Nitrogen doping of fused silica and silicate glasses: a study of transport and optical properties

Journal of Non-Crystalfine Solids 102 (1988) 181-195 North-Holland, A m s t e r d a m

181

N I T R O G E N D O P I N G O F FUSED SILICA AND SILICATE GLASSES: A STUDY O F T R A N S P O R T AND OPTICAL P R O P E R T I E S J. S C H R O E D E R Physics Department, and Center for Glass Science and Technology, Rensselaer Polytechnic Institute, Troy, N Y 12181, USA

Joseph J. J A R E K Inorganic Materials Laboratory, Corporate Research and Development, General Electric Company, Schenectady, N Y 12301, USA

Our work has demonstrated for the first time that relatively large a m o u n t s of atomic nitrogen can be incorporated in fused silica, rather than Vycor, by a high temperature N H 3 cracking process. The nitrogen analysis process shows that the nitrogen is chemically dissolved in the molecular structure of the fused silica glass and we obtain a product which is a stable and transparent glass. A kinetic study of this nitriding process has shown that nitrogen incorporation in a glass will occur even at temperatures as low as 500 o C, although the best efficiencies in the reaction are attained at temperatures near or slightly below the glass transition temperature of the glass. The kinetic study also shows that o p t i m u m nitriding conditions for fused silica are a N H 3 / N 2 ratio of 3:1 at a nitriding temperature of - 1 0 0 0 ° C. We discuss the possible structural scheme that allows increased lattice stiffness of a nitrided glass relative to a non-treated glass. This in turn has important consequences for understanding the problem of glass stability (i.e. devitrification) in silicate based glasses.

1. Introduction

Glass may devitrify by a process involving the nucleation of cristobalite crystal on external surfaces and a subsequent growth of these crystals into the body of the glass specimen. This nucleation is impurity catalyzed. In most glasses the rate of homogeneous crystal nucleation (nucleation unassisted by foreign substances) is too small to be measurable and nucleation is primarily of the heterogeneous type (impurity catalyzed). According to Turnbull et al. [1] the crystal growth velocity varies inversely with the shear viscosity and is given by: U = fkJ-T-T111 3rra0z

e-Ag/k"r].

(1)

Here Ag is the free energy per molecule, T the temperature, " a 0 " a jump distance (on the order of a hard sphere diameter) and 71 is the shear viscosity. If all parameters in the above equation were held constant with the exception of the shear viscosity, then a viscosity change (increase) of one order of magnitude would change the crystal growth velocity by one order of magnitude (slow0022-3093/88/$03.50 © Elsevier Science Publishers B.V. (North-Holland Physics Publishing Division)

ing down). As an example SiO 2 (GE 106) at 1200 ° C has a shear viscosity of 5.5 X 1012 P while a nitrided borosilicate glass at the same temperature has a shear viscosity of 1.2 x 1013 P [2] and one may easily envision the effects on devitrification that a nitrided glass may have. Devitrification is an important limiting factor governing the maximum working temperature of fused silica and silica glasses for which these glasses remain stable to phase transformations. To increase these operating temperatures and to increase the devitrification resistance of fused silica and other silicate based glasses, atomic nitrogen is inserted in the glass. That atomic nitrogen insertion in a glass would be advantageous with respect to devitrification is shown by the work of other investigators on porous Vycor [2] and soda-limesilica glasses [3-5]. Their work shows that several weight per cent of nitrogen may be taken up by reconstructed high silica glasses at elevated temperatures and the viscosities of these nitrided glasses attain values greater than normal fused silica. In addition alkali ion impurity migration in glasses, thought to be a contributor to the devitrification process, may be reduced. Dietzel and

182

J. Schroeder, J.J. Jarek / Nitrogen doping of glasses

Wickert [3] have shown that the mobility of alkali ions is reduced in a nitrided glass.

and interface rearrangement, nSiO2(glass) + O(non-bridging) + OH kB+ kB

2. Theoretical background

n SiO 2 (crystalline) + O(non-bridging) + OH, (2b)

Spaepen and Turnbull [6] have proposed a new model for the regrowth of crystals in thin molten or amorphous solid layers of four-coordinated covalently bonded elemental systems, such as Ge and Si [6]. In this model the activated defect is a broken covalent bond, and once this bond is ruptured the two halves can move rapidly by a specific rearrangement process that involves a simultaneous bond formation and breaking. Breaking one bond suffices to alter the five and seven fold rings, representative of the amorphous state, to six membered tings which make up the crystal planes [6]. The dangling half bonds propagate along a crystal ledge by such a method until they are annihilated by recombination with each other or by another half bond. Spaepen and Turnbull [6] found that a single half bond was not sufficient. The cooperative motion of two halves of the broken bond was necessary. In fused silica, bond rupture makes two types of dangling half-bonds, a non-bridging oxygen and a three coordinated silicon [7], of which the non-bridging oxygen is the most mobile. Fratello et al. [7,8] propose a model for the mechanism of growth of quartz crystals into fused silica where the active defect is a non-bridging oxygen attached to the same silicon group as a hydroxyl group. The action of the hydroxyl group on the fourth bond of the silicon allows higher mobility for the non-bridging oxygen, since the base silicon is allowed greater movement than a network bond would permit. Also a reorientation of the structure is permitted by a hydrogen jump between the defects [8]. The above described mechanism is formulated quantitatively in the following way [7,8]. The crystal growth can be described as a two step process of creation and annihilation of defects. Si-O ~ Si(3 coordinated) + O(non-bridging) (2a)

with n being the number of SiO 2 molecules crystallized per activated defect. Here a non-bridging oxygen is the activated defect. From the interface rearrangements and application of chemical reaction rate theory [9] the rate of crystallization is calculated by Spaepen and Turnbull [6] and Fratello et al. [7,8] to have the following form,

u = A B e-~C~/RT(1 -- e"aC"/RT)C, oCoH,

(3)

with AG B being the Gibbs free energy of crystallization per mole of SiO 2 and AG~ is the standard molar free energy of activation [7]; Cno and COH are the concentration of non-bridging oxygen and hydroxil groups respectively. The concentration of non-bridging oxygen is determined by Fratello et al. [7,8] by assuming equilibrium in eq. (2) and the result is the following:

