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Effect of composition on the properties of glasses in the K2O–BaO–MgO–SiO2–Al2O3–B2O3–MgF2 system
Journal of Non-Crystalline Solids 325 (2003) 193–205 www.elsevier.com/locate/jnoncrysol
Effect of composition on the properties of glasses in the K2O–BaO–MgO–SiO2–Al2O3–B2O3–MgF2 system K. Greene a
a,b
, M.J. Pomeroy
a,b,*
, S. Hampshire
a,b
, R. Hill
c
Department of Materials Science and Technology, University of Limerick, Limerick, Ireland b Materials and Surface Science Institute, University of Limerick, Limerick, Ireland c Materials Department, Imperial College of Science and Technology, London W2 3SF, UK Received 11 July 2002
Abstract Twenty five glasses in the(1 Z)BaO:ZK2 O:(6 X )MgO:X MgF2 :(3 Q)Al2 O3 :QB2 O3 :8SiO2 system have been prepared and the effect of systematic changes in composition (Z ¼ 0, 0.25, 0.5, 0.75 and 1.0, X ¼ 2, 2.5 and 3.0 and Q ¼ 0, 0.5 and 1) on molar volume (MV), fractional glass compactness (C), microhardness (lHV ), glass transition temperature (Tg mid ) and coefficient of thermal expansion (a) determined. As potassium is substituted for barium (increasing Z), increases in MV and a, and decreases in C, lHV and Tg mid arise which are attributed to the replacement of ONB –Ba–ONB ionic bridges by two ONB –K terminations (ONB ¼ non-bridging oxygen). When fluorine is substituted for oxygen, reductions in lHv and Tg mid and increases in MV and a values occur. These property changes occur because of reduced crosslink densities associated with the replacement of Al–O–Si crosslinks by Al–F terminations. In general, the substitution of aluminium by boron results in decreases in MV, lHV and Tg mid values and increases in a values. These effects are attributed to boron assuming a tri-coordinated state in the glass, giving rise to reduced crosslink densities, together with the release of modifier cations (Mg, K and Ba) from their charge balancing role for four-coordinated aluminiums leaving them free to cause greater network disruption. 2003 Elsevier B.V. All rights reserved.
1. Introduction Glasses based on the system SiO2 –Al2 O3 –B2 O3 – MgO–MgF2 –K2 O typically crystallise to trisilicic alkaline mica (KMg3 AlSi3 O10 F2 ). Such glassceramics are an important class of material as they are machinable to high (±10 lm) tolerances [1]. The microstructure of these materials, comprising highly interlocked K-fluorophlogopite crystals
*
Corresponding author. Tel.: +353-61 202 200. E-mail address: [email protected] (M.J. Pomeroy).
with a Ôhouse of cardsÕ structure embedded in a glass matrix, facilitates microfracture along the weak mica–glass interface and mica basal planes, which prevents macroscopic failure during machining. Mica glass-ceramics containing Ba-fluorophlogopite crystals are attracting interest because they have shown improvements in mechanical properties when compared with more conventional Kfluorophlogopite glass-ceramics [2,3]. However, very little is known about the effect which exchanging barium for potassium has on the properties of the base glasses. Wallace and coworkers
0022-3093/$ - see front matter 2003 Elsevier B.V. All rights reserved. doi:10.1016/S0022-3093(03)00337-5
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[4,5] have shown that replacing an alkaline earth cation such as Ca2þ with an alkali cation such as Naþ in bioactive glasses reduces both the glass transition temperature and the glass density. These reductions are related to decreases in levels of association between adjacent network regions due to the elimination of ionic bridges afforded by divalent cations and non-bridging oxygens (ONB – MII –ONB ) and their replacement by two terminal ONB –MI entities. However, the work of Dietzel [6,7] indicates that the effect of alkali/alkaline earth cation exchange on the properties of glasses can also be related to cationic field strengths. The cationic field strength is a measure of a cationÕs effective force in attracting anions and is given by z=r2 where z is valency and r the ionic radius of the cation. In Ln–Si–Al–O–N glasses (Ln ¼ Y, Ce, Nd, Sm, Eu, Dy, Ho and Er), the variation in glass transition temperatures for glasses of constant composition has been directly related to the cationic field strength of the lanthanide cation being substituted [8]. Thus, it has been reported that glass transition temperatures increase linearly with increasing cation field strengths of the rare earth modifiers. Investigations into the effect of fluorine on the properties of glasses have shown that it is network disruptive [9–11]. The reduction in glass transition temperature of pre-cerammed cannasite glassceramics has been related to the creation of nonbridging fluorines which replace bridging oxygens in the glass network [9]. In apatite–mullite glassceramics [10] and ionomer glass compositions [11], the glass transition temperature decreases with increasing fluorine content as a result of network disruption by fluorine. There is, unfortunately, a major problem with glass systems containing fluorine that is the volatilisation of silicon tetrafluoride (SiF4 ) during glass synthesis which can lead to fluorine deficient glasses and increased production costs. However, Hill et al. [11] have shown that volatilisation can be inhibited by incorporating appropriate amounts of aluminium and basic network modifying oxide to allow aluminium to achieve fourfold coordination. Thus, if the Al:F ratio is greater than 1:1 fluorine loss can be obviated. Boric oxide is typically included in the starting composition of machinable glass ceramics in order
to increase the aspect ratio of the resultant fluorophlogopite crystals [12] which yields a glass-ceramic with improved machinability [13]. However, the effect of boric oxide on the properties and structure of the uncerammed glasses has not been investigated. Research on the structure of boric oxide containing glasses has generally concentrated on the Na2 O–B2 O3 –SiO2 system [14] and to a lesser extent the Na2 O–B2 O3 –Al2 O3 –SiO2 system [15] used in the production of two phase glasses which can subsequently treated to yield porous materials. Even in these relatively simple systems, detailed relationships between the composition, structure and properties of the glasses have not been derived. The structure of aluminosilicate and boroaluminosilicate glasses can be described by applying an inorganic polymer model developed by Ray [16,17]. From studying a typical oxide network such as vitreous silica, Ray postulated that inorganic glasses could be regarded as polymer chains with networks of oxygen atoms linked by mainly covalent bonds through multivalent atoms of network-forming elements. An important factor relating the properties of an inorganic glass to its composition and structure are the density of crosslinks in the structure. The crosslink density can be defined as the average number of ÔbridgingÕ oxygen atoms in excess of two per network former, where bridging oxygens are defined as those that link together two other atoms that are part of the network, e.g. Si, B, Al. If there are not more than two bridging oxygens per network former atom, the structure is equivalent to that of a linear polymer and will therefore have a crosslink density of zero. The crosslink density of pure silica is two which represents a fully crosslinked inorganic polymer. The effect of glass forming, intermediate and modifying ions on the crosslink density can be calculated from the molar glass composition. Given the absence of clear information relating to glasses in the (Ba,K)–Mg–Si–(Al,B)–(O,F) system, the aim of this work was to systematically investigate the effects of substituting potassium for barium, fluorine for oxygen and boron for aluminium on their physical and thermal properties and from the property trends observed, infer possible structural changes arising in the glasses.
