CHAPTER 2

LITERATURE REVIEW

2.1 Fe-Ni-Co alloy

The Fe-Ni-Co alloy was developed from the Fe-Ni or the Invar alloy. Charles

Edouard Guillaume discovered the Invar alloy in the 1890s. For his work on the Fe-

Ni system and the discovery of the Invar alloy, he was awarded the Nobel Prize in

Physics in 1920. The Invar alloy is the forerunner of a family of controlled expansion

alloys and the Invar alloy itself is still used today in vast numbers of household

appliances, computer terminals, TV screens, cathode ray tube, advanced electronic

components, filter mobile phone networks, telecommunication, aerospace

engineering, cryogenic engineering, which require either high dimensional stability or

expansion characteristics with variation in temperature. The diversity of these

requirements therefore led to the development of a wide range of Fe-Ni alloys in two

major groups; (i) low expansion alloys, the Invar and the N42 alloys, widely used in

the manufacture of electronic components in the integrated circuits and (ii) sealing

alloys such as the Fe-Ni-Cr and the Fe-Ni-Co alloys, which have been produced for

optical parts in a wide range of temperatures especially to associate with specific

glasses.

All Fe-Ni alloys consist of an austenitic phase. Among this system, the Fe-

Ni-Co alloy, also known as the Sealing Alloy, the Kovar alloy or the Alloy-F15,

belongs to the special thermal expansion alloy group (Huang et al., 1999). A third or

further elements are added to improve their physical, mechanical and chemical

5

properties, depending upon the purpose of the application (Shiga, 1996). Therefore,

the thermal expansion coefficient of the Kovar alloy depends strongly on the relative

amounts of Ni and Co. The addition of Co in the Fe-Ni system was first reported by

H. Scott in 1930. The Co addition provided the alloy with low thermal expansion and

also the desirable higher transition temperature. The Fe-Ni-Co alloy is widely used in

forming glass-to-metal, vacuum-tight seals especially for joining with borosilicate

glasses in hermetic seals (Ikeda and Sameshima, 1964). Apart from its thermal

expansion, the main advantages of this alloy are its machineability and the possibility

to solder for discrete translators, diodes and integrated circuits (Huang et al., 1999).

The Fe-Ni-Co alloy is available in various forms i.e. sheet, wire, bars, tubes or

forged and drawn cups. The chemical requirement for the Kovar alloy after ASTM

F15-78 is shown in Table 2.1.

Table 2.1 Chemical composition of the Kovar alloy (ASTM F15-78, 1995).

Element wt %

Iron, nominal Nickel, nominal Cobalt, nominal Manganese, max Silicon, max Carbon, max Aluminum, max Magnesium, max Zirconium, max Titanium, max Copper, max Chromium, max Molybdenum, max

53A

29A

17A

0.50 0.20 0.04

0.10B

0.10B

0.10B

0.10B 0.20 0.20 0.20

A The iron, nickel, and cobalt requirements listed are nominal. They shall be adjusted by the manufacturer so that the alloy meets the requirements for coefficient of thermal expansion. B The total of aluminum, magnesium, zirconium, and titanium shall not exceed 0.20%.

6

Properties of the Fe-Ni-Co alloy have been reported e.g. by Rosebury (1993)

as followed;

Physical properties;

Density 8.359 g/cm3

Annealed temper (Rockwell hardness) B82 (max)

Cold-worked temper (Rockwell hardness) B100 (max)

Mechanical properties;

0.5% yield strength 410.23 MPa

Ultimate strength 534.34 MPa

Yield strength 303.37 MPa

Uniform elongation 16.8%

Total elongation 35.4%

Thermal properties;

Melting point 1450˚C

Thermal expansion (See Figure 2.1 and Table 2.2 below)

Thermal conductivity (cgs);

30˚C 0.0395 (determined)

300˚C 0.0485 (calculated)

400˚C 0.053 (calculated)

500˚C 0.0585 (calculated)

Curie point 435˚C

7

Specific heat;

0˚C 0.439 J/g ˚C

430˚C 0.643 J/g ˚C

Transformation point (γ- to α- phase) below -80˚C

Figure 2.1 The linear thermal expansion of the Kovar alloy, after annealing in

hydrogen for 1 hour at 900˚C and for 15 minutes at 1100˚C, compared with other

materials (Rosebury, 1993).

8

Table 2.2 The average linear coefficients of thermal expansion of the Fe-Ni-Co alloy

in various temperature ranges (ASTM F15-78, 1995).

Temperature range, ˚C Average linear coefficient of thermal expansion (˚C-1)

30 - 200 30 - 300 30 - 400 30 - 450 30 - 500 30 - 600 30 - 700 30 - 800 30 - 900

5.5 x 10-6

5.1 x 10-6

4.9 x 10-6 5.3 x 10-6 6.2 x 10-6 7.9 x 10-6 9.3 x 10-6 10.4 x 10-6 11.5 x 10-6

The chemical composition of the Fe-Ni-Co alloy used in this experiment is

marked by the star ( ) as shown in Figure 2.2. Two phases of γ- and α-iron should

be present in the alloy microstructure, which was shown as the solid lines for the

temperature at 500˚C. The dot lines in the phase diagram marked the boundaries of

the γ- and α-iron phases at 800˚C. At high temperature, only γ- iron is stable.

9

Figure 2.2 Iron-nickel-cobalt equilibrium diagrams at 500 and 800˚C (Butteridge,

1984).

Fe

Co

Ni

10

2.2 Oxidation of alloys

In hermetic seals, metals are usually preoxidised before joining to glasses.

Therefore, mechanisms of oxidation of iron and its alloys are reviewed here to give

general background. Discussion related to results in the present work will be given in

Chapter 4.

Difference in alloy compositions can cause different rates of diffusion and

oxidation in a complicated way during oxidizing (Birchenall, 1970). Oxidation by

gaseous oxygen is an electrochemical process. The scale of metal oxide (MO) at the

base metal and gas interface is formed as described in equation (2.1) to (2.3). From

these equations, the metal (M) ions are formed at the base metal/MO scale interface

and oxygen is reduced to oxygen ions at the MO/gas interface.

M = M2+ + 2e- (2.1)

12 O2 + 2e- = O2-

(2.2)

M(s) + 12 O2(g) = MO(s) (2.3)

In oxidation process, ionic and electronic transport processes occur through

the oxide scale accompanied by ionizing phase boundary reactions and formation of

new oxide. The high temperature oxidation of metal can be described by Wagner’s

theory, which assumes that the oxide layer is a compact layer and the oxidation rates

are determined by a steady-state diffusion process with interface composition locally

equilibrated. Since the metal oxide is ionic in nature, the diffusion can be therefore

explained in several mechanisms. In stoichiometric ionic compound, the oxidation

process occurs by ionic and electronic migration between two phase boundaries, the

scale-gas and the metal-scale boundaries, as shown in Figure 2.3. This scale can grow

11

by cation migration from metal and react as the scale-gas interface. In case of the

scale formation by anion migration, anion must penetrate and react at the oxide-metal

interface.

Figure 2.3 Schematic of the interfacial reactions and transportation process; (a) cation

transportation and (b) anion transportation in the Wagner’s model (Birks and Meier,

1983).

In the non-stoichiometric ionic compounds, even though the compound is

electrically neutral, the metal/non-metal ratio is not exactly as given in the chemical

formula. These ionic compounds can therefore be classified by their behavior as

negative (n-type) or positive (p-type) semiconductors. The chemical formula for n-

type semiconductors is given as M1+δO. This compound can be formed by either the

existence of metal cation and excess amount of electron over the interstitial sites or by

a non-metal deficit in the structure. The p-type behavior arises from the excess

vacancies and electron hole in the lattice. Figure 2.4 shows examples of n-type and p-

type non-stoichiometric oxides. The value of δ in the chemical formula, M1-δO, can

be in the wide range from 0.05 in FeO, 0.001 in NiO and very small deviations in

12

Cr2O3 and Al2O3, depending on the intrinsic point defect concentration of ion

(Birchenall, 1970; Birks and Meier, 1983).

Figure 2.4 Schematic representation of the structure of non-stoichiometric oxides: (a)

a n-type metal excess semiconductor ZnO and (b) a p-type metal deficit

semiconductor NiO (Birks and Meier, 1983).