Cno= - ~ ÷ [(~2-b KAC, t)] 1/2,

(4)

where 8 is a parameter given by the degree of reduction, K A is an equilibrium constant and Cst is the concentration of strained bonds. The equilibrium constant becomes KA= Cnn(C3si/Cst ) = e - ~ ° , / R T with C3si being the concentration of three coordinated silicons and Gst is the free energy difference between a strained bond and the resulting defects [7,8]. The degree of reduction is determined by the stoichiometry of the specimen and for relation (2) has the form C3si = C,o + 2& A complete description of crystal growth rate is given in the following equation [7,8],

u = A B e-aC~ (1 - e ~aa"/RT)

×

+ [8 ÷ (e- s,J" Cs,)llJ2)Co.,

(5)

where A and B are constants and all other parameters have been defined above. The predictions in eq. (5) are consistent with various experimental observations on the devitrification kinetics of fused silica [1,10-12]. The greater the concentration of O H - the larger the crystal growth rate, whereas

J. Schroeder, J.J. Jarek / Nitrogen doping of glasses

the more a glass is in a reduced state (nonstoichiometric) the lower the crystal growth rate. Eq. (5) clearly shows that ( O H - ) may act catalytically in the crystal-glass interface to bring about increased devitrification, while impurity doping (nitrogen insertion) may certainly increase the resistance to crystallization in a glass former by the reducing action of the dopant on the base glass.

3. Experimentalprocedure Some of the nitriding was performed in a radio frequency (RF) heated high pressure high temperature furnace. In fig. 1, a schematic of a sample space is shown. Since ammonia gas (electronic grade) was the source of atomic nitrogen a tungsten susceptor and boron nitride cloth and felt were the high temperature insulator materials. The pressure furnace walls are stainless steel and are water cooled. The sample is contained in a crucible made of pyrolytic boron nitride and the crucible walls are lined with boron nitride cloth to

..----Meosurlng lllermocouple (W-Re)

Cylinder

Insulation

o

O

~

0

//~//~

0

"~~

D

PowderSable

0

R'F'COIIs(G°ld°nd

~ OJ ",~

Tefo l nCooe td)

B-NnisuoltnC ond O lh tFelt

Fig. 1. Schematic of sample space with R F heating and N H 3 atmosphere.

183

prevent the sample from "wetting" the crucible wall and causing a crucible failure due to a thermal expansion mismatch. The furnace design permitted static gas pressures up to 1500 psi and the temperature upper limit was determined by the crucible melting point of - 2200 ° C (the RF furnace was rated to attain 4000 ° C). Six types of samples were used in this study. SiO 2 powder of 53 /~m or 27 btm particle size, C o m i n g No. 7930 (porous Vycor) with a particle size of about 35 # m and a surface area of about 200 mZ/g; Cabosil (99.8% SiO2) with a surface area of about 400 m 2 / g and particle size about 70 A; mixtures of SiO2, MgO and Si3N 4 powders; cylinders of G E 214 fused silica and thin tubes (multivapor arc tubes) made from G E 214. A listing of all samples used for this study is given in table 1. Finely divided powder batches of SiO 2 or porous Vycor were exposed to a static nitrogen atmosphere (N 2) at various pressures and temperatures. The results showed that very little nitrogen was chemically bonded to the glass by these techniques (0.02%) and a number of the glasses fused from the powders showed strong evidence of devitrification. Some oxynitride glasses were prepared by melting mixtures of SiO 2, MgO and Si3N 4 in boron-nitride crucibles under a static nitrogen atmosphere. The analysis of these glasses by the micro Kjeldahl method showed as much as 3 wt% of nitrogen. X-ray diffraction also confirmed that we had a glass by the absence of the known diffraction lines of the crystal phase of the constituents and the existence of a broad peak characteristic of an amorphous sample. However, the glass obtained in this way was not transparent. Samples of SiO 2 • M g O - Si3N 4 at 1 m m thickness were just barely translucent. The next process was to take SiO 2 powders of various particle sizes and heat it to about 1000 ° while exposed to N H 3 and N 2 atmospheres. After a certain time duration the sample was quenched to room temperature. These nitrided powders where then packed and sealed in platinum cans and cohesive glasses were formed in a high-temperature, high-pressure press. U p o n cooling we found transparent glass containing a certain wt% of nitrogen chemically dissolved in the SiO2 host glass. The same process was repeated with porous

J. Schroeder, J.J. Jarek / Nitrogen doping of glasses

184

4. Results and discussion

Vycor powder or Cabosil powder. In every case we were able to insert atomic nitrogen into the host glass lattice by this NH 3 cracking procedure. We a c h i e v e d n i t r o g e n c o n c e n t r a t i o n s u p t o 7.1 w t % i n s o m e o f t h e s e glasses. I n t a b l e 1 t h e e x p e r i m e n t a l results are summarized with respect to method, nitrogen content in the glass and pressure, flow and temperature parameters. The nitrogen content

4.1. K i n e t i c studies N i t r i d i n g o f p o r o u s V y c o r a n d silica a n d silic a t e g l a s s e s as a f u n c t i o n o f n i t r i d i n g t e m p e r a t u r e , time and NH 3 to N 2 flow rates was done and the results showed that the amount of nitrogen that may be incorporated in the host glass structure s h o w s a m a x i m u m w i t h t e m p e r a t u r e a n d t i m e . See

i n all g l a s s e s w a s d e t e r m i n e d b y w e t c h e m i c a l analysis (micro-Kjeldahl method).

Table l a Summary of experimental parameters for nitrogen doped silica glasses Sample No.

Starting material

TN ( o C)

Time (h)

Gas atm.