K. Greene et al. / Journal of Non-Crystalline Solids 325 (2003) 193–205
2. Experimental 2.1. Glass design In designing this series of novel glass compositions, the following considerations were taken into account: ii(i) The glasses should have enough of each specified compound to ensure that preferential crystallisation to fluorophlogopite would occur. i(ii) The Si:Al ratio was maintained >1:1 in order to adhere to the ÔLowenstein aluminium avoidance principleÕ [18], allowing Al to take up a predominantly fourfold coordination state, thereby favouring glass formation during synthesis. (iii) The Al:F ratio was maintained P 1:1, with enough modifier(s) present to ensure that Al is bonded to fluorine and that local electroneutrality around the Al3þ cation is maintained. This was expected to significantly reduce fluorine loss during glass synthesis [11]. (iv) The need to study, systematically, the effects of the following substitutions on glass properties: (a) K for Ba, (b) F for O and (c) B for Al. The glasses studied were therefore based on the following parent composition: 8SiO2 ð3 QÞAl2 O3 QB2 O3 ð6 X ÞMgO X MgF2 ð1 ZÞBaO ZK2 O; where Z values of 0.0, 0.25, 0.5, 0.75 and 1.0, X values of 2.0, 2.5 and 3.0, and Q values of 0.0, 0.5 and 1.0 were investigated. The minimum X value was limited to 2.0 as this corresponded to a glass having a stoichiometry similar to that of fluorophlogopite whilst the maximum X value was limited to 3.0 as it was expected that the amount of silicon tetrafluoride loss during synthesis would become significant since for X > 3:0 the Al:F ratio would be less than 1:1 [11].
195
oxide (B2 O3 ), magnesium carbonate (MgCO3 ), magnesium fluoride (MgF2 ), potassium carbonate (K2 CO3 ) and barium carbonate (BaCO3 ) powders, thoroughly mixed by ball milling. The purity and supplier of each reagent is shown in Table 1. After mixing, the powders were placed in sintered mullite crucibles (Zedmark Refractories, Earlsheaton, Dewsbury, UK) and fired in an electric furnace (Carbolite HTC 1600, Sheffield, UK) at temperatures between 1400 and 1450 C for 1.0– 1.5 h depending on the composition. Batch weights prior to firing were approximately 450 g. The glass melts were subsequently shock quenched into water to produce glass frit. This frit was dried in a fan oven at 100 C for 24 h, ground in a vibratory mill (Gyro-Mill, Glen Creston Ltd., UK) and sieved to produce three particle size fractions (i) <45 lm, (ii) >45 < 105 lm (coarse) and (iii) >105 lm (frit). Cast glass was produced by remelting frit particles in alumina crucibles (VZS, Fife, Scotland, UK) at 1450 C for 1.5 h in an electric furnace (Carbolite RHF 1500, Sheffield, UK). The molten glass produced was poured into preheated graphite crucibles, annealed for 1 h at a temperature 50 C below the midpoint of the glass transition region and allowed to furnace cool to room temperature overnight. The glass rods produced were 12 mm in diameter and 40 mm in length.
Table 1 Materials and suppliers Component
Percentage of purity
Supplier
SiO2
99.9
Al2 O3
99.86
MgCO3
99.85
MgF2
99.9
B2 O3
99.8
2.2. Glass preparation
BaCO3
99.9
The glasses were synthesised using batches of reagent grade silica (SiO2 ), alumina (Al2 O3 ), boric
K2 CO3
99.9
Tilcon Industrial Minerals, UK GPR Grade, BDH Laboratory Supplies, UK Alfa, Johnson and Matthey, Germany Alfa, Johnson and Matthey, Germany Aldrich Chemical Company Inc, USA Alfa, Johnson and Matthey, Germany E. Merck, Darmstadt, Germany
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2.3. Glass characterisation X-ray diffraction (XRD) data for each of the glasses was recorded using a Philips X-ray diffractometer (Philips XÕPert, Eindhoven, NL) using ), an anode current Cu Ka radiation (k ¼ 1:5406 A of 30 mA and a voltage of 40 kV. Diffraction patterns were collected between 5 and 70 2h. XRD was carried out on both <45 lm size fraction particles and also cast specimens to ensure that the glasses were completely amorphous. The densities of pore free cast glass rods were measured using a water-density balance. Samples, 12 mm in diameter and 10 mm in length, were weighed in both air and distilled water using a Sartorius balance. The densities were calculated using the following equation: q ¼ qwðT Þ Ws =ðWs Ww Þ; where q is the sample density (g cm3 ), qwðT Þ the density of water at the measured temperature (g cm3 ), Ws the sample weight in air (g) and Ww the sample weight in water (g). The accuracy of this technique is ±0.01 g cm3 . Molar volumes were calculated from the density data using the formula P xi M i MV ¼ ; q where xi and Mi represent the atomic fractions and atomic weights of K, Ba, Mg, Si, Al, B, O and F respectively and q is the measured glass density (g cm3 ). Fractional glass compactness (C) values were calculated using the expression C ¼ Vion
total =MV;
where Vion total ¼ xi Vi and Vi is the ionic volume calculated using ionic radii data published in [19]. Microhardness tests were carried out on polished glass samples using a Leco microhardness tester (Akashi Corporation, Japan). A load of 300 g was applied to the sample for 15 s. The resulting indentation diagonals were measured and the hardness was calculated using the following equation: lHV ¼ 1:854P =ðd 2 Þ; where lHV is the Vickers microhardness in GPa, P is the load (N) and d is the mean diagonal of the microhardness indentation (m).
Differential thermal analysis (DTA; Stanton Redcroft STA 1640, Rheometric Scientific, Epsom, UK) was used to determine the midpoint of the glass transition temperature range (Tg mid ). DTA experiments were carried out using a heating rate of 10 C/min in a flowing nitrogen atmosphere with matching pairs of platinum–rhodium alloy crucibles and 50 mg alumina as the reference material. Measurements to determine coefficients of thermal expansion (a) were carried out on cast glasses using a Netsch DIL 402 dilatometer. Samples 10 mm · 4 mm · 4 mm were heated at a rate of 5 C/min in air and a was calculated in the range 200–500 C using the following equation: a ¼ ðDl=lo Þ=DT ; where lo is the original length, Dl the change in length of the specimen and DT the temperature change.
