabrication Processes, W. D. Kingery
ed., The M. I. T. Press, Cambridge, 1958.

 

D. R. Uhlmann, in Materials Science Research, Vol. 4, T. J.
Gray and V. D. Frechette eds. , Plenum Press, New York, 1969.

 

 

G. 0. Jones, Glass, Methuen, London, 1956.

J. W. Christian, The Theory of Transformations in Metals and

 

Alloys, Pergamon Press, 1965.

R. Becker, Proceedings of the Physical Society of London, 52,
71(1940). ""

H. R. Swift, J. Am. Ceram. Soc., 39, 170(1947).

Robert H. Doremus, Glass Science, Wiley-Interscience, New
York, 1973.

 

W. D. Scott and J. A. Pask, J. Am. Ceram. Soc., 14, 181(1961).
J. G. Morley, Glass Tech., _6_, 77(1965).

Daniel R. Stewart, in Introduction to Glass Science, L. D. Pye,

 

108

109

H. J. Stevens, and W. C. LaCourse eds., Plenum Press, New
York, 1972.

16. S. D. Stookey, Corning Research, 73(1961).
17. R. D. Maurer, J. Appl. Phys., _3_3_, 2132(1962).

18. T. J. Barry, D. Clinton, L. A. Lay, R. A. Mercer, and R. P.
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1‘). T. J. Barry, D. Clinton, L. A. Lay, R. A. Mercer, and R. P.
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20. R. F. Geller, A. S. Creamer, and E. N. Bunting, J. Research
Nat. Bur. Standards, _13, 237(1934).

21. H. C. Cooper, L. 1. Shaw, and N. E. Loomis, Am. Chem. J.,
:13, 461(1909).

22. Siegfried Hilpert and Richard Nacken, Berichte. der Deut. Chem. ,
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23. H. C. Cooper, E. H. Kraus, and A. A. Klein, Am. Chem. J.,
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24. L. G. Berezkina and D. M. Chizhikov, Russ. J. Inorg. Chem.,
1, 442(1962).

25. H. F. McMurdie and Elmer N. Bunting, J. Research Nat. Bur.
Standards, 23, 543(1939).

26. J. F. Argyle and F. A. Hummel, J. Am. Ceram. Soc., :43,
452(1960).

27. E. V. Smirnova, in The Structure of Glass, Vol. 7, E. A. Porai-
Koshits ed., Consultants Bureau, New York, 1963.

 

28. E. V. Smirnova, Izv. Akad. Nauk. SSSR, Neorg. Mater., E,
2045(1966).

29. E. V. Smirnova, ibid., 3, 1230(1967).

30. W.Richard Ott and Malcolm G. McLaren, J. Am. Ceram. Soc.,
53, 374(1970).

31. R. M. Smart and F. P. Glasser, ibid., 22, 378(1974).

32.

33.

34.

35.

36.

37.

38.

39.

40.

41.

42.

43.

44.

45.

110

R. A. Rita and C. G. Bergeron, ibid., 22, 274(1976).

T. F. Beals and W. C. Bigelow, in Symposium on Advances in
Electron Metallography and Electron Probe Microanalysis, ASTM

 

 

Spec. Tech. Publ. No. 317,1962.

H. R. Limb, The Review of Scientific Instruments, _4_1, 874(1970).
H. W. Billhardt, The American Mineralogist, 24, 510(1969).

H. W. Billhardt, Glastech. Ber... :15, 498(1969).

D. G. Burnett and R. W. Douglas, Phys. Chem. Glasses, E,
117(1971).

George J. Bair, J. Am. Ceram. Soc., _1_?_, 339(1936).
P. F. James and S. R. Keown, Phil. Mag., Ser. 8, 39, 789(1974).

D. A. Jackson, in Liquid Metals and Solidification, ASM,
Cleveland, 1958.

 

K. A. Jackson, in Growth and Eerfection of Crystals, R. H.
Doremus, B. W. Roberts, and D. Turnbull eds., Wiley, New
York, 1958.

 

K. A. Jackson, in Progress in Solid State Chemistry, Vol. 4,
H. Reiss ed. , 1967.

 

Donald R. Uhlmann, in Advancesin Nucleation and Crystallization
in Glasses, American Ceramic Society, Columbus, 1972.

 

 

L. Shartsis and E. S. Newman, J. Am. Ceram. Soc., 31, 213
(1948). “—

H. D. Keith and F. J. Padden, Jr., J. Appl. Phys., 34, 2409
(1963). "‘

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