Wt% N d)

Comments

1 2 3 4a 4b 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30a 30b 31 32

7930 a~

1200 1200 1750 1207 1207/2000 1800 1770 1770 1370 1500 1250 1620 2090 1500 1000 1000 1000 1470 1000 1000 1000 1000 900 1000 1200 800 600 1400 900 1100 500 500 700 1400

3.0 2.0 9.0 9.0 6.0 4.0 4.0 0.5 2.0 1.0 1.0 0.5 4.0 1.0 3.0 7.0 14 4.3 7.0 30.0 20.0 10.0 20.0 20.0 20.0 20.00 2.0 20.0 20.0 20.0 20.0 20.00 20.00

N2

0.01 0.016 0.008 0.020 2.97 Mo impure Mo impure 0.067 equipment failure 0.021 0.032 < 0.01 0.002 0.217 0.417 0.643 0.055 1.420 0.306 2.45 0.64 contaminated 1.245 to be analyzed 1.045 0.400 0.051 1.850 to be analyzed 0.290 0.172 0.851 0.156

physical solubility physical solubility complex glass physical solubility physical solubility complex glass Mo crucible melted Mo contamination evidence of chem. sohib.

a) b) c) d)

SiO2 b) S i O 2 • Si 3 N 4

SiO 2 SiO2 S i O 2 • Si 3 N 4 • M g O

SiO 2. Si 3 N 4 • M g O SiO 2 •Si 3N4. MgO SiO 2 SiO 2 SiO 2 SiO 2 SiO 2 7930 7930 7930 7930 7930 Cabosil c) Cabosil 7930 SiO2 7930 7930 7930 7930 7930 7930 7930 7930 7930 7930 7930 7930

Porous Vycor (Corning -#7930). SiO2_GE 204 (54 ~ particle size). Cabosil (400 m2/g, average particle size - 70 A). Determined by Micro-Kjeldahl Analysis.

<

N2

N2 N2 N2 N2 N2 N2 NH3/N NH3/N NHa/N NH3/N NHa/N N2 NH3/N NH3/N NH3/N NH3/N NH3/N NH3/N NH3/N NHa/N NH3/N NH3/N NH3/N NH3/N NH3/N

2 2 2 2 2 2 2 2 2 2 2 2 2 2 2 2 2 2

N2

NH3/N NH3/N NH3/N NH3/N NH3/N NHa/N

2 2 2 2 2 2

chemical solubility chemical solubility temperature too high control sample chemical solubility chemical solubility chemical solubility chemical solubility chemical solubility chemical solubility chemical solubility chemical solubility chemical solubility chemical solubility chemical solubility chemical solubility chemical solubility control sample chemical solubility chemical solubility chemical solubility chemical solubility chemical solubility chemical solubility

J. Schroeder,J.J. Jarek / Nitrogen doping of glasses fig. 2 w h e r e t h e a m o u n t o f n i t r o g e n c h e m i c a l l y d i s s o l v e d i n p o r o u s V y c o r a n d G E 2 1 4 g l a s s as a f u n c t i o n o f n i t r i d i n g t e m p e r a t u r e is p l o t t e d . I t a l s o appears that glasses containing some boron may b e a b l e t o t a k e u p m o r e n i t r o g e n t h a n p u r e S i O 2. To further treat the problem of chemical solubility of gases in glass a model for chemical solub i l i t y b y S h a c k e l f o r d et al. [13,14] w a s a p p l i e d t o

185

chemical bonding of the dissociation species with t h e h o s t glass. I n t h e m o d e l o f S h a c k e l f o r d et al. it is a s s u m e d t h a t all t h e d i s s o l v e d s p e c i e s a r e l o c a l i z e d a n d t h a t is a d i r e c t c o n s e q u e n c e o f t h e r i g i d ity requirement, the glasses must be below their glass transition temperatures. The Shackelford, S t u d t a n d F u l r a t h [13,14] e x p r e s s i o n f o r c h e m i c a l solubility has the form:

N2, H 2 a n d N H 3 i n S i O 2 glass. T h e b a s i c m o d e l o f S h a c k e l f o r d a n d c o w o r k e r s is v a l i d i n g l a s s e s below their glass transition temperature. This model assumes an ideal gas in equilibrium with t h e s p e c i e s i n s o l u t i o n . T h e i d e a l g a s is r e p r e sented by free particles in a three dimensional box with any molecular motion (rotation and vibration) given by independent rigid rotors and simple harmonic oscillators, respectively. Chemical solubility infers molecular dissociation and very strong

×[(1 --e--Or'/T)) ×{2Ns[ e

1/2

O,/2T/(l_ e

o,/r)]3}

Xe%,/Zk.T¢ E(O)/RT

Table lb Sample

Material

TN ( o C)

1-S 2-S 3-S 33 34a 34b

SiO 2 SiO 2 GE 214 GE 214 GE 214

1100 1100 1100 1000 1000 1000

35-T 36-T 37 38 39 40 41 42 43a 43b 44 45a 46 47 48 49 50 51 52 53

GE 214 GE 214 GE 214 GE 214 GE 214 GE 214 GE 214 GE 214 GE 214 GE 214 GE 214 GE 214 GE 215 Cabosil GE 214 GE 214 GE 214 GE 214 GE 214 GE 214

1000 800 900 1000 1000 600 800 700 900 900 1000 1000 1000 1000 600 700 800 900 1000 1000

31 20 20 97.25 50 20 20 20 20 20 30 30 30 30 20 20 20 20 20 209.67

54

Ge 214

1000

209.0

SiO 2

Time (h) 20 20 10 10 20 20

Particle size (/~)

Wt% N

Comments

cylinder cylinder cylinder 37 37 37

0.374 0.564 0.568

solid cylinder solid cylinder 37 cylinder cylinder 37 37 37 37 37 4 37 54 0.007 37 37 37 37 37 arc-tube and cylinders arc-tube and cylinders

0.007 0.503 0.146 0.168 0.288 0.359 0.259 0.487 0.872 0.607 7.17 0.073 0.158 0.301 0.425 0.801 -

coated with # 7930 HF-etched analyzed 50 h in vac. furnace-analyzed devitrification study devitrification study Used for interfer, micro. Used for interfer, micro. -

-

-

-

(6)

J. Schroeder, J.J. Jarek / Nitrogen doping of glasses

186 Table lc Sample

Starting material

TN ( o C)

55 56 57 58 59 60 61 62 63 64 65 66

1000 1000 1000 1000 1000 1000 1000 1000 1000 1000 1000 1000

68 69 70 71 72 73 74

GE214(37/z) GE214(37 #) GE214(37 ~) GE214(37 ~t) GE214(37/z) GE214(37/z) GE214(37/~) GE214(37/L) GE214(37/L) GE214(37/L) GE214(37 la) GE214(cube) ( 1 2 × 1 2 × 3 0 m m 3) GE214(cube) ( 1 2 × 1 2 × 3 0 m m 3) GE214A(37/~) GE214A(37/~) GE214A(37/~) GE214A(37/~) GE214A(37/~) GE214A(37 #) GE214A(37/~)