3. Results 3.1. General observations Weight change measurements following glass synthesis showed weight losses consistent with the loss of carbon dioxide from the magnesium, potassium and barium carbonates used in glass preparation. Because of this, it is assumed that all fluorine added to the glass formulation is present in the glass after firing. Furthermore, no weight losses were recorded after the casting of glass bars for property measurements, indicating that no fluorine losses occurred. XRD analysis of all glasses, as melted and fritted and as cast, showed that all materials were amorphous. 3.2. Boron free glasses (Q ¼ 0) Table 2 gives the relevant glass compositions and results for the property measurements for this series of glasses where Z and X were varied. It is seen from Table 2 that as barium is replaced by potassium (increasing Z) then glass densities decrease from 2.90 to 2.63 g cm3 for each fluorine content. Whilst these density decreases may be associated with the lower atomic weight of potas-
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Table 2 Glass compositions, property values and calculated crosslink density (CLD) values for boron free glasses Glass code GðX ; Q; ZÞ
Density (g cm3 )
Molar volume (cm3 mol1 )
Microhardness (GPa)
Tg mid (C)
CTE (106 K1 )
G(2.0, G(2.0, G(2.0, G(2.0, G(2.0,
0.0, 0.0, 0.0, 0.0, 0.0,
0.00) 0.25) 0.50) 0.75) 1.00)
2.91 2.84 2.76 2.71 2.64
7.66 7.72 7.81 7.82 7.89
6.59 6.53 6.39 6.38 6.11
679 671 665 660 653
6.00 5.98 6.24 6.54 6.65
1.14 1.14 1.14 1.14 1.14
G(2.5, G(2.5, G(2.5, G(2.5, G(2.5,
0.0, 0.0, 0.0, 0.0, 0.0,
0.00) 0.25) 0.50) 0.75) 1.00)
2.90 2.84 2.77 2.69 2.63
7.68 7.72 7.78 7.88 7.93
6.54 6.43 6.28 6.19 6.14
653 650 645 642 635
6.31 6.34 6.48 6.56 7.00
1.07 1.07 1.07 1.07 1.07
G(3.0, G(3.0, G(3.0, G(3.0, G(3.0,
0.0, 0.0, 0.0, 0.0, 0.0,
0.00) 0.25) 0.50) 0.75) 1.00)
2.90 2.83 2.77 2.69 2.63
7.68 7.75 7.78 7.88 7.93
6.44 6.30 6.20 6.13 5.95
638 638 639 636 620
6.25 6.45 6.50 6.60 7.12
1 1 1 1 1
8.00 7.95
3
-1
molar volume (cm .mol )
sium, which is significantly lower than that of barium (39.1 compared with 137.3), increases in glass volume also arise. Fig. 1 shows the variation in molar volume with increasing levels of potassium substitution for barium and indicates that as barium ions are replaced by two potassium ions, molar volume increases in a reasonably linear manner. Fig. 2 clearly shows that the compactness of the glass decreases with increasing potassium for barium substitution indicating that the glass structure expands due to changes in the strength of modifier – glass structural unit interaction. The effect of increased fluorine substitution for oxygen has virtually no effect on density (see Table 2). Effects of this increased substitution on molar volume can be interpolated from Fig. 1 which shows that molar volume changes with increased X value are almost within experimental error (±0.03 cm3 mol1 ). That said, there appears to be an increase in molar volume associated with the initial increase in fluorine content (X ¼ 2:0–2.5) but values for the X ¼ 2:5 and 3.0 substitution levels are almost identical. Microhardness values decrease linearly with increasing substitution of potassium for barium substitution levels (see Fig. 3). In addition, it can be seen from both Table 2 and Fig. 3 that as fluorine substitution levels increase microhardness values decrease.
CLD
7.90 7.85 7.80 7.75 7.70 7.65 7.60 0
0.25
0.5
0.75
1
Z value Fig. 1. Effect of potassium substitution for barium (Z) on molar volume of boron free (Q ¼ 0) glasses with different fluorine substitution levels ( (X ¼ 2:0), M (X ¼ 2:5), } (X ¼ 3:0)).
Glass transition range midpoint temperatures (Tg mid ) data for frit samples are given in Table 2. Fig. 4 shows that whilst they are lowered by potassium substitution, the extents of these decreases are dependant on fluorine substitution level. Thus,
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690
glass transition temperature (°C)
fractional glass compactness
0.585
0.580
0.575
0.570
680 670 660 650 640 630 620 610
0.565 0
0.25
0.5
0.75
1
0
0.25
Fig. 2. Effect of potassium substitution for barium (Z) on fractional glass compactness of boron free (Q ¼ 0) glasses with different fluorine substitution levels ( (X ¼ 2:0), M (X ¼ 2:5), } (X ¼ 3:0)).
6.9
microhardness (GPa)
6.7 6.5 6.3 6.1 5.9 5.7 0.25
0.5
0.75
1
Z value
Z value
0
0.5
0.75
1
Z value Fig. 3. Effect of potassium substitution for barium (Z) on microhardness of boron free (Q ¼ 0) glasses with different fluorine substitution levels ( (X ¼ 2:0), M (X ¼ 2:5), } (X ¼ 3:0)).
for X ¼ 2:0 and X ¼ 2:5 glasses, Tg mid values decrease linearly with increasing potassium for barium substitution (Z), but for X ¼ 3:0 glasses glass
Fig. 4. Effect of potassium substitution for barium (Z) on glass transition temperature of boron free (Q ¼ 0) glasses with different fluorine substitution levels ( (X ¼ 2:0), M (X ¼ 2:5), } (X ¼ 3:0)).
transition temperatures are almost constant up to 75% substitution of potassium for barium. The effect of increasing fluorine substitution levels is to cause significant decreases in Tg mid as X is increased from 2.0 to 2.5. However, for further fluorine substitution (X ¼ 2:5–3:0) decreases are less significant, indicating that the degradation of this property, induced by fluorine, is diminished, except when all barium is substituted by potassium. Fig. 5 shows the effect of Z value on coefficient of thermal expansion (CTE) and indicates that as potassium replaces barium then CTE values also increase. Furthermore, it can also be seen from Fig. 5 that CTE values for the same Z value are only affected by initial increases in fluorine substitution levels (X ¼ 2:5). Further increases (X ¼ 3:0) have little effect. 3.3. Glasses (X ¼ 2:5)
with
constant
fluorine
content
Table 3 shows the compositional variations used to investigate the effects of variations in potassium for barium and boron for aluminium substitution on glass properties, together with
K. Greene et al. / Journal of Non-Crystalline Solids 325 (2003) 193–205
7.1
-6
-1
CTE (10 K )
7.6
6.6
6.1
5.6 0
0.25
0.5
0.75
1
Z value Fig. 5. Effect of potassium substitution for barium (Z) on coefficient of thermal expansion of boron free (Q ¼ 0) glasses with different fluorine substitution levels ( (X ¼ 2:0), M (X ¼ 2:5), } (X ¼ 3:0)).