In the case of iron and oxygen, it is possible that the oxide layer was developed

by Fe diffusion through the oxide layer toward the surface and formed several stable

compounds including wustite (FeO), magnetite (Fe3O4) and hematite (Fe2O3) at the

temperature above 570˚C as interpreted from the phase diagram in Figure 2.5. The

wustite phase is a p-type semiconductor with a wide range of existing stoichiometry

from Fe0.95O to Fe0.88O. The mobility of cations and electrons via vacancies and

electron holes is extremely high in this phase, because of high concentration of cation

vacancies. Magnetite structure composes of Fe2+ and Fe3+ occupying octahedral sites

and tetrahedral sites with vacancies occuring on both sites. Migration of Fe ions

would take place via these defects. Birks and Meiers (1983) proposed that both the Fe

(a) (b)

13

ions would diffuse outwards through Fe3O4, but only Fe3+ would diffuse outward

through Fe2O3, resulting in outward growth of these oxides.

Figure 2.5 The iron-oxygen phase diagram (Birks and Meier, 1983)

Figure 2.6 represents schematically the iron oxide formation mechanism at the

temperature above 570˚C. Magnetite may form at the magnetite-hematite interface by

the following reaction:

Fen+ + ne- + 4Fe2O3 = 3Fe3O4 (2.4)

New hematite may form at the hematite-gas interface according to the reaction

2Fe3+ + 6e- + (3/2)O2 = Fe2O3 (2.5)

Since the rapid rate of iron reaction and due to a much greater mobility of defect in

FeO, the FeO layer is therefore appeared as a very thick layer compared to the Fe3O4

and Fe2O3 layers. Figure 2.7 shows the relative ratio of FeO : Fe3O4 : Fe2O3 which is

approximately 95 : 4 : 1 at 1000˚C.

14

Figure 2.6 Schematic diagram of the diffusion steps and interfacial reactions of the

three-layer iron oxide scale, which occurred at temperature above 570˚C (Birks and

Meier, 1983).

Figure 2.7 The three-layered scale formed on iron at high temperature oxidation in air

(Fontana, 1986).

15

Effects of additional elements to iron were studied in order to prevent

oxidation of iron. The goal was to slow growth of wustite and prevents depletion of

alloy agents which could lead to acceleration of mechanical failure in wustite scale.

The adding elements such as Ni, Cr, Al and Mo dissolved in the alloy phase but not in

wustite (Birchenall, 1970). In the Fe-Ni-O system, iron was oxidised preferentially

(Birchenall, 1970). Only Co and Mn appear to stabilize the wustite scale. CoO and

MnO can form continuous series of solid solutions with wustite. Small amount of Co

added on the iron solid solution resulted in a decrease of oxidation rate of the alloy

(Driessens and Bunsenges, 1968). MnO is much more stable than wustite and Cr2O3.

The natural growth rate of MnO on manganese is greater than that of wustite on iron

except at very high temperature. However, there are no kinetic effect benefits on the

MnO dissolution on wustite (Birchenall, 1970).

For the oxidation in Fe-Cr alloys, the native cation vacancy levels in pure

FeO, CoO and NiO are approximately 10, 1 and 0.1%, respectively (Wood, 1970).

By adding a small content of Cr in the Fe base alloy, the oxidation could be

sufficiently improved. In the temperature range between 850-1000˚C, the

concentration of 0.3%Cr was enough to prevent the initial formation of wustite and

2%Cr could maintain the oxidation rate about an order of magnitude below that of the

pure iron.

16

2.3 Borosilicate glass

2.3.1 Formation of glass and general properties

Glass is a uniform material which is commonly produced when the viscous

molten material is cooled very rapidly to a rigid condition below its glass transition

temperature without crystallizing. The most familiar form of glass is the silica-based

material used for household objects such as light bulbs and windows. Under tension,

glass is brittle and will break easily. Under compression, pure silica can withstand a

great amount of force. The properties of glass can be modified or changed with the

addition of other compounds or heat treatment. Since glass was developed on the

basis of major commercial used, a large percentage of these are silica-based and more

than 99% of glass compositions are oxides. Most glasses contain about 70–72 % by

weight of silicon dioxide (SiO2). The most common form of glass is the soda-lime

glass, which contains nearly 30 % sodium and calcium oxides or carbonates. Pyrex is

a borosilicate glass containing about 10 % boric oxide. Lead glass is commonly

contains a minimum of 24 % lead oxide. The typical compositions and physical

properties of some commercial glasses are shown in Table 2.3

17Table 2.3 Typical compositions and physical properties of some commercial glasses.

Composition (%wt) Physical properties

Type of glasses SiO2 Na2O K2O CaO MgO B2O3 Al2O3 Fe2O3 PbO BaO

Glass transition

temperature

Tg (ºC)

Annealing

point*

(ºC)

Softening point**

(ºC)

Thermal expansion

coefficient(10-6/K)

(20-300ºC)

Electrical Resistance (Ω-cm at 250ºC)

Ref.

Fused silica 99.9 <0.001 <0.001 <0.001 - - 0.005 - - - 1100 1150 1650 0.54 12.0 Pfaender, 1983

Vycor 96.0 <0.03 - - - <3.5 <0.3 - - - 1050 1020 1530 0.75 9.7 Pfaender, 1983; Holloway, 1973

Soda-lime sheet glass 72.2 14.2 - 10.0 2.5 - 0.6 - - - No

information 548 730 8.5 6.5 Pual, 1982; Varshneya, 1994

Soda-lime plate glass 72.5 13.0 0.3 9.3 3.0 - 1.5 0.1 - - No

information 553 735 8.7 6.7 Holloway, 1973: Pual, 1982

Soda-lime container glass 73.0 15.0 - 10.0 - - 1.0 0.05 - - No

information 548 730 8.5 7.0 Holloway, 1973: Pual, 1982

Soda-lime bulb glass 73.6 16.0 0.6 5.2 3.6 - 1.0 - - - No

information 510 696 9.2 6.4 Pual, 1982; Varshneya, 1994

Lead alkali silicate 63.0 7.6 6.0 0.3 0.2 0.2 0.6 - 21.0 - 435 430 626 9.1 8.9 Pfaender, 1983;

Pual, 1982

High lead alkali silicate 30.0 - 3.0 - - - - - 65.0 - 435 430 580 9.1 11.8 Pfaender, 1983;

Pual, 1982

Alumino silicate 54.6 0.6 17.4 4.5 8.0 14.8 - - No

information 715 9.1 11.4 Holloway, 1973; Pual, 1982

Borosilicate (Pyrex) 80.6 4.2 - 0.1 0.05 12.6 2.2 0.05 - - No

information 565 862 No information

No information

Holloway, 1973; Pual, 1982

Low expansion borosilicate 81 4 - - - 12 2 - - - No

information 565 820 3.2 8.1 Pual, 1982; van Vlack,

1964 Low electrical loss borosilicate

- - - - - - - - - - No information 495 No

information 3.2 11.2 Pual, 1982

Borosilicate for tungsten seal 75.1 4.3 1.4 0.7 0.4 16.7 1.3 - - - 523 535 760 4.0 8.3 Pfaender, 1983

Borosilicate for Molybdemun seal

75.6 6.6 - 0.3 - 8.8 4.5 - - 3.9 565 570 570 4.9 6.9 Pfaender, 1983

Borosilicate for Kovar seal 68.7 0.8 7.5 - 0.6 18.6 3.0 - - - 495 507 715 5.0 10.3 Pfaender, 1983

17

* Annealing point corresponds to a viscosity of 1013 poises and represents a temperature at which internal strains are reduced to an acceptable limit in 15 minutes.

** Softening point corresponds to a viscosity of 107.5 to 108 poises.

18

SiO2 is the most important raw material in glass manufacture. It is a basic

glass former (network former). Crystalline silica has a very high melting point and

liquid silica has high viscosity compare to the other glasses. High concentration of

silica in glasses gives high softening temperature, low thermal expansion, and good

chemical durability.