75 76 77 78 79 80 81

GE214A(37/z) GE214A(37 ~) GE214A(37 ~) GE214A(37 ~t) GE214A(37/.t) GE214A(37/~) SiO 2 powder

67

Time (h)

Gas atmosphere

Wt% nitrogen

Comments

20 1 4 9 16 25 1 25 4 2 1 50

NHa/N 2 60/20 60/20 60/20 60/20 60/20 60/20 60/20 60/20 60/20 60/20 60/20

0.180 0.188 0.094 0.256 0.260 0.343 0.066 0.994 0.684 0.579 0.405 -

W paddles as stirrer W paddles as stirrer W paddles as stirrer W paddles as stirrer W paddles as stirrer W paddles as stirrer W paddles as stirrer no stirrer no stirrer no stirrer no stirrer bulk solid of SiO e

1000

51

60/20

-

bulk solid of SiO 2

1000 1000 1000 1000 1000 1000 1000

1 4 25 9 16 49 100

60/20 60/20 60/20 60/20 60/20 60/20 60/20

-

1000 1000 800 800 800 800 800

73.25 33.25 16 24 49 9 100

60/20 60/20 60/20 60/20 60/20 60/20 60/20

0.29/0.31 0.51/0.56 0.91/0.94 0.72/076 0.76/0.67 0.99/0.98 (contaminated) with B-N cloth 0.998 0.965 0.280 0.294 0.470 0.210 0.52

-

3.C

2.4

1.8

/ / / /

u

=o

,,

/

1.2

~.\

•

/

0.6

\ \

/

I--

Z O.C 500

7~X)

NITRIDING

9~

L

11(X)

e

13OO

~

1500

TEMPERATURE (°C }

Fig. 2. Nitrogen concentration versus nitriding temperature for G E 214 fused silica and Corning-7930 Vycor glass with a total nitriding time of 20 hours. • - is the symbol for GE-214A (SiO2) and • - is C o m i n g # 7 9 3 0 (Vycor).

J. Schroeder, J.J. Jarek / Nitrogen doping of glasses

where Or= h2/8~r2kBI is the characteristic temperature for rotation of the polyatomic molecule; 0vi is the characteristic temperature for vibration of the molecule's internal vibration mode; 0v is the characteristic temperature for vibration of thermoatomic species; %1 is the negative of dissociation energy of the diatomic molecule and E(0) is the binding energy between monatomic species and glass. N~ is the number of solubility sites per cm 3, n~ the number of gas a t o m s / c m 3, p the pressure of the gas atmosphere, m, the mass of one molecule, h, k B and R the Planck Boltzmann and gas constants, respectively, and T the absolute temperature. Values for Or, 0~ and (el of the H 2, N 2 and N H 3 molecules were all obtained from the J A N A F table [15], Herzfeld and Litovitz [16] and Stuart [17]; N~, the site density was calculated to be 2.239 × 1022 sites/cm3; 0v values were calculated from I R absorption peaks (at 2.94 /~m), this leaves only E(0), the binding energy of the dissociated species to the silica structure and these could be estimated for H 2 from the data of Shackelford et al. [13], for N 2 and N H 3 from Stuart [17] and Herzfeld and Litovitz [16]. In fig. 3 we plot the solubility of H or N a t o m s / c m 3 of glass divided by the square root of pressure as a function of inverse temperature for hydrogen, nitrogen and N H 3 gas and compare it to the measured results for porous Vycor and SiO 2 glass. It is immediately evident that none of the calcula-

iOz2 _ _

_

r

i

T

i

i

i

3

187

tions coincide exactly with the measured results, however, the calculated model with N H 3 does show some agreement. This shows that the N H 3 cracking technique may be the most efficient process. N 2 dissociation at high temperature might eventually work but the temperatures required would be such that the process would be technically difficult to attain• One other surprising fact that is nicely demonstrated is the amount of gas that can chemically dissolve in a glass at a relatively low temperature of 500 o C.

4.2. Structural considerations A question of major consequence unanswered up to now is how does the nitrogen bond to the SiO 4 tetrahedron. Possibilities of dangling bonds for existing defect states or interstitial trapping exist. Both of these models would provide us with a value for the m a x i m u m amount of nitrogen that is dissolved chemically by the following scheme for the a m m o n i a cracking process [4,5,18].

I I. - Si - N 2- (comparable to non-bridging J oxygen)

I -

I

Si -O

I

I

II. - Si - N - - Si - (comparable to a bridging ] ] oxygen)

I

~ - - - ~ -

-Si

I N E iO~s --

iO~2 ~ . ~ . - ~ - - - - ~ r ~ -- ~ r --A . . . . .

-A----______ iOlO

~ i0'6 _

"~.\

I

III. - S i - N - Si - (not found in oxide glass)•

I

I -

I 106 ~"

• - H

~"

I 04

2~ :SiO 2 (CALCULATED}

i 06

t 05

F i g . 3. T h e c h e m i c a l inverse

Si-

I

xq~ "-.

- - )04

o -NH 3, ~02 ~ (MEASURED) v -NR$ ~SiO2 (MEASURED} 05

I

I

- - 108

0 -NH:"1

I

O-Si-

treatment

I 07

solubility

temperature.

~

-- i02 \~

i l 0.8 09 IO~/(TN) OK-L of various

I I0

I II

gas atoms

Both calculated

values

I "I~ 12

I0o

5

versus

the

from

eq.

(6) and measured values for nitrogen in SiO2 glass are shown.

The above reactions demand a low partial pressure of water vapor and shift to the right as this partial pressure is decreased. If the partial pressure of water vapor reaches zero then the nitrogen will be incorporated exclusively as dissolved nitrides [4]. Hence, for mode III, the only one not possible for the oxide glass, the degree of linkage is above that of pure oxide glasses• With this type

J. Schroeder, J.J. Jarek / Nitrogen doping of glasses

188

of bonding a much more rigid structure arises with rather different physical properties.