property values and crosslink density values for the glasses assuming boron is in either three or fourfold coordination. It is seen from Table 3 that for each boron substitution level (Q ¼ 0, 0.5 and 1.0), densities of glasses decrease as levels of po-
199
tassium for barium substitution increase. The extents of the increases decrease with increasing boron substitution for aluminium. Furthermore, it is seen that for constant Z values the effect of boron substitution for aluminium has a minor effect on density with density decreases of the order of 0.02–0.03 g cm3 accompanying the substitution of one mole of alumina by one mole of boric oxide. Fig. 6 shows the effect of Z value on molar volume for each boron substitution level. It is seen that, for all boron levels, molar volumes increase with increasing potassium for barium substitution. The effect of increasing boron substitution on molar volume is not significant as differences are within the calculation error of ±0.03 cm3 mol1 . Fig. 7 indicates that fractional glass compaction values decrease with increasing potassium substitution levels. Again the effects of aluminium substitution by boron are within the calculation error band (±0.003). Microhardness values for these glasses decrease linearly with increasing potassium substitution level for the Q ¼ 0 and Q ¼ 0:5 compositions (see Fig. 8). However, for the Q ¼ 1:0 substitution level the microhardness values of the baria rich glasses (Z ¼ 0 and Z ¼ 0:75) are similar to each other and to the values for the Q ¼ 0:5 substitution. The
Table 3 Glass compositions, property values and calculated crosslink density values for glasses with fixed (X ¼ 2:5) fluorine content Glass code GðX ; Q; ZÞ
Density (g cm3 )
Molar volume (cm3 mol1 )
Microhardness (GPa)
Tg mid (C)
CTE (106 K1 )
CLD B in fourfold
B in threefold
G(2.5, G(2.5, G(2.5, G(2.5, G(2.5,
0.0, 0.0, 0.0, 0.0, 0.0,
0.00) 0.25) 0.50) 0.75) 1.00)
2.90 2.84 2.77 2.69 2.63
7.68 7.72 7.78 7.88 7.93
6.54 6.43 6.28 6.19 6.14
653 650 645 642 635
6.31 6.34 6.48 6.56 7.00
1.07 1.07 1.07 1.07 1.07
1.07 1.07 1.07 1.07 1.07
G(2.5, G(2.5, G(2.5, G(2.5, G(2.5,
0.5, 0.5, 0.5, 0.5, 0.5,
0.00) 0.25) 0.50) 0.75) 1.00)
2.87 2.80 2.72 2.67 2.61
7.66 7.72 7.82 7.83 7.88
6.23 6.17 6.10 6.06 5.88
627 622 624 624 610
6.53 6.51 6.63 6.91 7.53
1.07 1.07 1.07 1.07 1.07
0.96 0.96 0.96 0.96 0.96
G(2.5, G(2.5, G(2.5, G(2.5, G(2.5,
1.0, 1.0, 1.0, 1.0, 1.0,
0.00) 0.25) 0.50) 0.75) 1.00)
2.85 2.76 2.69 2.64 2.58
7.61 7.73 7.80 7.81 7.86
6.17 6.16 5.94 5.78 5.70
614 611 606 600 581
6.28 6.33 6.68 6.83 7.46
1.07 1.07 1.07 1.07 1.07
0.85 0.85 0.85 0.85 0.85
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K. Greene et al. / Journal of Non-Crystalline Solids 325 (2003) 193–205
6.9
8.00
6.7
7.90
microhardness (GPa)
3
-1
molar volume (cm .mol )
7.95
7.85 7.80 7.75 7.70
6.3 6.1 5.9
7.65
5.7
7.60 0
0.25
0.5
0.75
1
Fig. 6. Effect of potassium substitution for barium (Z) on molar volume of glasses containing 2.5 mol MgF2 (X ¼ 2:5) for different boron substitution levels ( (Q ¼ 0:0), N (Q ¼ 0:5), r (Q ¼ 1:0)).
0.5850
0.5800
0.5750
0.5700
0.5650 0
0.25
0.5
0
0.25
0.5
0.75
1
Z value
Z value
fractional glass compactness
6.5
0.75
1
Z value Fig. 7. Effect of potassium substitution for barium (Z) on fractional glass compactness of glasses containing 2.5 mol MgF2 (X ¼ 2:5) for different boron substitution levels ( (Q ¼ 0:0), N (Q ¼ 0:5), r (Q ¼ 1:0)).
effect of increasing boron substitution level up to an initial level of Q ¼ 0:5 is to decrease microh-
Fig. 8. Effect of potassium substitution for barium (Z) on microhardness of glasses containing 2.5 mol MgF2 (X ¼ 2:5) for different boron substitution levels ( (Q ¼ 0:0), N (Q ¼ 0:5), r (Q ¼ 1:0)).
ardness by 0.2–0.3 GPa. For the Q ¼ 1 substitution level, similar effects arise for the glasses with Z ¼ 0:5, 0.75 and 1.0 but not for the Z ¼ 0 and 0.25 compositions where hardness values are unaffected. The effect of potassium for barium substitution on glass transition temperature is shown in Fig. 9. It is seen that for each boron substitution level there is a decrease in Tg mid , although the effect of Z value on Tg mid for glasses with boron substitution levels corresponding to Q ¼ 0:5 is negligible until complete substitution of barium by potassium (Z ¼ 1:0) is reached. In addition, it can be seen from Fig. 9 that, as boron substitution for aluminium increases, Tg mid values decrease quite markedly. For glasses with Z ¼ 0:5, 0.75 and 1.0, the decreases in Tg mid with Q are linear. However, for the other Z value glasses the decreases in Tg mid are less as boron is substituted for aluminium. Coefficients of thermal expansion (see Table 3 and Fig. 10) increase with increasing levels of potassium substitution, for each boron content. As the level of boron substitution is increased from Q ¼ 0 to 0.5, CTE values also increase. For further increases in boron substitution, the effect on CTE
K. Greene et al. / Journal of Non-Crystalline Solids 325 (2003) 193–205
pansion are similar to those containing no boron and for lower Z (high Ba) values the values of CTE are similar to those for glasses with a boron substitution level corresponding to Q ¼ 0:5.
glass transition temperature (°C)
660 650 640 630
4. Discussion of results
620 610 600 590 580 570 0
0.25
0.5
0.75
1
Z value Fig. 9. Effect of potassium substitution for barium (Z) on glass transition temperature of glasses containing 2.5 mol MgF2 (X ¼ 2:5) for different boron substitution levels ( (Q ¼ 0:0), N (Q ¼ 0:5), r (Q ¼ 1:0)).
7.6
7.1
-6
-1
C TE (10 K )
201
6.6
6.1
5.6 0
0.25
0.5
0.75
1
Z value Fig. 10. Effect of potassium substitution for barium (Z) on coefficient of thermal expansion of glasses containing 2.5 mol MgF2 (X ¼ 2:5) for different boron substitution levels ( (Q ¼ 0:0), N (Q ¼ 0:5), r (Q ¼ 1:0)).