B2O3 normally functions as a network former in glasses. It is an important

composition in special glasses which are used in electrotechnology, especially in the

fields of heat and chemical resistance, excellent electrical insulation, low electrical

loss, and gaseous impermeability. B2O3 will join the network structure of silica

glasses without producing adverse change in the thermal expansion and durability. In

the Kovar sealing, the higher percentages of B2O3 (17%-23%) are necessary if the

glass-transition temperature of glasses must be reduced below 510˚C (Pfaender,

1983).

Na2O is a network modifier which normally added in the form of soda ash

(Na2CO3). This fluxing agent lowers softening point in glasses, raises thermal

expansion and ionic conductivity, and reduces the glass durability.

K2O, normally added in the form of potassium carbonate (K2CO3), is a

network modifier similar to the Na2O. It does not only contribute to the optical or

thermal properties that are often desired, but also increases the workability of the

glass by increasing its fluidity (van Vlack, 1964).

Al2O3 is an intermediate. It is usually added to the glass batch in the form of

felspars to join the network as AlO4 tetrahedra. Al2O3 improves the chemical

durability and increases viscosity in the lower temperature ranges, strongly suppresses

devitrification, and makes melting and refining of the glasses more difficult.

19

According to Zachariasen’s rules, it has been considered that a substance can

form extended three-dimensional networks lacking of periodicity with energy content

comparable with that of the corresponding crystal network. These rules were

remarkably successful in predicting new glass-forming oxides. The rules are as

follows (Doremus, 1973);

(1) An oxygen atom links to not more than two cations.

(2) The number of oxygens surrounding these cations must be small.

(3) The oxygen polyhedra share with one another by their corners, not by their

edges or faces.

(4) At least three corners of each polyhedra should be shared.

Oxides, such as SiO2, B2O3, P2O5, GeO2 and BeF2, are called network formers

because of their ability to form branching network structures. These network formers

are generally 3 to 4 in coordination numbers. Goldschmidt also considered the crystal

structures and their relation to the ionic sizes, and postulated a correlation between the

ability to form glass and the relative sizes of the oxygen and cation atoms. The ratios

between the radius of cation (RC) and radius of anion (RO) in glass-forming oxides are

in the range of about 0.2 to 0.4. The radius ratios of typical glass formers are shown

in Table 2.4.

20

Table 2.4 Radius ratios for typical network-formers (Babcock, 1977).

Compounds Radius ratio(Ra/Ro)

SiO2 RSi : RO = 0.39 Å : 1.4 Å ≈ 0.28

B2O3 RB : RO = 0.20 Å : 1.4 Å ≈ 0.15

P2O5 RP : RO = 0.34 Å : 1.4 Å ≈ 0.25

GeO2 RGe : RO = 0.44 Å : 1.4 Å ≈ 0.31

BeF2 RBe : RO = 0.34 Å : 1.36 Å ≈ 0.25

Following Zachariasen’s constraints, the network structure of pure SiO2 glass

is such that each silicon ions bonded to four oxygen ions and each oxygen ions

bonded to two silicon ions with the oxygen-silicon-oxygen bond at the φ angle of

approximately 109°28´, as shown in Figure 2.8. The continuous random network

structure occurs through corner-to-corner connections of the SiO4 tetrahedra. The β

angle is the bond angle between two adjacent SiO4 tetrahedra. The mutual orientation

of the adjacent tetrahedral is defined by ψ angle (Allen and Thomas, 1999).

Figure 2.8 The structure of pure SiO2 glass. A bond-and-stick model of two SiO4

tetrahedra, φ ≈ 109˚28´ (adapted from Allen and Thomas, 1999).

21

The pure SiO2 structure is a very dense network structure and has a very high

glass transition temperature of about 1430 K. The uniformly connected atomic

arrangement represents the crystalline silica (Figure 2.9a). On the other hand, the

formation of disordered structure represents the vitreous structure (Figure 2.9b). To

reduce the glass transition temperature of the SiO2 glass, network modifiers, such as

Na2O, K2O, CaO, and BaO, have been added to the glass. The coordination numbers

of these network modifier cations are generally equal to or more than 6. These

modifiers weaken the pure SiO2 network structure by alternating the bonding of the

oxygen atoms. The alkali oxide (M2O) enters the glass as singly charged cations and

occupies the interstitial sites in the glass structure (Varshneya, 1994). Some oxygen

ions in the modified glass are then ionically bonded to the adjacent modifier cations

rather than forming a rigid ionic-covalent bond with the silicon atoms as shown in the

reaction below.

Each alkali ion is therefore expected to create one ‘non-bridging oxygen’ (NBO). A

schematic represented the alkali silicate network is shown in Figure 2.8c. Other

oxides called intermediates, such as Al2O3, TiO2, and ZnO, can behave either as

network formers of network modifiers depending on the basicity of the other oxides in

the glass.

Si O Si + M2O = Si O-M+ M+ -O Si

22

(c)

Figure 2.9 Schematic of two-dimensional SiO4 tetrahedra structure; (a) regularly

ordered structure in the crystalline SiO2, (b) disordered structure of the amorphous

SiO2 glass and (c) formation of NBO structure in sodium silicate glass (Allen and

Thomas, 1999; Harper, 2001).

23

In the glass, there is no sharp discontinuous transition into the liquid state but

there observes the progression in viscosity as the temperature increases through the

transformation range. The reference conditions of various transition types have been

defined in terms of the viscosity (η) and temperature as shown in Figure 2.10. The

viscosity of some commercial glasses in the melting tanks is at the order of 102 poise.

The viscosity in between 103 - 106 poise (typically at is 104 poise) when glass are

pressed into molds or drawn into tubing or rod is called the working point. The

softening point is the temperature at which η = 107.6 poise, whereas the strain point at

η = 1014.5 poise and the annealing point at η = 1013.4 poise, respectively. The strain

point corresponds to the highest temperature at which the glass can be rapidly cooled

without a serious internal stress, while the annealing point is the temperature at which

any internal stress will be relieved in a few minutes.

24

Figure 2.10 Variations of viscosity with temperature of some commercial glasses

(Pfaender, 1983).

25

2.3.2 Structure of borosilicate glasses

A large number of commercial glasses are produced for specialized

application based on borosilicate systems which are primarily known for their thermal

shock resistance and excellent chemical durability. By adding B2O3 and small

amounts of alkali to silica resulted in their desirable properties. The improvement in

thermal shock resistance results from a lower thermal expansion coefficient which

therefore make it an excellent material for laboratory glassware and for use in large-

scale technological plants in the chemical apparatus industry. The borosilicate glasses

are resistant to corrosion and remain absolutely neutral, even to aggressive chemicals

in nearly all fields of chemistry. They possess high strength and heat resistance.

These give them great advantages over other materials. Therefore, this type of glasses

is suitable for making x-ray tube and also other evacuated tubes in a wide range of

technology, because of its special qualities as low x-ray absorption, high UV-

transmission, and high electrical insulation.

As mentioned above, boron and silicon are network formers in the borosilicate

glass structure. The alkali reacted as the network modifiers entering in the glass as

single charged cations occupying interstitial sites. There are two possibilities of an

oxygen ion from a modifier oxide to modify the structure of boric oxide glass; 1) by

creating NBO as shown in Figure 2.11a and 2) by converting from a 3-coordination

(BO3) state to a 4-coordination (BO4 ) state as shown in Figure 2.11b, respectively

(Varshneya, 1994). Furthermore, Vogel (1985) also suggested that the borate glasses

can form as various structures as shown in Figure 2.12.

26

Figure 2.11 Schematic of two possibilities of an oxygen from a modifier oxide to

modify the structure of boric oxide glass (Varshneya, 1994).

Figure 2.12 Schematic of the possible structure in borate glasses (Vogel, 1985).

(a)

(b)

27

In the borosilicate glass suitable for glass-metal sealing, Al2O3 is necessary.

The trivalent aluminum ion does not always act as a network former but the structural

configuration depends on the Al2O3/M2O ratio (Varshneya, 1994). The Al3+ ion will

act as the network former in tetrahedral coordination only when the Al2O3/M2O ratio

< 1. By adding some of aluminum ions into alkali silicate glass, NBO in the structure

can therefore be removed. On the other hand, when the Al2O3/M2O ratio > 1, the Al3+

ions may enter the network as a modifier as an octahedral coordination. Figure 2.13

illustrates an Al3+ ion reacting with six groups of oxygen, three of NBO and another

three of BO.