4.3. Optical studies A Zeiss interferometer microscope was used to try to find the depth of nitrogen penetration in bulk glass samples, however, the results are inconclusive up to now due to surface roughness, surface strains and the lack of good optical flatness of the bulk samples. Observations have been made by examining the ultra violet (UV) transmission spectrum of nitrided glass and comparing it with untreated fused silica. Fig. 4 gives the results of these studies. Upon nitriding the UV absorption edge of the glass does move to longer wavelengths. For untreated G E 214 the UV absorption edge is around 170 nm [19,20], whereas the nitrided sample shows an absorption edge starting around 200 nm. The region of most interest in fig. 4 is in the rectangular area. At 200 nm, G E 214 transmits about 32% of the incoming UV while the nitrided G E 214 transmits only about 4% of the UV light. Various authors [21-24] have observed that upon introducing a conventional modifier O H into fused silica the tail of the intrinsic UV absorption edge shifts to longer wavelengths when compared to the unaltered fused silica. It is gener80

I

r

l

I

I

GE 214

2O

/

l

t40

200

I

_

II

~ /

0

I

/)K'x

I I

$00

400 WAVELENGTH- ), (rim)

500

600

Fig. 4. Percentage transmission of nitrided G E 214 SiO 2 with respect to wavelength is compared to non-nitrided G E 214 SiO 2 glass. Note the red shift for the nitrided sample, the ultra violet transmission of the nitrided sample is reduced at the shorter wavelengths.

ally agreed that the location of the ultraviolet edge of oxide glass may be discussed on the basis of the strength of S i - O - S i bridging in the network [25,26]. Hence, the fewer non-bridging oxygens in the network the stronger the bridging and the greater the UV transmission. Adding a modifier or O H - does increase the number of non-bridging oxygens, weakening the S i - O - S i bond, with the result that the UV transmission is decreased and the absorption band edge moves to longer wavelengths [22]. A similar effect is observed when UV transmission measurements are performed nitrogen doped G E 214 fused silica. These transmission measurements (see fig. 4) on the nitrogen doped G E 214 glass compared to normal G E 214, estabfished that the UV absorption edge shifted to the longer wavelengths for the doped sample compared to the untreated fused silica. These observations may indicate that in the nitriding process some of the bridging oxygens are replaced by nitrogen in the G E 214 glass. What the exact nature of the structure of the nitrogen incorporation in the glass is, has not definitively been established and groups such as =NH, = N and O N - are also likely candidates. Whatever the resulting structure may be, one thing is certain it must be of the type that allows the whole lattice structure to be stiffened. Consider the work of Philipp [27,28] on various fused silicas and silicon nitrides in which their optical properties were explored. Especially important are measurements of the absorption curves of any SiOx (x < 2) shifted to lower energies (longer wavelengths) likewise any siO~Ny or SiN z had absorption curves that were shifted to longer wavelengths when compared to stoichiometric fused silica. The silicon nitride structure may be ruled out since its absorption curve would have to move to longer wavelengths than silicon-oxynitride. In fig. 3 we also plot two points, one for SiO0.96N0.7o, the other for SIN4/3, with the data taken from the paper of Philipp [27]. It is immediately obvious that our nitrided samples of fused silica are in between these two values for the silicon-oxynitride and silicon nitride. It seems that for our nitrided glass more oxygen bonds are broken than exist in a SiOo.96No.70 glass but not so many that we reach the SIN4/3 case. The above

J. Schroeder, J.J. Jarek / Nitrogen doping of glasses Table 2 Fluorescence study results a) Excitation wavelength

Filter (wavelength pass)

5145 ,~ 5145 ~, 5145 A

>/5100 ~. >1 5250 A >~ 5820 ~.

Fluorescence GE 214 green yellow-orange yellow-orange red

Nitrided Ge 214 blue-green faint green none

a) Conclusions: (1) Nitrided sample fluoresces weakly ( < 5820 A), localized at - 5400 ,~. (2) Untreated sample fluoresces ( > 5100, > 5820), localized at - 5600 ,~ and - 6200 ,~..

observations could be employed in determining the thickness of nitrogen penetration in a bulk sample. Normally produced fused silica is in a chemically reduced state. At about 2410 ]~ the normal G E 214 (fused silica) absorbs strongly in the ultraviolet and this absorbed energy is reemitted as visible fluorescence. This absorption and consequent reemission as visible fluorescence can be regarded as a measure of the degree of reduction [23]. The nitrided fused silica shows even higher absorption and consequently more fluorescence in the visible than the non-nitrided fused silica. Table 2 shows this result. The near infrared spectra for the same set of samples confirms that nitriding tends to remove water ( O H - ) at the same time that nitrogen is inserted into the lattice. The near infrared spectrum also contains a weak line at 2.98 g m that is indicative that some N H groups are formed in the nitrided glass. That the nitrided glass goes even further from stoichiometry is also evident in the shift of the ultraviolet band edge to a higher wavelength as nitrogen is applied [28]. Surface nitriding and bulk nitriding was carried out in this study. The philosophy behind the surface nitriding was that most of the devitrification is of the heterogeneous type and thus originates at the surface of the sample. By preventing or slowing down nucleation at the surface the major aspect of devitrification could be eliminated. The samples of G E 214 that were nitrided at various temperatures and for various time durations showed interesting I R and ultraviolet spectral response. Fig. 5 shows that heating a sample

189

for many hours at a specified temperature in a N H 3 atmosphere greatly reduced the amount of dissolved ( O H - ) in the G E 214 glass. This was determined by measuring the strength of the ( O H - ) stretch band at 2.73 g m and calculating the amount of ( O H - ) from the extent of this band. Actually whether the atmosphere was nitrogen or ammonia, in either case the amount of dissolved ( O H - ) in the glass was severely reduced from the initial starting value. The ultraviolet transmission spectra (from 260 nm to 185 nm) also showed interesting results. The absorption fine centered at about 240 nm increased in strength as the glass was heat treated. The longer the heat treatment the greater the absorption. The UV absorption edge does shift to higher wavelengths as the heat treatment time is increased. This absorption can also be regarded as a measure of the degree of reduction [23]. Hence, both the nitriding and the heat treatment process produces nonbridging oxygens and removes ( O H - ) from the glass and the UV spectra shows that the fused silica is reduced (made more non-stoichiometric) in the process. Both aspects are a desirable occurrence as predicted from eq. (5) with regard to devitrification. A third process also occurs in the cross linking effect brought about by the nitrogen incorporation in the lattice of the glass. Nitrogen can replace oxygen in the crystalline silicates, hence it is feasible that a similar replacement may occur in vitreous sihcates [29]. It is thought that even small concentrations of nitrogen (on the order of - 1 A / O ) in oxide glass tend to increase the softening temperature and viscosity [30]. The mechanism for this may very well be that in the tetrahedral network the nitrogen is coordinated by three linkages and the structure should be more rigid and consequently have a higher viscosity than non-nitrided silicate glasses. This nitrogen incorporation may allow the production of bulk glasses that are more refactory and more resistant to devitrification and therefore more stable than ordinary fused silica. A method of determining whether nitriding stiffens the lattice of SiO 2 glass producing a more rigid structure and increasing the viscosity of the glass is by Brillouin scattering measurements. Extremely small samples can be used with this tech-