depends upon Z value. Thus, for the high Z value (high K) glasses with one third of the aluminium substituted by boron coefficients of thermal ex-
In general, for each of the boron free glass series in which potassium for barium substitution was investigated, density, fractional glass compactness, microhardness and glass transition temperature values decreased, whilst molar volume and coefficient of thermal expansion values increased. Each of these property trends indicates two structural effects. One is that barium and potassium ions act as network modifiers rather than charge balancing aluminium in the network. The other is that the substitution of potassium for barium results in a decrease in the strength of association between modifier ions and their surrounding elements of the glass network. It is of course non-bridging oxygens (ONB ) which are most strongly coordinated to modifier cations. As indicated in the introduction, the effects of alkali/alkaline earth modifier cation substitutions on glass structure, and therefore properties, can either be viewed from the point of view of a sequential replacement of ONB –MII –ONB bridges with two ONB –MI terminations [4,5] or changes in cation field strength [6–8]. If changes in cation field strength for the substitution of Ba2þ by 2Kþ , are examined (using ionic radii values given by Chiang et al. [19]) then very little change is observed, since the cation field strength of Ba2þ is 110 nm2 and that of 2Kþ is 105 nm2 . Because of this, it appears that the sequential replacement of ONB –Ba–ONB bridges with two ONB –K terminations is responsible for the degradation in properties observed. Such replacements would certainly result in a decrease in the degree of association of the glass network which would facilitate decreases in fractional glass compactness and hardness as the rigidity with which adjacent parts of the glass network could be associated would be substantially decreased. A similar argument would be expected to apply to decreases in glass transition temperatures since segments of the glass would be expected to be able
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to move more readily as the degree of network association decreases due to ONB –K termination effects. With respect to glass transition temperatures, there is a clear anomaly in the effect of Z for the glasses containing the maximum level of fluorine substitution (X ¼ 3). This is discussed later. Increasing thermal expansion values may also be attributable to decreased association levels. However, the situation is somewhat more complex since this property is also related to non-symmetrical vibrations of ions and structural entities, the exact analysis of which is outside the scope of this work. For the glasses containing boron with fluorine contents corresponding to X ¼ 2:5, the general effect of potassium substitution for barium is virtually the same as that for the boron free glasses. Accordingly, the same explanations can be invoked as for boron free glasses. Anomalous property trends (e.g. Tg mid for Q ¼ 0:5 glasses) are addressed below in the paragraphs dealing with the effects of boron content. The effects of increasing fluorine substitution for oxygen on glass properties are similar to those for the substitution of barium by potassium, i.e. all property values disimprove. However, the reasons for the property degradation induced by increased levels of fluorine substitutions are different as fluorine resides in the glass network. As oxygen is replaced by fluorine, potential Al–O–Si crosslinks in the network are replaced by Al–F terminations thus reducing crosslink densities as indicated in Table 2. The property disimprovements observed are therefore consistent with decreased crosslink density trends as has been previously reported for other fluorine containing glass systems [9–11,20–22]. It is clear from the results presented, that, with the exception of hardness, the property values for glasses containing 2.5 and 3.0 mol of MgF2 are very similar for the same potassium content and that glass transition temperatures for the glasses with X ¼ 3:0 are independent of potassium/barium content until all barium is replaced. Given that no weight loss attributable to fluorine was observed for these glasses and the Al:F ¼ 1:1 rule of Hill et al. [11] was adhered to, the level of fluorine substitution can be safely assumed to be that present in the starting glass composition. Thus,
fluorine loss is not the reason for the phenomena observed. Two other possible explanations arise. One involves separation of the glass into a fluorine and modifier (K, Ba and Mg) rich phase and an aluminosilicate rich phase as proposed by Hill et al. [11], Rabinovich [23] and Maeda et al. [24], whilst the other may be related to levels of glass network disruption reaching a maximum at the fluorine content corresponding to X ¼ 2:5. With respect to the possibility of a two phase glass, additional work has shown that the optimum nucleation temperatures of the glasses investigated here are close to their glass transition temperatures indicating that they nucleate via a phase separation process and thus that chemical inhomogeneity may exist in the glasses prior to their nucleation. If this is the case, then the overall glass microstructure would comprise a modifier plus fluorine rich droplet phase in a continuous modifier and fluorine depleted aluminosilicate phase. This might well explain the distinct similarities in molar volume and thermal expansion coefficient values for glasses with X ¼ 2:5 and 3.0, since if much of the excess fluorine (i.e. that corresponding to 0.5 mol MgF2 ) is accommodated in the droplet phase then the matrix phase which controls these particular properties of the glass would be very similar to the bulk glass corresponding to X ¼ 2:5. Such a glass structure would also be consistent with the effects of increased fluorine levels on microhardness. This is because the microhardness values would represent those obtained from a droplet phase–matrix phase composite and as the modifier plus fluorine rich droplet phase would be comparatively soft the overall microhardness values for X ¼ 3:0 glasses would be less than for the X ¼ 2:5 glasses. However, it must be noted that no evidence of modifier plus fluorine crystalline phases were observed during X-ray diffraction and thus the modifier plus fluorine rich droplet phase must be either amorphous or present in very small levels (<2–3%). The alternative explanation is that as the fluorine level is increased from X ¼ 2:0 to 2.5, the extent of network disruption reaches a maximum. This would result in subsequent fluorine for oxygen substitution (X ¼ 3:0) having no further effect on properties such as molar volume and thermal expansion coefficient. Such an argument is how-
K. Greene et al. / Journal of Non-Crystalline Solids 325 (2003) 193–205
ever at variance with the crosslink density data given in Table 2 and therefore it strongly appears that the similarity in molar volume and coefficient of thermal expansion values between glasses with fluorine contents corresponding to X ¼ 2:5 and 3.0 is due to the location of the extra fluorine in a modifier rich second phase. For the barium rich boron containing glasses with low Z values (0 and 0.25), increased levels of boron substitution (Q ¼ 0:5–1:0) have no effect on microhardness and decreases in Tg mid are less than for the potassium rich higher Z value glasses. In addition, the coefficients of thermal expansion for the Q ¼ 0 and Q ¼ 1 substitution levels are similar. For all other Z values, the effect of increasing Q value on glass properties is degradative. The most plausible explanation for the anomalous effects of boron for aluminium substitution is that during the melting and subsequent bar casting of these low Z-value glasses, fluorine loss occurred with the result that the property changes due to boron substitution were negated by fluorine loss. This reasoning would certainly be consistent with the weight losses observed for these glasses and the fact that the Al:F ratio for the Q ¼ 1:0 substitution level is 0.8:1 which on the basis of the work of Hill et al. [11], would be expected to demonstrate fluorine loss. The reason why these anomalies only arise for barium rich glasses is thought to be related to the fact that the electronegativity difference between barium and fluorine is less than that between fluorine and potassium. This would result in greater K–F bond strength and thus when insufficient aluminium is present in the glass to fulfil Hill et al.Õs [11] criteria then the presence of excess fluorine in the glass could be strongly associated to potassium with the result that fluorine losses do not occur. Al29 and F19 MASNMR analyses of the potassium rich glasses would be of great benefit in deriving more categoric explanations. If the anomalies referred to above are taken into account, then in general, the substitution of aluminium by boron results in decreases in microhardness and glass transition temperature and increases in coefficients of thermal expansion. The decreases observed for microhardness and glass transition temperature and increases in thermal expansion coefficient are expected to be due to
203
changes in glass structure. If fourfold coordinated aluminium is substituted by fourfold coordinated boron in the glass network the crosslink density values remain the same and therefore, assuming the charge balancing and modification effects are the same, little change would be expected in glass property values. If on the other hand the boron is in threefold coordination as argued in [14] and [15], then two effects arise. These are (a) a reduction in the crosslink density (see Table 3) and (b) release of cations (e.g. 0.5 Mg2þ per AlO4 tetrahedron substituted) from their charge balancing role allowing them to act as modifiers. Thus if boron is in threefold coordination then the glass network becomes more disrupted and the glasses contain greater modifier levels. Both of these structural changes would lead to property disimprovements as the results obtained show. Accordingly, it is possible to conclude that boron is in threefold coordination in the glasses with Q values of 0.5 and 1.0. An interesting effect arises for glasses with Al:F ratios of 1:1. For these glasses (X ¼ 3:0, Q ¼ 0 and X ¼ 2:5, Q ¼ 0:5) the effect of potassium for barium substitution on glass transition temperature is negligible until all barium is replaced by potassium (see Figs. 4 and 9). These effects may be due to the occurrence of a fluorine plus modifier rich phase within an aluminosilicate rich matrix as discussed above. If this is the case, then clearly it only arises when barium is present in the glasses. When barium is fully replaced by potassium then glass transition temperatures fall in a manner which might be expected on the basis of results collected for glasses of lower fluorine content. It is therefore concluded that when barium is present in these glasses then it causes a two phase glass structure to arise and stabilises this glass microstructure.