Figure 2.13 Schematic of octrahedrally coordinated network modifier of Al3+ in alkali

aluminosilicate glass structure (Varshneya, 1994)

28

However, there are also other three models of the tetrahedrally coordinated

Al3+ as shown in Figure 2.14. In the electrically neutral structure, the tricluster may

form as one AlO4 and two SiO4 (Figure 2.14a) as well as two AlO4 and one SiO4

(Figure 2.14b). Figure 2.14c represented the tricluster of an Al3+ ion acting as a

modifier and bonding to three NBOs.

Figure 2.14 Schematic of tetrahedrally coordinated Al3+ in tricluster arrangement

(Varshneya, 1994)

2.3.3 The volume-temperature diagram and the glass transition

The volume-temperature (V-T) diagram for glass-forming liquid is shown in

Figure 2.15. When a liquid that normally does not form a glass is cooled, the volume

gradually decreases along the path ‘abc’ which point ‘b’ corresponding to melting

point (Tm). Crystallization occur at or slightly below the Tm if there are a sufficiently

large number of nuclei present and a large enough crystal growth rate exists. The

location of the point ‘c’ below Tm varies depending upon (i) when the thermodynamic

29

driving force is created by the undercooling causing a particular group of atoms to

transform from the liquid state to the crystal state and (ii) the velocity at which the

atoms from the liquid can be transported to the crystal-liquid interface. The shaded

region shows the varying probability representing the crystallization path. Upon

further cooling, volume shrinkage generally accompanies the crystallization leading to

shrink along the crystal line to the point ‘e’.

However, crystallization will not occur below Tm if the cooling rate is high.

The liquid mass moves into the supercooled liquid state along the line ‘bcf’, which is

the extrapolation of the line ‘abc’. No discontinuity in the V-T is observed. The

volume shrinks continuously, so the structure of the liquid rearranges itself into a

lower volume along the line ‘bcf’ required by the lower energy corresponding to a

lowered temperature. As cooling continues, the molecules become less and less

mobile, the viscosity of the system rapidly increases. At sufficiently low

temperatures, the molecules cannot rearrange themselves fast enough to reach the

volume characteristic of that temperature. The state line then starts a smooth

departure from ‘bcf’ and soon becomes a near-straight line often roughly parallel to

‘de’. This line ending at ‘g’ when cooling fast or at ‘h’ when cooling is slow. The

material then behaves as a glassy solid state.

The smooth curve between the onset of departure from the supercooled liquid

line to a seemingly rigid condition is termed the glass transition region, or the glass

transformation range. The intersection of the extrapolated glass line and the

supercooled liquid line is termed the fictive temperature (Tf). At this temperature, the

structure of the supercooled liquid is instantly frozen into the glass.

30

The departure from the supercooled liquid line is dependent upon the rate of

cooling. Slower cooling allows the structure to rearrange itself to stay on the path

‘bcf’ somewhat longer, and hence the more slowly cooled glass at ‘h’ would be

expected to have a lower volume (higher density) and a lower fictive temperature than

a more quickly cooled glass at point ‘g’.

When the glass at ‘h’ is reheated, the state smoothly moves through the

transition region along the dashed curve to the supercooled liquid state ‘fcb’ and

ultimately to the liquid state. The V-T curve never retraces its path in the transition

region. When the crystals at ‘e’ are heated, the state will move along the crystal line

up to ‘d’ (at Tm) past the shaded region, melt at Tm to reach point ‘b’ and subsequently

follow the liquid path ‘ba’.

Figure 2.15 The volume-temperature diagram of glass (Varshneya, 1994)

31

2.3.4 Phase transformation in amorphous glass

The types of phase transformation which is usually observed in glass are

shown in Figure 2.16. The definitions are as follows.

Figure 2.16 Phase transformation in amorphous glass (Pual, 1982)

Crystallization: The growth of a crystalline phase which may or may not have the

same composition as the original liquid. The structure is obviously different.

Surface crystallization: Crystal growth begins from the glass/atmosphere interface,

and crystals usually grow perpendicular to this interface.

Volume crystallization: Crystal growth begins from the nucleation sites within the

body of the materials. The initiating site for crystallization may be e.g. a substance

foreign to the bulk material, when it is called ‘heterogeneous nucleation’. On the

other hand, if the nucleus is the same as the bulk material, it is called ‘homogeneous

nucleation’.

Phase Transformation

Crystallization Liquid-liquid Phase Separation

Volume Crystallization

Surface Crystallization

Homogeneous Nucleation

Heterogeneous Nucleation

Spinodal Decomposition

32

Liquid-liquid phase separation: The growth of the second liquid phase which will

have a different composition from the original liquid phase. A single-component

system cannot separate in this way. This separation can occur either by nucleation

and growth mechanism or spinodal decomposition mechanism.

Spinodal decomposition: Within a region which separates into two liquid phases,

there will be a region where there is no energy barrier to nucleation and phase

separation is, therefore, limited by diffusion only.

There are four factors affecting the crystallization or devitrification of glasses;

(1) time, (2) temperature, (3) nucleation, and (4) internal structure (van Vlack, 1964).

The crystallization must eventually take place in the glass if the time is long enough

because the free energy of the glassy phase is higher than that of the crystalline phase.

When the temperature is increased, the faster rate of crystallization occurs. The

activation energy relationship is approximated by an Arrhenius equation (van Vlack,

1964).

Rate = Ae -E/kT (2.4)

or

loge (1/t) = A´ - E/kT (2.5)

Where

t = time taking to achieve any given state of reaction

E = the activation energy

k = Boltzman’s constant

T = the absolute temperature

A = constant

A´ = loge A

33

The higher temperature favors the devitrification because the particular

network bond receives the sufficient energy to be ruptured. The faster rate of

formation of the lower-energy and long-range crystalline structure is therefore

expected. However, it should be noted that the maximum rate of devitrifiation is in

the range below the true crystallization temperature. This is because the

devitrification depends not only on the probability of bond rupture and long-range

ordering, but also on the new surface formation between the glassy and crystalline

phases which requires thermodymically driving force. At the true crystallization

temperature, there is a very little driving force to support the nucleation forming new

surfaces. Figure 2.17 shows a Time-Temperature Transformation (TTT) diagram of

the devitrification of silica and fayalite (Fe2SiO4) from glasses. The maximum rate of

crystallization occurs at moderate temperature slightly below the true freezing

temperature. Too high temperature leads to limitation in thermodynamic driving

force to form new crystalline/glass interface by nucleation, while too low temperature

leads to limitation in kinetic energy required for mobility of atoms in growth

according to the Arrhenius relationship. The devitrification rate of SiO2 crystalline is

slower than Fe2SiO4 because the SiO2 crystalline has an excellent network structure,

whilst the Fe2SiO4 crystalline has a large percentage of discrete SiO44- ions and

relatively little polymerization.

34

Figure 2.17 TTT diagram of devitrification in glasses to form (a) network silica and

(b) partially ionic Fayalite (Fe2SiO4) (van Vlack, 1964)

Apart from nucleation and growth mechanism, spinodal decomposition is

another special process of phase transformation in glasses. To describe this type of

phase transformation, a binary system of which components show complete solid

solubility above a critical temperature Tc is illustrated in Figure 2.18a. Below Tc the

system exhibits immiscibility gap. Homogeneous single-phase solid within the

immiscibility gap will decompose into two phases, α1 and α2, which have different

compositions but the same structure. An area within the immiscibility dome, which is

known as spinodal, is indicated by dash line. A melt within this dome (e.g. X

between X′s and X′′s at the temperature T′) will spontaneously separate if the mobility

(kinetic energy) of the ions is great enough. However, if the initial composition of the

melt within the immiscibility gap is outside the spinodal dome (e.g. X′ between X*e

X′s at the temperatureT′), phase separation will not take place spontaneously, but will

35

require the formation of nuclei (Pual, 1982). Schematic variation of the free energy of

single-phase solids at the temperature T΄ is given in Figure 2.18b. The relationship of

free energy versus composition in Figure 2.18b can be written as follows;

For stability as well as metastability:

⎟⎟⎠

⎞⎜⎜⎝

⎛∂∂

2

2

xG > 0 (2.6)

For the boundary of spinodal (X′s and X′′s):

⎟⎟⎠

⎞⎜⎜⎝

⎛∂∂

2

2

xG = 0 (2.7)

For the instability:

⎟⎟⎠

⎞⎜⎜⎝

⎛∂∂

2

2

xG < 0 (2.8)

With these conditions, the α1 and α2 phases form a stability phase. The

undecomposed solids between phase boundary and spinodal are in metastable

equilibrium, while those within the spinodal are unstable. In many silicate and borate

melts, two-liquid phase formations can be observed according to spinodal

decomposition.