190

J. Schroeder, J.J. Jarek / Nitrogen doping of glasses

nique. One measures the Brillouin shifts of each sample and determines in what way the shifts go as a function of changing the various parameters of interest. For this study bulk nitriding was performed on finely divided powders of SiO 2 in the form of Cabosil (particle size about 70 ,~) and then the Brillouin spectra were obtained for sintered samples. A typical nitrided sample of Cabosil containing 7.17 wt% of nitrogen has a longitudinal Brillouin shift of 28.3 G H z while the non-nitrided Cabosil had a Brillouin shift of 21.79 GHz. At 4880 ,~ fused silica has a longitudinal Brillouin shift of 25.20 GHz. The normal Cabosil sample did contain some ( O H - ) hence it is reasonable that it has a value of 21.70 GHz. It is most interesting that the nitrided sample had a value of 28.3 GHz, a definite indication that the lattice has been stiffened by the nitrogen addition. The Brillouin shifts are essentially proportional to the refractive index multiplied by the sound velocity of the sample [31]. Hence, nitriding may effect both of the quantities but only if the lattice is stiffened in the process. Actually the incorporation of nitrogen into the host silica network seems to tighten the network structure, which in turn increases the strength of the ionic restoring forces resulting in an increased sound velocity. It may be of interest to add that small amounts of chemically dissolved nitrogen seems to effect the sintering properties of SiO2 powders. The higher the nitrogen content of the particles to be sintered the higher temperature required for proper sintering. This indicates that the softening point is affected by the nitriding. The changes in sintering behavior were observed in nitrided G E 214 glass particles of 37 /~m size and in nitrided Cabosil powder of 0.007 /~m diameter. When nitrided Cabosil and untreated Cabosil were held in a furnace under a nitrogen atmosphere and at -1250°C, the untreated Cabosil sintered to a cohesive glass but also showed evidence of severe devitrification while the nitrided Cabosil powder did not sinter nor showed any visual evidence of crystallization. The change in the sintering response of several nitrided powders seems to be direct evidence that the nitriding process does affect the softening temperature and consequently the viscosity of the glass. The resistance to sinter-

ing of the nitrogen doped powders at temperatures where normal fused silica powders will sinter is certainly an indication that nitriding raises the viscosity of the doped glass. In addition, the fluorescence spectra of nitrided and non-nitrided G E 214 was compared and the results are summarized in table 2. Clearly nitriding changes the fluorescent response of a glass, but at this stage it is premature to draw too many conclusions on the scant amount of evidence presently available.

4.4. Surface nitriding G E 214 fused silica cylinders have been nitrided by the high temperature ammonia cracking process. The micro-Kjeldahl analysis, that worked well for the nitrogen analysis of finely divided powders where the nitrogen content was distributed uniformly, can no longer give a quantitative determination of concentration for the solid samples in which the nitrogen is taken up in an as yet undefined diffusion layer. However, qualitatively the Kjeldahl method is still an indicator of nitrogen in a glass. For example, sample 35-T in table 1 shows about 0.007 wt% of nitrogen which is about seven times the lowest limit of detectability of the analysis method. The kinetic studies on the powder samples showed that the nitrogen concentration does increase linearly with the square root of the nitriding time (see figs. 5 and 6). This indicates that the nitriding process in our samples is diffusion controlled, and in the bulk samples (thin flat plates, cylinders, etc) only thin surface layers are nitrogen doped by this process. This may still not be a severe limitation to the solution of our problem since the nucleation that will eventually lead to devitrification is primarily of the heterogeneous, surface type (impurity catalyzed) and nitriding only thin surface layers of a sample may suffice. In fig. 5 the amount of nitrogen in wt% is plotted as a function of the square root of nitriding time and the nitriding temperatures of 800 ° C and 1000 o C. The amount of nitrogen that can be inserted initially into the G E 214 fused silica glass goes linearly with the square root of time. The kinetics of the nitriding process are also evident

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~- (hrs) Fig. 5. (a) Wt% nitrogen versus the s q u a r e root of n i t r i d i n g time for G e 214 SiO 2 at two n i t r i d i n g temperatures. • a n d O d e n o t e T h = ( 1 0 0 0 ° C ) a n d • is the s y m b o l for T N = 8 0 0 e C . All s a m p l e s u s e d are particles 37 ~ m in d i a m e t e r c o n t a i n i n g < 1 5 P P M of O H . (b) N i t r o g e n c o n c e n t r a t i o n in wt% with respect to the s q u a r e root of the total n i t r i d i n g t i m e for three different silicate glasses at a n i t r i d i n g t e m p e r a t u r e of 1000 o C. [] - is C a b o s i l of 0.007/~m p a r t i c l e size c o n t a i n i n g a b o u t 103 P P M of O H - ; • - is C o r n i n g :#:7930 (Vycor) w i t h u n k n o w n O H - c o n t e n t ; a n d • - d e n o t e s G e - 2 1 4 A (SiO2) of particle size 37 tLm c o n t a i n i n g less t h a n 15 P P M of

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since the higher the temperature the more nitrogen is chemically bonded to the glass. A 200 ° C rise in the nitriding temperature gives a three fold increase in the amount of atomic nitrogen that is retained in the glass. The nitrogen front seems to propagate into the glass with a square root time dependence which of course is normal for a diffusion controlled process. In fig. 5 the 1000 ° C plot shows a linear dependence up to about 36 h of nitriding time, beyond this time the amount of dissolved atomic nitrogen is constant with time. To explain this observed behavior we shall first discuss the diffusion controlled reaction model as given by Shelby and co-workers [33,34]. One assumes that the interaction between the diffusing gas and the glass is very fast compared to the diffusion rate. Reaction site concentrations are independent of temperature and are constant throughout the material. Sharp boundaries will exist between the reacted and non reacted zones. This is evident from our secondary ion mass spec-

troscopy (SIMS) work. The concentration of nitrogen at the gas-solid interface will be determined by the chemical solubility of the nitrogen in the material and by the applied gas pressure. It can be shown that the concentration of reacted material is given by [33,34] (C(Cf-