5. Conclusions For the (BaO, K2 O)–SiO2 –(Al2 O3 , B2 O3 )– (MgO, MgF2 ) glasses studied here the following conclusions can be drawn: 1. Substitution of barium by potassium results in increases in molar volume and coefficients of
204
2.
3.
4.
5.
6.
K. Greene et al. / Journal of Non-Crystalline Solids 325 (2003) 193–205
thermal expansion and decreases in fractional glass compactness, microhardness and glass transition temperature values. These effects are due to the replacement of ONB –Ba–ONB bridges by ONB –K terminations, with attendant decreases in the level of association of the glass network and thus its rigidity. Increasing the fluorine content of the glasses results in increases in molar volume and coefficients of thermal expansion and decreases in microhardness and glass transition temperature values, all of which are consistent with a reduction in glass network connectivity due to the replacement of Al–O–Si by Al–F terminations which reduces the rigidity of the glass structure. For boron free glasses with fluorine substitution levels corresponding to X ¼ 3:0, molar volume and coefficients of thermal expansion have similar values to those for glasses containing lower fluorine levels (X ¼ 2:5). This is attributed to the additional fluorine residing in a modifier rich phase resulting in glasses with similar properties since this phase, which makes up only a small volume of the high fluorine glasses exerts little influence on these properties. For glasses containing the maximum level of boron substitution for aluminium, anomalous property values arise when they are barium rich (X ¼ 0 and 0.25). These anomalies can be attributed to fluorine losses during melting and casting which occur because the aluminium to fluorine ratio is less than 1:1. For glasses where fluorine loss does not occur, the substitution of aluminium by boron results in decreases in molar volume, microhardness and glass transition temperature values and increases in coefficients of thermal expansion. These effects of boron substitution on properties arise because boron is in threefold coordination and modifier cations, not required for charge balancing substituted fourfold coordinated aluminium atoms (or indeed boron), are able to modify the glass network even more. Both of these effects reduce the structural rigidity of the boron containing glasses. For boron free glasses containing aluminium and fluorine in the proportion 1:1, glass transi-
tion temperatures are independent of potassium substitution level (Z) until barium is wholly replaced, where there is a marked reduction. It is thought that this phenomenon arises because of the occurrence of a two phase microstructure where fluorine plus modifier regions within the glass are barium stabilised. Acknowledgements Part of this research was funded by the European Commission under contract no. BE4462. Additional funding for a PhD studentship (K.G.) was provided by Enterprise Ireland through a research scholarship and by the University of Limerick. References [1] D.S. Baik, K.S. No, J.S. Chun, J. Mater. Sci. 30 (1995) 1801. [2] S.N. Hoda, G.H. Beall, in: J.H. Simmons, D.R. Uhlmann, G.H. Beall (Eds.), Nucleation and Crystallization in Glasses, Advances in Ceramics, vol. 4, American Ceramic Society, 1982. [3] T. Uno, T. Kasuga, K. Nakajima, J. Am. Ceram. Soc. 74 (1991) 3139. [4] K.E. Wallace, PhD thesis, University of Limerick, 2000. [5] K.E. Wallace, R.G. Hill, J.T. Pembroke, C.J. Brown, P.V. Hatton, J. Mater. Sci.: Mater. Med. 10 (1999) 697. [6] W. Vogel, Structure and Crystallization of Glasses, Pergamon, 1971. [7] W. Vogel, Chemistry of Glass, American Ceramic Society, 1985. [8] R. Ramesh, E. Nestor, M.J. Pomeroy, S. Hampshire, J. Eur. Ceram. Soc. 17 (1997) 1933. [9] S. Likitvanichkul, W.C. Lacourse, J. Mater. Sci. 30 (1995) 6151. [10] A. Rafferty, A. Clifford, R. Hill, D. Wood, B. Samuneva, M. Dimitrova-Lukacs, Paper submitted to J. Am. Ceram. Soc. [11] R. Hill, D. Wood, M. Thomas, J. Mater. Sci. 34 (1999) 1767. [12] G.H. Beall, Mica glass-ceramics, US patent no. 3,689,923, 1972. [13] D.G. Grossman, J. Am. Ceram. Soc. 55 (1972) 446. [14] J. Zhong, P.J. Bray, J. Non-Cryst. Solids 84 (1986) 17. [15] W.F. Du, K. Kuraoka, T. Akai, T. Yazawa, J. Mater. Sci. 35 (2000) 4865. [16] N.H. Ray, Inorganic Polymers, Academic Press, 1978. [17] N.H. Ray, Br. Polym. J. 7 (1975) 307.
K. Greene et al. / Journal of Non-Crystalline Solids 325 (2003) 193–205 [18] W. Lowenstein, Am. Miner. 39 (1954) 92. [19] Y-M. Chiang, D. Birnie III, W.D. Kingery, Physical Ceramics, John Wiley, New York, 1997. [20] J.E. Shelby, W.C. Lacourse, A.G. Clare, in: Engineered Materials Handbook, Vol. 4: Ceramics and Glasses, ASM International, 1991, p. 845.
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[21] G.M. Singer, M. Tomozawa, Phys. Chem. Glasses 30 (1989) 86. [22] H.M. Mulvihill, PhD thesis, University of Limerick, 2000. [23] E.M. Rabinovich, Phys. Chem. Glasses 24 (2) (1983) 54. [24] T. Maeda, S. Matsuya, M. Ohta, Dent. Mater. J. 17 (1998) 104.
Effect of composition on the properties of glasses in the K2O–BaO–MgO–SiO2–Al2O3–B2O3–MgF2 system K. Greene a
a,b
, M.J. Pomeroy
a,b,*
, S. Hampshire
a,b
, R. Hill
c
Department of Materials Science and Technology, University of Limerick, Limerick, Ireland b Materials and Surface Science Institute, University of Limerick, Limerick, Ireland c Materials Department, Imperial College of Science and Technology, London W2 3SF, UK Received 11 July 2002
Abstract Twenty five glasses in the(1 Z)BaO:ZK2 O:(6 X )MgO:X MgF2 :(3 Q)Al2 O3 :QB2 O3 :8SiO2 system have been prepared and the effect of systematic changes in composition (Z ¼ 0, 0.25, 0.5, 0.75 and 1.0, X ¼ 2, 2.5 and 3.0 and Q ¼ 0, 0.5 and 1) on molar volume (MV), fractional glass compactness (C), microhardness (lHV ), glass transition temperature (Tg mid ) and coefficient of thermal expansion (a) determined. As potassium is substituted for barium (increasing Z), increases in MV and a, and decreases in C, lHV and Tg mid arise which are attributed to the replacement of ONB –Ba–ONB ionic bridges by two ONB –K terminations (ONB ¼ non-bridging oxygen). When fluorine is substituted for oxygen, reductions in lHv and Tg mid and increases in MV and a values occur. These property changes occur because of reduced crosslink densities associated with the replacement of Al–O–Si crosslinks by Al–F terminations. In general, the substitution of aluminium by boron results in decreases in MV, lHV and Tg mid values and increases in a values. These effects are attributed to boron assuming a tri-coordinated state in the glass, giving rise to reduced crosslink densities, together with the release of modifier cations (Mg, K and Ba) from their charge balancing role for four-coordinated aluminiums leaving them free to cause greater network disruption. 2003 Elsevier B.V. All rights reserved.