36

Figure 2.18 (a) Schematic phase diagram of the miscibility gap, two phase boundary

and spinodal of a binary system. (b) Variation of free energy with composition of

single-phase solids at temperature T΄ (Jena and Chaturvedi, 1992).

(a) (b)

37

2.4 Theory of glass-metal sealing and bonding mechanism

The Fe-Ni-Co alloy to borosilicate glass seals have been used for many years

as hermetic and electrically insulating seals. It is not only excellent in thermal

expansion matching but also in good wettability and bond strength (Macey et al.,

1985). It was well known that the thermal expansion coefficient of the metal needs to

be as close as possible to the glass and the oxide layer should be developed on the

metal surface before sealing (Mantel, 2000). These mean the strong glass-metal

adhesion requires a continuous transition from the metallic structure of the metal,

gradually, to ionic-covalent structure of the glass (Hong and Holland, 1989; Mantel,

2000). The transition layer usually consists of metal oxides because they are

compatible with both glass and substrate metal. Adherence oxide is developed at the

glass-metal interface by achieving and maintaining equilibrium compositions at the

interfaces.

2.4.1 Glass-metal sealing

The metals commonly used in glass-metal sealing are platinum, tungsten,

molybdenum, copper, iron, nickel, chromium, and iron alloys. The most physical

requirement is a thermal expansion matching between two materials. For glass to

metal seals, the materials are usually selected in terms of the thermal expansion

coefficient. The classification system of the glass-metal seals based on relative

thermal expansion is shown in Figure 2.19.

38

Figure 2.19 Types of glass-metal seal (adapted form Lancaster, 1993)

There are two types of glass-metal joints, matched and unmatched seals. In

matched seals, the thermal expansion coefficient of glass is matched as close as

possible to the metal. Chemical bonding occurs at the glass-metal interface. The

hermetic seal is an example and is typically made by forming a bond between the

glass and the oxide created on the metal components prior to sealing.

For unmatched seals, it can be divided into two categories.

(1) The compression seal, the glass surrounds the metal duct and itself is

surrounded by the metal ring as shown in Figure 2.20. The glass which contains

residual tensile stress is under compression at room temperature and hence chemical

bonding is not necessary.

(2) The ductile seal or Housekeeper, as illustrated in Figure 2.21, is the most

commonly used in joining glass and copper. Copper can be joined with glass to

produce a seal if it is thin enough. The tip of the copper tube is tapered down to give

Glass-Metal

Unmatched

Compression Ductile metal (Housekeeper)

Matched

39

a very thin wedge-shaped section. The copper and glass bead are heated to about

1000˚C, and then bring together and seal the matching glass to the beaded glass. The

copper-glass seals must be annealed immediately after sealing. There are no

significant stress in the glass and the metal. The thin-sectioned copper makes it weak

enough to follow the expansion and contraction of the glass (Shand, 1958; Lancaster,

1993; Rosebury, 1993). These techniques have also been adapted to the case of the

Kovar and stainless steels, but the chemical bonging is also required. The various

types of Housekeeper seal are given in Figure 2.22.

Figure 2.20 Typical designs of compression glass-metal seal (Lancaster, 1993)

Glass bead metal

40

Figure 2.21 Steps in making a ductile or Housekeeper seal. (1) the copper tube is

thinned; (2) glass bead is applied to the copper edge; (3) the glass tubing is sealed to

the bead (Shand, 1958).

Figure 2.22 The various types of Housekeeper seal (Rosebury, 1993)

(1)

(2)

(3)

41

In order to make a glass-metal seal, there are four basic requirements to get a

good adherence (Shand, 1958).

1. The glass should wet and adhere to the metal.

2. The linear thermal expansion of the glass must match as close as possible

with the metal over a wide range of temperature.

3. The metal should have no thermal critical points within the range of the

highest temperature reached in sealing and the lowest temperature reached

in service.

4. When heated for making the seal, the glass should not reboil and the metal

should not degas.

Formation of the glass-metal seal has been described by Pask (1964) and

Rosebury (1993). Initially, the specific sealing glass is molten on or around the metal

and does not support any stress. The strain is occurring when the glass is cooled.

Therefore, the metal and glass to be jointed successfully should have a difference in

thermal expansion coefficient less than 10%. Figure 2.23 shows a schematic of

relative thermal expansion coefficient of glass and metal. Before joining, the glass

and metal surface must be cleaned to remove all materials that are not soluble in the

glass at sealing temperature. This cleaning causes both etching and roughening of the

metal surface. For good bonding, the cleaned metal is usually preoxidised if the

amount of oxidation on the sealing process seems to be insufficient. The linear

thermal contraction curves are more important for glass sealing applications (Harper,

2001). The glass contraction curve is overlaid on the metal contraction curve and

vertically displaced to have the curves coincide at the set point, which is often taken

42

as somewhere between the strain point and the annealing point of glass. The offset on

the y-axis is the yield (ε) which is the different in contraction or mismatch of the glass

and the metal.

Figure 2.23 Thermal matching of glass-to-metal seal. The contraction curves of the

glass and the metal are matched at the set point (Harper, 2001).

2.4.2 Glass-metal bonding mechanisms

There are three types of bonds possible between solid materials; (1) van der

Waals forces, (2) mechanical bonds and (3) chemical bonds (King, 1959). The field

in which glass-metal bonding mechanisms was first studied was probably the enamel-

metal bonding. The adherence mechanisms between enamel and metal interface can

be explained by various theories. These theories can be divided by the

phenomenological observation into two categories; a mechanical bond (an

interlocking between the glass-metal interface) and a chemical bond (an intermediate

oxide layer between the glass and metal) (Pask, 1964). For poor adherence, the

43

interface exhibited a van der Waals bond, which is weaker than a chemical bond.

Details of mechanical and chemical mechanisms are described below.

2.4.2.1 Mechanical bonding

The mechanical theories of adherence originally gained acceptance due to the

fact that the interface between a metal and a porcelain enamel often appeared rough,

even if the substrate was quite smooth. Bonding was occurred due to the mechanical

effect between the roughening substrate and the glass. A number of theories arose

based on the roughening mechanism, e.g. the dendrite theory which favored the

precipitation of a metallic phase within the glass and the electrolytic theory which

proposed that electrolytic attack roughened the surface. The result of these bonding is

an interlooking metal-glass structure. Therefore, all mechanical theories require the

existence of a rough glass-metal interface.

1) The dendrite theory

The dendrites of a metallic phase, e.g. cobalt or Co-Fe intermetallic phase,

were often found in the interfacial region. It was believed that the dendrites formed

due to the reaction of a metal oxide presented in the glass with a metallic element in

the metal substrate. The dendrites performed anchor points for the glass and the metal

substrate, as illustrated in Figure 2.24a. For example, when a glass containing CoO

was bonded to an iron substrate, the reaction was believed to occur as follows;

CoO(glass) + Fe(substrate) Co(dendrite) + FeO (glass) (2.9)

44

2) The electrolytic or galvanic-corrosion theory

The roughening of the metal substrate was believed to occur during coating of

the enamel due to an electrolytic or a galvanic-corrosion mechanism. According to

the Dietzel theory, in a coating of CoO-containing glass to an iron substrate in air,

bonding resulted in the precipitation of metallic cobalt, as proposed in the dendrite

theory. However, it was believed that the localized galvanic or electrolytic cells were

formed in contact of Co and iron substrate. The current (positive electricity) flowed

from the iron through an enamel to the cobalt and back to the iron, and hence, the

direction of electron flow was vice versa. The result was that the iron dissolved

continuously into a solid solution of the solid enamel. Thus, the surface became

roughened and the glass anchored itself into the holes. The net effect of the

dissolution of iron and the formation of a pitted surface was illustrated in Figure

2.24b. The reaction required to support the mechanism are as follows;

Fe(substrate) + CoO(glass) FeO(glass) + Co (to establish cell at the interface) (2.10)

2Co (precipitates at the interface) + O2 (from atmosphere) 2Co2+(glass)+ 2O2-

(glass) (2.11)

Co2+(glass) + 2e- Co(interface) (2.12)

Fe(substrate) - 2e- Fe2+(glass) (2.13)

Eventhough, the galvanic corrosion can occur under normal firing condition,

the theory has a weakness when the other metal pairs with their galvanic behavior

different from iron and cobalt are considered.