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J. Schroeder, J.J. Jarek /Nitrogen doping of glasses

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If all the particles that we are nitriding have the same distribution of diameter then even L may be eliminated and the concentration of nitrogen in the glass should be proportional to the v~-. This is indeed the case as is evident from figs. 5 and 6, up to a certain nitriding time. At a nitriding temperature of 1000°C for Ge 214 (with most of the water removed) the amount of nitrogen chemically dissolved in the glass is linearly proportional to the square root of the time up to about 36 h of nitriding time, beyond this time duration the amount of nitrogen does not increase with increasing time. This seems to indicate a saturation effect, namely, each 37/~m particle has taken on as much nitrogen as possible. At 800 ° C the rate of nitriding is again linearly proportional to the square root of time but with lesser slope and saturation has not been reached even at about 50 h of nitriding time. The amount of dissolved nitrogen versus v~plots are interesting when the amount of ( O H - ) in

the G E 214 starting sample is a variable. Fig. 6 shows the effect that dissolved ( O H - ) has on the nitriding procedure. It should be noted that our basic starting material G E 214 contained - 2 0 0 ppm of ( O H - ) as determined from IR measurements. The same material after heat treatment at 1000°C for 50 h in a pure nitrogen atmosphere showed a reduced water content of < 15 ppm. The Ge 214 powder with 200 ppm of ( O H - ) had a much slower rate of nitriding than the Ge 214 powder that had most of its water removed. Low ( O H - ) containing fused silica powder seemed to saturate at about 1 wt% of nitrogen after about 30 h of nitriding while for the same time duration the glass containing about 200 ppm of ( O H - ) accumulated only about 0.38 wt% of nitrogen. The cause of this behavior may be that the additional water occupies some of the possible defect sites (reaction sites) and the nitriding procedure first removes the water and then reacts with the reaction sites to form the nitrogen containing bonds, all of which takes much longer than the case where most of the water has been removed beforehand. It should also be noted that the nitriding seems to progress best at temperatures corresponding approximately to the glass transition temperature (Tg) of each sample type and attempting to nitride much above Tg is a very inefficient process, if not impossible. Some of the nitriding runs done on SiO 2 and Coming 7930 powders at temperatures above 1400°C support this conclusion. At the higher temperatures the reaction between the diffusing nitrogen ( N - ) and the glass is no longer fast compared to the diffusion rates, and the concentration of the reaction sites may become temperature dependent [33]. The species to react may now leave the reactant site without being involved in a complete reaction, with the consequence that nitriding is no longer efficient and at high enough temperatures, nitriding ceases completely. Structures in the bulk form of thin flat plates, cylinders, tubes and spheres were nitrided at various temperatures and for varying nitriding times. Table 1 summarizes these samples with an exact description of these nitrided samples. Secondary ion mass spectroscopy (SIMS) was used to measure the diffusion of nitrogen in the fused silica.

J. Schroeder, J.J. Jarek / Nitrogen doping of glasses SIMS ANALYSIS OF NITROGEN DIFFUSION IN NITRIDED SILICA IO0 i

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Fig. 7 shows the compositional profiles of a nitrogen doped sample and a control sample consisting of an undoped fused silica plate. From a SIMS in-depth profile for nitrogen and silicon content as a function of sputter time we determine the (N3/Si) mass ratio. This is plotted in fig. 7 versus relative sputter time (proportional to the penetration depth). A thin gold film had been put over the sample surface prior to a SIMS run mainly to prevent any charge buildup on the non-conducting sample and this can be seen in fig. 7 up to a relative sputter time of 20. Beyond this penetration depth the nitrogen content is quite evident in the treated sample and missing in the untreated sample. Beyond a sputtering time of 80 very little nitrogen seems to have penetrated the treated sample and the (N3/Si) mass ratio values of the treated sample and control sample are about the same. Different portions of the same starting fused silica (GE 214) were nitrided at temperatures of 600 o C, 800 o C, 1000 o C and

193

1200°C for various time intervals. These treated samples were then analyzed for the SIMS profiles. From the SIMS profile measurements it is possible to obtain reliable diffusion coefficients for nitrogen diffusing into fused silica. A typical value for nitrogen at 1000 ° C in fused silica in D m. . . . = 4.7 × 10-14 c m 2 / / S . Calculating the self diffusion coefficient from a jump model for fused silica one finds Dj~mp ~ ( l ' 2 / 6 ~ ) = (l *2/6~s)C44-, 1.16 × 10 -2° c m / s , where 1 is the average jump distance, "9 approximately the shear relaxation time, ~s the shear viscosity and G4 the transverse elastic constant. A measured value for molecular nitrogen diffusing into glass at 1100 ° C is D -- 6.5 × 10 8 cmZ/s. We see that our measured value for atomic nitrogen falls just about into the middle of both extreme cases. The quality of the SIMS profiles was such that the diffusion coefficients could not be determined better than having error bars to the extent of at least one order of magnitude. The reason for such inaccuracy was primarily the establishment of the nitrogen zero baseline and the determination of where the gold film ended and the nitrogen profile began. Another problem is that SIMS only looked at (N 3) of mass 42 and not N of mass 14.

5. Conclusion and summary

It is possible to synthesize simple Si-oxynitride glasses. We have incorporated relatively large amounts of nitrogen into silica glasses and have obtained as a final product a durable and highly transparent glass. A kinetic study of the nitriding of silicate glasses has shown that the chemical solubility of nitrogen occurs even at temperatures as low as 500 o C, although higher efficiencies in the process are achieved at higher temperatures which must remain near or below the glass transition temperature for each glass. Hence, the kinetic study showed the existence of a nitriding temperature maximum and a N H 3 / N 2 ratio maximum for the most favorable nitriding conditions, but also that it is quite feasible to nitride simple glasses at relatively low temperatures. A number of realistic structural models for nitrided amorphous SiO 2 have been suggested, but none have rigorously