1. Introduction Glasses based on the system SiO2 –Al2 O3 –B2 O3 – MgO–MgF2 –K2 O typically crystallise to trisilicic alkaline mica (KMg3 AlSi3 O10 F2 ). Such glassceramics are an important class of material as they are machinable to high (±10 lm) tolerances [1]. The microstructure of these materials, comprising highly interlocked K-fluorophlogopite crystals
*
Corresponding author. Tel.: +353-61 202 200. E-mail address: [email protected] (M.J. Pomeroy).
with a Ôhouse of cardsÕ structure embedded in a glass matrix, facilitates microfracture along the weak mica–glass interface and mica basal planes, which prevents macroscopic failure during machining. Mica glass-ceramics containing Ba-fluorophlogopite crystals are attracting interest because they have shown improvements in mechanical properties when compared with more conventional Kfluorophlogopite glass-ceramics [2,3]. However, very little is known about the effect which exchanging barium for potassium has on the properties of the base glasses. Wallace and coworkers
0022-3093/$ - see front matter 2003 Elsevier B.V. All rights reserved. doi:10.1016/S0022-3093(03)00337-5
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[4,5] have shown that replacing an alkaline earth cation such as Ca2þ with an alkali cation such as Naþ in bioactive glasses reduces both the glass transition temperature and the glass density. These reductions are related to decreases in levels of association between adjacent network regions due to the elimination of ionic bridges afforded by divalent cations and non-bridging oxygens (ONB – MII –ONB ) and their replacement by two terminal ONB –MI entities. However, the work of Dietzel [6,7] indicates that the effect of alkali/alkaline earth cation exchange on the properties of glasses can also be related to cationic field strengths. The cationic field strength is a measure of a cationÕs effective force in attracting anions and is given by z=r2 where z is valency and r the ionic radius of the cation. In Ln–Si–Al–O–N glasses (Ln ¼ Y, Ce, Nd, Sm, Eu, Dy, Ho and Er), the variation in glass transition temperatures for glasses of constant composition has been directly related to the cationic field strength of the lanthanide cation being substituted [8]. Thus, it has been reported that glass transition temperatures increase linearly with increasing cation field strengths of the rare earth modifiers. Investigations into the effect of fluorine on the properties of glasses have shown that it is network disruptive [9–11]. The reduction in glass transition temperature of pre-cerammed cannasite glassceramics has been related to the creation of nonbridging fluorines which replace bridging oxygens in the glass network [9]. In apatite–mullite glassceramics [10] and ionomer glass compositions [11], the glass transition temperature decreases with increasing fluorine content as a result of network disruption by fluorine. There is, unfortunately, a major problem with glass systems containing fluorine that is the volatilisation of silicon tetrafluoride (SiF4 ) during glass synthesis which can lead to fluorine deficient glasses and increased production costs. However, Hill et al. [11] have shown that volatilisation can be inhibited by incorporating appropriate amounts of aluminium and basic network modifying oxide to allow aluminium to achieve fourfold coordination. Thus, if the Al:F ratio is greater than 1:1 fluorine loss can be obviated. Boric oxide is typically included in the starting composition of machinable glass ceramics in order
to increase the aspect ratio of the resultant fluorophlogopite crystals [12] which yields a glass-ceramic with improved machinability [13]. However, the effect of boric oxide on the properties and structure of the uncerammed glasses has not been investigated. Research on the structure of boric oxide containing glasses has generally concentrated on the Na2 O–B2 O3 –SiO2 system [14] and to a lesser extent the Na2 O–B2 O3 –Al2 O3 –SiO2 system [15] used in the production of two phase glasses which can subsequently treated to yield porous materials. Even in these relatively simple systems, detailed relationships between the composition, structure and properties of the glasses have not been derived. The structure of aluminosilicate and boroaluminosilicate glasses can be described by applying an inorganic polymer model developed by Ray [16,17]. From studying a typical oxide network such as vitreous silica, Ray postulated that inorganic glasses could be regarded as polymer chains with networks of oxygen atoms linked by mainly covalent bonds through multivalent atoms of network-forming elements. An important factor relating the properties of an inorganic glass to its composition and structure are the density of crosslinks in the structure. The crosslink density can be defined as the average number of ÔbridgingÕ oxygen atoms in excess of two per network former, where bridging oxygens are defined as those that link together two other atoms that are part of the network, e.g. Si, B, Al. If there are not more than two bridging oxygens per network former atom, the structure is equivalent to that of a linear polymer and will therefore have a crosslink density of zero. The crosslink density of pure silica is two which represents a fully crosslinked inorganic polymer. The effect of glass forming, intermediate and modifying ions on the crosslink density can be calculated from the molar glass composition. Given the absence of clear information relating to glasses in the (Ba,K)–Mg–Si–(Al,B)–(O,F) system, the aim of this work was to systematically investigate the effects of substituting potassium for barium, fluorine for oxygen and boron for aluminium on their physical and thermal properties and from the property trends observed, infer possible structural changes arising in the glasses.
K. Greene et al. / Journal of Non-Crystalline Solids 325 (2003) 193–205
2. Experimental 2.1. Glass design In designing this series of novel glass compositions, the following considerations were taken into account: ii(i) The glasses should have enough of each specified compound to ensure that preferential crystallisation to fluorophlogopite would occur. i(ii) The Si:Al ratio was maintained >1:1 in order to adhere to the ÔLowenstein aluminium avoidance principleÕ [18], allowing Al to take up a predominantly fourfold coordination state, thereby favouring glass formation during synthesis. (iii) The Al:F ratio was maintained P 1:1, with enough modifier(s) present to ensure that Al is bonded to fluorine and that local electroneutrality around the Al3þ cation is maintained. This was expected to significantly reduce fluorine loss during glass synthesis [11]. (iv) The need to study, systematically, the effects of the following substitutions on glass properties: (a) K for Ba, (b) F for O and (c) B for Al. The glasses studied were therefore based on the following parent composition: 8SiO2 ð3 QÞAl2 O3 QB2 O3 ð6 X ÞMgO X MgF2 ð1 ZÞBaO ZK2 O; where Z values of 0.0, 0.25, 0.5, 0.75 and 1.0, X values of 2.0, 2.5 and 3.0, and Q values of 0.0, 0.5 and 1.0 were investigated. The minimum X value was limited to 2.0 as this corresponded to a glass having a stoichiometry similar to that of fluorophlogopite whilst the maximum X value was limited to 3.0 as it was expected that the amount of silicon tetrafluoride loss during synthesis would become significant since for X > 3:0 the Al:F ratio would be less than 1:1 [11].
