2.1 Fe-Ni-Co alloy
Transcript of 2.1 Fe-Ni-Co alloy
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
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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%.
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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
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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).
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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.
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Figure 2.2 Iron-nickel-cobalt equilibrium diagrams at 500 and 800˚C (Butteridge,
1984).
Fe
Co
Ni
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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
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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
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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)
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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.
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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).
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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.
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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
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* 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.
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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.
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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.
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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).
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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
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(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).
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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.
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Figure 2.10 Variations of viscosity with temperature of some commercial glasses
(Pfaender, 1983).
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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.
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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)
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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)
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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
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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.
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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)
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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.