45

Figure 2.24 Schematic diagrams of mechanical anchor points of glass-to-metal

adhesion; (a) the dendrite theory, (b) the electrolytic theory (Donald, 1993)

In purely mechanical bond, the metal is visualized as keying the glass to its

surface in a sort of dovetail arrangement (King, 1959) as shown in Figure 2.25. The

strength of the interface may be strong (approach to the chemical bonds) or weak

depending on the local flaws and resultant stress concentrations.

(a) (b)

46

Metal

Glass

Glass

Metal

(a)

(b)

Figure 2.25 Metal-glass interface that might promote mechanical bonding; (a) ideal

dovetail joint and (b) generalized metal-glass dovetail (adapted from King, 1959)

2.4.2.2 Chemical bonding

A strong chemical bond was formed if the ‘transition zone’ occurred in the

glass-metal interface. In this zone, the metallic bonding of the metal is gradually

substituted by the ionic-covalent bonding of the glass. The glass can become

saturated with an oxide of the substrate. King et al. (1959) suggested that when the

appropriate metal oxide is dissolved in the glass up to its saturation point, the metal

ions remain at the surface and perform the metal-metal bonding across the interface.

When the temperature increased, metal ions from the glass will diffuse into the metal

where they will gain electron and become zero valence metal atoms. On the other

hand, metal atoms in the substrate will also diffuse into the glass and become ionized.

If the glass is not already saturated with the appropriate oxide, the dynamic

47

equilibrium of the glass-metal interface cannot occur. At the lower temperature

where the atoms and ions are less mobile, the exchange mechanism of electron at the

interface may alternate depending on the concentration and ionization of metal atoms.

The metal oxide is usually obtained by preoxidation of the metal interface.

The surface oxide will dissolve into the glass during sealing and making the

adherence layer onto glass-metal interface. Different possible structures at the

interfacial region of glass-to-metal bonding are compared in Figure 2.26. In Figure

2.26a, glass is saturated with substrate metal oxide in the interfacial region to give the

strong chemical bonding via a ‘mono-oxide’ layer. In Figure 2.26b, the same

situation is held as in Figure 2.26a but the chemical bonding occurs via a ‘bulk’ oxide

layer where the strength of the bonding depending on the properties of the bulk oxide

layer. Finally, in Figure 2.26c, the unsaturated glass having a weakly van der Waals

bonding with the metal.

48

Figure 2.26 Simplified schematic representations of glass-to-metal bonding; (a) a

mono-oxide layer, (b) a bulk oxide layer, and (c) unsaturated glass (Donald, 1993)

(a) (b) (c)

49

2.5 Joining process of borosilicate glass to Fe-Ni-Co alloy

For the Kovar alloy, Pask (1964) reported the degree of metal preoxidation

that is suitable for glass-metal sealing as shown in Figure 2.27. The optimum weight

gain should be in between 0.33-0.66 mg/cm2. It was also mentioned that the strength

of the seal is dependent upon the strength of the surface oxide and its adherence to the

metal. If the oxide layer is too thin, the sealing is poor in strength, and if the oxide

layer is too thick, an amount of oxide remains after sealing. After sealing, the

samples must be annealed or cooled slowly down to the room temperature in order to

avoid destructive thermal strains.

Figure 2.27 Degree of preoxidation of the Kovar alloy as a function of time and

temperature (Pask, 1964)

50

The Fe-Ni-Co alloy to borosilicate glass seals were made by Ikeda and

Sameshima (1964). The alloy parts were degreased and annealed at 1,000˚C for 15

minutes in wet H2, preoxidised in the air and then sealed with glass in N2 atmosphere.

Three kinds of reagents, HF, HCl and HCl + H2O2, were selected in order to etch the

glass phase, the alloy phase and the oxide phase, respectively. The seals made by

preoxidising the alloys to 0.45 mg/cm2 in weight gain at 750˚C in air and applying the

glass to the alloy at 1,000˚C in N2 for 3, 7, 15 and 30 minutes, respectively.

Zanchetta et al. (1995a) used the small discs of the Kovar alloy (10 mm

diameter, 0.9 mm thickness) with about 580 mg in weight. A preoxidation treatment

in air was first performed by heating to 700˚C with a rate of 25 K/min, maintained for

5 minutes and followed by cooling at the same rate. Porcelain discs were obtained by

slip casting in the plaster mould. The green pieces were fired in air at 1280˚C and

covered after cooling. Discs were dried at 100˚C for 30 minutes and heated at 910˚C

in air. During sealing, the Kovar disc was put on the enameled porcelain piece and

they were first heated under vacuum (10 Pa) at 300˚C and then under flowing argon.

After the thermal cycle, the soaking temperatures of 910 or 940˚C were reached at a

rate of 15 K/min except for the last 20 or 30 K where the rate was reduced to 3 K/min.

After the time varying from 3 to 60 minutes, samples were rapidly cooled to 620˚C,

then slowly to 550˚C where they were annealed for 15 minutes.

The oxide scales of a Fe-Ni-Co alloy formed by preoxidation with reducing

LPG/O2 or C2H2/O2 flame before direct fusion to a borosilicate glass were studied by

Piyavit et al. (2006). Preoxidation time was in the range of 1-3 minutes. The degree

of oxidation was measured by weight gain of the specimens. The glass rod was cut

into pieces with 1.5 cm in length and 1.5 g in weight. Specimens needed to be

51

cleaned before joining to eliminate organic substances on their surface. For the

borosilicate glass, this was done by immersing in 10 vol% hydrosulfuric solution. For

the Fe-Ni-Co alloy, cleaning was done by immersing in distilled water with 5 vol%

hydrogen peroxide addition and heated to the boiling point for 30 minutes. After

immersion, the specimens were rinsed with distilled water and dried in the electrical

oven. Both the glass and the alloy were then joined by direct fusion using an

oxidising flame from a gas burner. Direct fusion was performed for 3 minutes during

which the glass piece was melted and wet onto the alloy. The joint was consequently

transferred to an electrical furnace and annealed at 550˚C for 30 minutes followed by

furnace cooling. Good adherence was achieved at the preoxidation time of 2 minutes

due to the saturation of FeO in the interfacial glass and the gain weight in these

experiment were in the range of 0.59-0.83 mg/cm2 which is comparable to the range

of weight gain suggested by Pask (1964).

2.6 Observation of the glass to metal joint

2.6.1 General observations of the glass/metal interface

Figure 2.28 shows the concentration distribution of Fe, Ni and Co at the alloy-

oxide boundary oxidised to 0.45 mg/cm2 in weight again before sealing to glass

(Ikeda and Sameshima, 1964). Fe was diffused through the oxide layer toward the

surface. Fe at the surface of the base alloy was decreased, leading to enrichment of Ni

and Co contents at the alloy surface.