194

J. Schroeder, J.J. Jarek /Nitrogen doping of glasses

been established as the right one for the process. One needs to shed more light on these models, and determine by a combination of spectroscopic techniques (i.e., Raman, Brillouin and IR) which is most plausible (i.e., SiNH 2, SiNHSi (a broken Si-Si bond with an N H group inserted), SiNSi (bridge form), nitrogen as an interstitial atom or perhaps an SiON structure). The devitrification aspects of SiO 2 glass will be treated by more systematic experiments and we shall make measurements whether nitriding a glass will substantially change its devitrification characteristics at elevated temperature. Since nitrogen is substituted for oxygen in the glass network, thus producing a more highly cross-linked structure, the viscosity or annealing point should be effected by the degree of nitriding. For this purpose a beam bending type viscometer will be used to measure the viscosity as a function of temperature of the nitrided sample glasses. Devitrification studies of bulk SiO 2 glasses and nitrided glasses synthesized from the SiO 2 powder and N H 3 cracking process have been initiated. At the moment the number of samples is still too few to draw any conclusions with regard to the rate of increase or decrease of devitrification of nitrided versus non-nitrided glasses. These studies if continued should be most conclusive in providing answers to the general devitrification process in glass. The degree of devitrification will be determined by means of light scattering (Rayleigh) and should provide evidence of devitrification long before the onset of devitrification can be detected by other optical techniques. In Rayleigh scattering we shall measure depolarization ratios and relate these to the anisotropy of the glass brought about by the existence of precursors to devitrification. To summarize it may be emphasized that we have achieved the nitriding of a simple glass, namely, fused silica and this is a new result. The kinetic study that was done on porous Vycor, fused silica and Cabosil powders has shown a rather surprising result that even as low as 500 ° C some chemical dissolving of nitrogen may take place in a glass and large amounts of nitrogen are incorporated into the host glass network. By a simple N H 3 cracking technique at elevated

pressures and temperature, we are able to incorporate trivalent nitrogen into the SiO 2 structure. The trivalent nitrogen is substituted for oxygen in the glass network, consequently producing a more highly cross-linked structure, which is harder, more refractory and less soluble than if no nitrogen were present. Bond angle changes may dominate the cooperative reorientation of the basic structural units in silicates. Where bond angle changes are restrained by symmetry, coordination or packing a much more rigid structure arises with rather different physical properties (i.e. very flexible SiO2 versus partially flexible nitrided SiO2). Thus it is possible to produce nitrided bulk glasses that have a higher melting point and are more resistant to devitrification than SiO 2. Nitrogen is chemically bonded to silicon in the glass network and the formation of bridging network nitrogens is a general phenomenon applicable to many glasses. It has been shown for the first time that nitriding a finely divided silica sample and then sintering it into a bulk sample, that the nitrided sample supports a much higher sound velocity than the untreated fused silica (shown by Brillouin scattering experiments). This can only be interpreted to suggest that the nitrided sample must have a stiffer lattice than the untreated sample and this leads to the conclusion that nitriding tends to raise the viscosity of the treated glass relative to the untreated glass. This research was supported by the National Science Foundation, Materials Research Group, under Grant Number DMR-85-10617 entitled "Stability of Glasses".

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J. Schroeder, J.J. Jarek / Nitrogen doping of glasses [6] F. Spaepen and D. Turnbull, in: Laser-Solid Interactions and Laser Processing (AlP Publ. (50), 1978) p. 78. [7] V.J. Fratello, J.F. Hays and D. Turnbull, J. Appl. Phys. 51 (1980) 4718. [8] V.J. Fratello et al., J. Appl. Phys. 51(12) (1980) 6160. [9] S. Glasstone, K.J. Laidles and H. Eyring, The Theory of Rate Processes (McGraw-Hill, New York, 1941). [10] F.E. Wagstaff and K.J. Richards, J. Am. Ceram. Soc. 49 (1966) 118. [11] F.E. Wagstaff and K.J. Richards, J. Am. Ceram. Soc. 48 (1965) 38. [12] F.E. Wagstaff, S.D. Brown and I.B. Cutler, Phys. Chem. Glasses 5 (1964) 76. [13] J.F. Shackelford, P.L. Studt and R.M. Fulrath, J. Appl. Phys. 43 (1972) 1619. [14] P.L. Studt, J.F. Shackelford and R.M. Fulrath, J. Appl. Phys. 41 (1970) 2777. [15] JANAF - Thermochemical Tables (NSRDS-NBS 37) Second Ed. (US Govt. Printing Office, Washington, DC, 1971). [16] K.F. Herzfeld and T.A. Litovitz, Absorption and Dispersion of Ultrasonic Waves (Academic Press, New York, 1959). [17] H.A. Stuart, Die Struktur des Freien Molekiils (Springer Verlag, Berlin, 1952). [18] T. Kelen and H.O. Mulfinger, Glastechn. Ber. 41 (1968) 230.

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[19] E. Kordes, Glastech Ber. 38(6) (1965) 242. [20] B.D. McSwain et al., Phys. Chem. Glasses 4(1) (1963) 1: R.D. Maurer, J. Appl. Phys. 33(6) (1962) 2132. [21] G.N. Greaves, J. Non-Cryst. Solids 32 (1979) 295. [22] G.H. Siegel Jr, J. Phys. Chem. Solids 32 (1971) 2373. [23] T. Bell, G. Hetherington and K.H. Jack, Phys. Chem. Glasses 3 (1962) 141. [24] G. Hetherington and K.H. Jack, Phys. Chem. Glasses 3 (1962) 129. [25] H. Scholze, Glastechn. Ber. 32 (1959) 81. [26] E. Kordes and E. Worster, Glastechn. Ber. 32 (1959) 267. [27] H.R. Philipp, J. Electrochem. Soc. 120(2) (1973) 295. [28] H.R. Philipp, J. Phys. Chem. Solids 32 (1971) 1935. [29] K.H. Jack, J. Mat. Sci. 11 (1976) 1135. [30] K.H. Jack, in: Nitrogen Ceramics, ed. F.L. Riley (Noordhoff, Leyden, 1977) p. 257. [31] J. Schroeder, J. Non-Cryst. Solids 40 (1980) 549. [32] J. Schroeder, Light Scattering of Glass; Treatise on Materials Science and Technology, Vol. 12: Glass I, eds. R.H. Doremus and M. Tomozawa (Academic Press, New York, 1977) p. 172. [33] J.E. Shelby, J. Appl. Phys. 51(5) (1980) 2584. [34] J.E. Shelby and J. Vitko Jr, J. Non-Cryst. Solids 45 (1981) 83.