52

Figure 2.28 Concentration distribution of Fe, Ni, and Co at the alloy-oxide boundary

(Ikeda and Sameshima, 1964)

Borom and Pask (1966) suggested that the good adherence of porcelain

enamels on iron results from saturation of the glass at the interface with FeO. An

important feature was the metallic layer that called “barrier layers” occurring at the

position of the original metal surfaces. These layers contained dendrite isolated from

the base metal through the diffusion zone in the glass. Heterogeneous nucleation of

these dendrites could occur at defect sites, such as bubbles. In their experiment, a

model system of Na2Si2O5 glass, Fe, and FeO (FexO where x = 0.875 to 0.946 at

1000˚C) was used to illustrate the potential reaction and the conditions under which

chemical bonding occurred. When the Na2Si2O5 was placed in contact with Fe

containing FeO layer, the solution of FeO in form of Na2Fez Si2O5+z took place

53

resulting in saturation at the interface with the oxide. The hypothetical activity of iron

through the cross section of glass-iron was shown in Figure 2.29. In curve t0, the

glass was just placed in contact with the oxidised iron, an equilibrium phase due to

equal activities between iron in the metal and the oxide at their interface. Curve t1

shows equal activities of iron in the glass and oxide at their interface. This condition

is maintained as long as oxide layer exists because the solution rate is higher than the

diffusion rate. Curve t2 represents the situation when the last discrete layer of the

oxide has been dissolved and the glass at the interface still retains an oxide-like

structure. Finally, the t3 curve represents the situation sometimes after the oxide has

been completely dissolved and after iron concentration in the glass at the interface has

been dropped because of the diffusion into the bulk glass.

Figure 2.29 Diagram of hypothetical iron activity and penetration distance in oxidised

iron-glass contact zone showing ferrous iron activity in the oxide and the glass

relative to metallic iron as the standard state (Borom and Pask, 1966)

54

Zanchetta et al. (1995b) reported that the oxidation layer of the Kovar alloy

was mainly composed of iron oxides. Underneath the iron oxides is the porous zone

which enriched in cobalt and nickel. The oxide layer was dissolved by the glass

during their reaction leading to improvement of the wetting quality and good

junctions. When the Kovar alloy was oxidised, the composition of the alloy near the

interface was changed and the areas existing in the interfacial zone are shown in

Figure 2.30.

Figure 2.30 Major components of the interfacial zone after bonding (Zanchetta et al.,

1995b)

55

Concentration profiles of elements in the four distinct parts in the interfacial zone

from to the Kovar alloy to glass are given by Zanchetta et al. (1995b) in Figure 2.31.

Area I is the original alloy with its nominal composition. Area II is the

Fe-impoverished alloy zone. Area III is in the narrow range which contains both

glass and iron oxide, partially or entirely dissolved. The last part is area IV where the

iron diffuses into the glass. The dissolution of oxide into the glass increases the

thermal expansion of the glass to the values approaching those of the iron-

impoverished alloy. Finally, the preoxidation creates an open porosity in the alloy in

which FeO-rich glass penetrates during bonding treatment leading to a strong physical

adhesion. The duration of the bonding thermal treatment is very important. If the

time is too short, the brittle oxide layer remains and the FeO-dissolved gradient are

very strong. If the time is too long, the homogenization of the glass leads to a

significant gap of the thermal expansion at the interface of the glass and the alloy so

that the bonding will be broken during cooling. For the best binding condition, there

should be a large gradient of FeO in the glass making a good fitting of the thermal

expansion coefficients into the interfacial zone and a sufficient penetration of the

glass in the superficial porosity zone in the alloy.

56

Figure 2.31 Concentration profiles of the iron, nickel, cobalt, silicon, and aluminum

by energy-dispersive x-rays spectrometry (EDS) of the Kovar alloy/glass interface

after thermal treatment for 15 minutes at 910˚C (Zanchetta et al., 1995b).

According to Zanchetta et al. (1995a), the complex mechanism occurred

during the sealing of the preoxidised Kovar alloy with porcelain through a glassy

interphase, by the following steps;

(1) A very rapid diffusion of iron into the glass led to the complete dissolution of

the oxide scale.

(2) If the oxide scale thickness is appropriate, the quantity of iron oxide dissolved

will give the same thermal expansion coefficient in the glass and the

underlying alloy.

(3) The glass penetrates into the open porosity of the alloy and good junction is

achieved.

57

(4) If there is not enough oxide, the fit of thermal expansion coefficients is not

good and the bonding fails. If the time is too long, the bonding is

mechanically very poor and unsticking takes place after cooling.

X-ray diffractograms of the preoxidised surface of the Fe-Ni-Co alloy at

preoxidation time in the range of 1-3 minutes by using a reducing LPG/O2 or C2H2/O2

flame were performed by Piyavit et al. (2006). The results suggested that the oxide

scale consists of hematite (Fe2O3) and magnetite (Fe3O4). An investigation on cross

sections of the joints by SEM backscattered electron imaging and SEM-EDS x-ray

line scanning (Figure 2.32) revealed that the oxide scale, formed by preoxidation with

both LPG/O2 and C2H2/O2 flame at the time of 2 minutes, mostly dissolved into the

borosilicate glass after direct fusion joining. Fe-rich zone in the interfacial glass was

developed. The extent of the Fe-rich zone in the glass was about 40 μm. The results

of this investigation are in agreement with Zanchetta et al. (1995b).

58

Co

Ni

Fe Si

Al

Ι ΙΙ ΙΙΙ ΙV Ι ΙΙ ΙΙΙ ΙV Ι ΙΙ ΙΙΙ ΙV

(a) (b)

Ι ΙΙ ΙΙΙ ΙV

bulk borosilicate glass bulk Fe-Ni-Co alloy

Fe-rich zone with no enrichment in Ni or Co. Depletion of Si and Al was observed.

Complex porous zone with penetration of the Fe-rich interfacial glass. Enrichment in Ni and Co was seemingly occurred.

(c)

Figure 2.32 (a) SEM backscattered electron image; (b) SEM-EDS x-ray line scanning

of the joint (pre-oxidation with LPG/O2 flame for 2 minutes) (Piyavit et al., 2006) and

(c) Schematic drawing of zones at the interface.

59

2.6.2 Devitrification of glass adjacent to the glass/metal interface

Devitrification of the small particles under the size of 1 μm in the glass phase

at the glass-metal boundaries was reported by Ikeda and Sameshima (1964) (Figure

2.33). The excess amount of the oxide and the prolonged sealing time were the cause

of the devitrification of the glass. These particles seemed to be fayalite (Fe2SiO4)

which was determined by the XRD pattern. Moreover, the development of the

crystals at interfaces was very harmful because the crystals were frequently the cause

of hair line cracks in the glass. Devitrification is unfavorable due to its detrimental

effects on oxide dissolution and thermal expansion matching.

Figure 2.33 Devitrification of fayalite in the glass phase (Ikeda and Sameshima, 1964)

Zanchetta et al. (1995a) also reported the devitrification of glass enriched in

iron (and may be cobalt), forming a fayalite-like phase during thermal treatment of

940˚C and soaking times from 8 to 30 minutes (Figure 2.34). However, for well-

chosen times and temperatures, the crystallization rate was slower than iron diffusion

and the oxide scale was entirely dissolved into the glass before any devitrification.

60

Figure 2.34 SEM micrographs of the Kovar alloy-glass interfaces after thermal

treatment for (a) 8 minutes, (b) 15 minutes and (c) 30 minutes at 940˚C, respectively

(Zanchetta et al., 1995a).

For steel-enamel interface, Lupescu et al. (1997) confirmed, by XRD, the

crystallization of the adherence compound, fayalite and magnetite, formed as

interfacial layer, while the FeO was not observed.

Piyavit et al. (2006) also found that the amount of the oxide scale formed by

preoxidation by direct flames for 3 minutes was excessive. The oxide was also

partially dissolved into the borosilicate glass leading to devitrification within the glass

at the vicinity close to the interface (Figure 2.35). The results from energy dispersive

x-ray microanalysis (Figure 2.36) and the x-ray diffractometry confirmed that the

devitrified crystals were fayalite.

61

Co

Ni

Fe Si

Al

Figure 2.35 Devitrification within the glass at the vicinity to the interface

(preoxidation with C2H2/O2 for 3 minutes). SEM backscattered electron image and

SEM-EDS x-ray line scanning of the joint (preoxidation with C2H2/O2 for 3 minutes)

(Piyavit et al., 2006)

Figure 2.36 Point analysis by energy dispersive x-ray microanalysis of devitrified

fayalite in the glass closed to the interfacial zone (preoxidation with C2H2/O2 for 3

minutes) (Piyavit et al., 2006).

Al

62

2.6.3 Microscopical investigation of the glass and metal joining

Electron microscopy (EM) has become an interesting tool for studying phase

transformation and devitrification phenomena in glasses. EM investigation in glasses

may be divided into two major groups (Zarzycki, 1990); (1) low-resolution studies of

texture resulting from phase separation and devitrification process and (2) high-

resolution studies of middle-range order in amorphous phase.

The cross-sectional TEM of the steel-enamel interface was performed by

Liu et al. (1992). It was indicated that the adherence oxide occurring as interfacial

zone is the R3O4 (R represents cations of iron and minor nickel) with the type of

spinel formed as particles about 20 nm in size and intimately mixed with α-Fe. These

particles performed the same single crystal pattern so that the phase has a definite

preferred orientation relationship with the matrix. The selected-area diffraction

pattern (SADP) in Figure 2.37 suggests that both the (010)α-Fe and the (110)oxide phase

are closely parallel. Moreover, TEM study indicated that the adherence spinel

particles are intimately mixed with iron substrate rather than forming a continuous

layer. In the vicinity of the enamel-steel interface, the oxidised steel, rather than

enamel, exhibited dislocation and complicated contrast because amorphous material

does not have lattice strain. The strain contrast of the steel near the interface could be

due to the combined thermal mismatches between α-Fe and enamel, and between α-

Fe and the epitaxial spinel particles.

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Figure 2.37 Schematic indexing of selected area diffraction pattern of epitaxial spinel

particle and α-Fe (Liu et al., 1992)

Two compositions of multicomponent ground enamels coated on a steel

substrate with a thickness higher than 10 mm were studied by Lupescu et al. (1997).

By using XPS analysis of the steel/enamel interfaces, the binding energies of Fe 2p3/2

and O 1s reveal the fayalite and magnetite forming at the steel/enamel interface.

The interface between the borosilicate glass and the Fe-Ni-Co alloy joined by

direct fusion of glass to metal was studied by Chairuangsri et al. (2002). By using

SEM, the joint performed without preoxidation of the alloy revealed the distribution

of the cavities along the alloy grain boundary and in the base alloy grains. While in

the case of preoxidation of the alloy before sealing, no cavities are observed but there

is an oxide layer with the thickness of 3-4 μm. In the case where alloy was not

preoxidised, the EDS line profile revealed Fe enrichment in the glass phase adjacent

to the interface with negligible of Ni and Co content. However, in the case where the

alloy was preoxidised, both Fe and Co diffused into the glass. Recently, the wetting

and sealing of the glass/Kovar interface was reported by Chern and Tsai (2007). The

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Kovar alloy was preoxidised at 700-900˚C for 0-60 minutes. The wetting experiment

was performed at 925˚C for 15 minutes. The glass and metal sealing were joined at

thermal treatment temperature of 825-1000˚C for 15 minutes. Wetting phenomenon,

interfacial observation as well as elemental analysis of the glass-to-metal junctions

were conducted by optical microscope and electron probe microanalysis (EPMA). It

was found that an oxide layer, mainly consisted of FeO, formed at 700˚C for 5-15

minutes with 4-7 μm in thickness is more compatible for good sealing.

2.7 TEM-EELS study of glasses and related oxides

The new generation of field emission electron microscopes equipped with an

energy filter enables excellent spatial as well as energy resolution. This allows the

acquisition of detailed information about the atomic structure, the chemical

composition and the local electronic states of the object. The excitation of atomic

inner shells by high energy electrons provides a method for studying the core-level

reaction. EELS is one of the most important of the energy-loss process to investigate

the local chemistry and bonding in solids. The area under the ionization edges can be

used to extract the chemical composition while threshold shifts are induced by

different bonding coordination and charge states. The shapes of the ionization edges

in EELS spectra reflect the available excited states which depend on the local bonding

environments.

EELS has been used to study various oxides especially of the transition

metals. For example, Colliex et al. (1991) reported the EELS experiment in the Fe-O

system in thin-film configuration of the characteristic O K and Fe L23-edges in FeO,

Fe3O4, γ-Fe2O3 and α-Fe2O3. This work confirmed the stability of the L3-L2 energy

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difference at about 13.1 ± 0.2 eV for the various phases of these oxides. The Fe L2-

edges obtained from TEM-ELNES were used to identify the characteristic of Fe

valence states in minerals (van Aken et al., 1998; van Aken and Liebscher, 2002). It

was mentioned that the Fe L3-edge of divalent iron (Fe2+) appeared as a white line at

707.8 eV followed by the broadening peak at 710.5 eV, while the Fe L3-edge ELNES

for the trivalent iron (Fe3+) consists of a white line at 709.5 eV. The spin-orbit

splitting reveals the separation of Fe L3- and L2- edges of 12.8 ± 0.1 eV and 13.2 ± 0.1

eV for the Fe2+ and Fe3+, respectively. The white line intensity of L3/L2 ratio also

allows the quantification of ferrous/ferric ratio in minerals. With a well-prepared

specimen and the nanometer-scale spatial resolution TEM, the Fe L3-edge ELNES

gave an advantage over the EPMA and x-ray absorption technique (van Aken et al.,

1998).

The EELS analysis for quantitative determination of the cation valence states

of Mn and Co oxides by identifying the L23-ionization edges of the various magnetic

oxides including CoO and Co3O4 were observed by Wang et al. (2000). Mn and Co

are the only transition metal elements whose the L3/L2 ratio is most sensitive to

valence state variation, while the white line of Fe is almost independent. However, it

should be noted that the determination of the crystal structure of nanoparticles require

high quality EELS spectra especially when the particles are less than 5 nm, because

the peak is broadened due to the crystal shape factor. By using TEM equipped with a

monochromator and a high-resolution imaging filter, the ELNES of the metal L23- and

O K-edges in a well-defined of 3d transition metal oxide such as TiO2, V2O5, Cr2O3,

Fe2O3, CoO and NiO have been measured with a new system (Mitterbauer et al.,

2003). It was concluded that the multiplet structure of the L23-edges are different due

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to both solid state (crystal field splitting) and atomic (Coulomb and exchange

interaction) of transition metal oxides.

Based on chemical consideration of bridging and non-bridging oxygen (NBO)

in oxide glasses, the EELS analysis is one of the techniques which were used for

identified the oxygen configuration of the O K-edge in a Ca aluminosilicate glass

(Jiang, 2002). In silicate glass, the NBO is bonded to only one Si atom, while the

bridging oxygen links with two Si tetrahedra. Since the bonding configuration of the

oxygen in the Ca aluminosilicate glass have at least three types; Si-O-Si, Si-O (NBO)

and Si-O-Al, therefore, the electron energy loss near edge fine structure (ELNES) of

the O K-edge can be used to identify the local structures of oxygen in silicate glass.

The advantages of EELS are also required for the various types of interfacial

investigation. Brydson et al. (1995) studied the bonding of Fe/Al2O3 and Ni/Al2O3

interfaces by using various techniques such as high-resolution electron microscopy

(HREM), scanning transmission electron microscopy (STEM) and parallel electron

energy-loss spectroscopy (PEELS). It was found that HREM and EELS confirmed

the appearance of Al2O3 amorphous phase at the area vicinity closed to the metal-

ceramic interface. STEM-point analyses also confirmed the formation of an iron-rich

spinel (FexAl1-x)3O4 (with x ≈ 0.5) interphase in the region between Fe and Al2O3.

Similarly to Brydson et al. (1995), HREM and EELS characterization of the atomic

and electronic structure in the system of Cu/α-Al2O3 interfaces was carried out by

Sasaki et al. (2005). HREM revealed the orientation relationship (OR) of the (111)

and (001) plans of Cu epitaxially oriented to the (0001) and (11−

2 0) of Al2O3,

respectively. In both cases of the Cu/Al2O3(0001) and Cu/Al2O3(11−

2 0) interfaces,

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the ELNES of O K-edges at the interface indicated the shoulder at about 533 eV in the

main broad peak around 528 eV. The appearance of the shoulder peak indicated that

the electronic states in the interface are different from the bulk Al2O3 which can effect

from the Cu-O interactions. EELS was therefore confirmed the attribution of Cu-O

bonding across the interface in both cases. Furthermore, it can be considered that the

bonding state at the interface play an important role for the difference in ORs.