195
oxide (B2 O3 ), magnesium carbonate (MgCO3 ), magnesium fluoride (MgF2 ), potassium carbonate (K2 CO3 ) and barium carbonate (BaCO3 ) powders, thoroughly mixed by ball milling. The purity and supplier of each reagent is shown in Table 1. After mixing, the powders were placed in sintered mullite crucibles (Zedmark Refractories, Earlsheaton, Dewsbury, UK) and fired in an electric furnace (Carbolite HTC 1600, Sheffield, UK) at temperatures between 1400 and 1450 C for 1.0– 1.5 h depending on the composition. Batch weights prior to firing were approximately 450 g. The glass melts were subsequently shock quenched into water to produce glass frit. This frit was dried in a fan oven at 100 C for 24 h, ground in a vibratory mill (Gyro-Mill, Glen Creston Ltd., UK) and sieved to produce three particle size fractions (i) <45 lm, (ii) >45 < 105 lm (coarse) and (iii) >105 lm (frit). Cast glass was produced by remelting frit particles in alumina crucibles (VZS, Fife, Scotland, UK) at 1450 C for 1.5 h in an electric furnace (Carbolite RHF 1500, Sheffield, UK). The molten glass produced was poured into preheated graphite crucibles, annealed for 1 h at a temperature 50 C below the midpoint of the glass transition region and allowed to furnace cool to room temperature overnight. The glass rods produced were 12 mm in diameter and 40 mm in length.
Table 1 Materials and suppliers Component
Percentage of purity
Supplier
SiO2
99.9
Al2 O3
99.86
MgCO3
99.85
MgF2
99.9
B2 O3
99.8
2.2. Glass preparation
BaCO3
99.9
The glasses were synthesised using batches of reagent grade silica (SiO2 ), alumina (Al2 O3 ), boric
K2 CO3
99.9
Tilcon Industrial Minerals, UK GPR Grade, BDH Laboratory Supplies, UK Alfa, Johnson and Matthey, Germany Alfa, Johnson and Matthey, Germany Aldrich Chemical Company Inc, USA Alfa, Johnson and Matthey, Germany E. Merck, Darmstadt, Germany
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2.3. Glass characterisation X-ray diffraction (XRD) data for each of the glasses was recorded using a Philips X-ray diffractometer (Philips XÕPert, Eindhoven, NL) using ), an anode current Cu Ka radiation (k ¼ 1:5406 A of 30 mA and a voltage of 40 kV. Diffraction patterns were collected between 5 and 70 2h. XRD was carried out on both <45 lm size fraction particles and also cast specimens to ensure that the glasses were completely amorphous. The densities of pore free cast glass rods were measured using a water-density balance. Samples, 12 mm in diameter and 10 mm in length, were weighed in both air and distilled water using a Sartorius balance. The densities were calculated using the following equation: q ¼ qwðT Þ Ws =ðWs Ww Þ; where q is the sample density (g cm3 ), qwðT Þ the density of water at the measured temperature (g cm3 ), Ws the sample weight in air (g) and Ww the sample weight in water (g). The accuracy of this technique is ±0.01 g cm3 . Molar volumes were calculated from the density data using the formula P xi M i MV ¼ ; q where xi and Mi represent the atomic fractions and atomic weights of K, Ba, Mg, Si, Al, B, O and F respectively and q is the measured glass density (g cm3 ). Fractional glass compactness (C) values were calculated using the expression C ¼ Vion
total =MV;
where Vion total ¼ xi Vi and Vi is the ionic volume calculated using ionic radii data published in [19]. Microhardness tests were carried out on polished glass samples using a Leco microhardness tester (Akashi Corporation, Japan). A load of 300 g was applied to the sample for 15 s. The resulting indentation diagonals were measured and the hardness was calculated using the following equation: lHV ¼ 1:854P =ðd 2 Þ; where lHV is the Vickers microhardness in GPa, P is the load (N) and d is the mean diagonal of the microhardness indentation (m).
Differential thermal analysis (DTA; Stanton Redcroft STA 1640, Rheometric Scientific, Epsom, UK) was used to determine the midpoint of the glass transition temperature range (Tg mid ). DTA experiments were carried out using a heating rate of 10 C/min in a flowing nitrogen atmosphere with matching pairs of platinum–rhodium alloy crucibles and 50 mg alumina as the reference material. Measurements to determine coefficients of thermal expansion (a) were carried out on cast glasses using a Netsch DIL 402 dilatometer. Samples 10 mm · 4 mm · 4 mm were heated at a rate of 5 C/min in air and a was calculated in the range 200–500 C using the following equation: a ¼ ðDl=lo Þ=DT ; where lo is the original length, Dl the change in length of the specimen and DT the temperature change.
3. Results 3.1. General observations Weight change measurements following glass synthesis showed weight losses consistent with the loss of carbon dioxide from the magnesium, potassium and barium carbonates used in glass preparation. Because of this, it is assumed that all fluorine added to the glass formulation is present in the glass after firing. Furthermore, no weight losses were recorded after the casting of glass bars for property measurements, indicating that no fluorine losses occurred. XRD analysis of all glasses, as melted and fritted and as cast, showed that all materials were amorphous. 3.2. Boron free glasses (Q ¼ 0) Table 2 gives the relevant glass compositions and results for the property measurements for this series of glasses where Z and X were varied. It is seen from Table 2 that as barium is replaced by potassium (increasing Z) then glass densities decrease from 2.90 to 2.63 g cm3 for each fluorine content. Whilst these density decreases may be associated with the lower atomic weight of potas-
K. Greene et al. / Journal of Non-Crystalline Solids 325 (2003) 193–205
197
Table 2 Glass compositions, property values and calculated crosslink density (CLD) values for boron free glasses Glass code GðX ; Q; ZÞ
Density (g cm3 )
Molar volume (cm3 mol1 )
Microhardness (GPa)
Tg mid (C)
CTE (106 K1 )
G(2.0, G(2.0, G(2.0, G(2.0, G(2.0,
0.0, 0.0, 0.0, 0.0, 0.0,
0.00) 0.25) 0.50) 0.75) 1.00)
2.91 2.84 2.76 2.71 2.64
7.66 7.72 7.81 7.82 7.89
6.59 6.53 6.39 6.38 6.11
679 671 665 660 653
6.00 5.98 6.24 6.54 6.65
1.14 1.14 1.14 1.14 1.14
G(2.5, G(2.5, G(2.5, G(2.5, G(2.5,
0.0, 0.0, 0.0, 0.0, 0.0,
0.00) 0.25) 0.50) 0.75) 1.00)
2.90 2.84 2.77 2.69 2.63
7.68 7.72 7.78 7.88 7.93
6.54 6.43 6.28 6.19 6.14
653 650 645 642 635
6.31 6.34 6.48 6.56 7.00
1.07 1.07 1.07 1.07 1.07
G(3.0, G(3.0, G(3.0, G(3.0, G(3.0,
0.0, 0.0, 0.0, 0.0, 0.0,
0.00) 0.25) 0.50) 0.75) 1.00)
2.90 2.83 2.77 2.69 2.63
7.68 7.75 7.78 7.88 7.93
6.44 6.30 6.20 6.13 5.95
638 638 639 636 620
6.25 6.45 6.50 6.60 7.12
1 1 1 1 1
8.00 7.95
3
-1
molar volume (cm .mol )
sium, which is significantly lower than that of barium (39.1 compared with 137.3), increases in glass volume also arise. Fig. 1 shows the variation in molar volume with increasing levels of potassium substitution for barium and indicates that as barium ions are replaced by two potassium ions, molar volume increases in a reasonably linear manner. Fig. 2 clearly shows that the compactness of the glass decreases with increasing potassium for barium substitution indicating that the glass structure expands due to changes in the strength of modifier – glass structural unit interaction. The effect of increased fluorine substitution for oxygen has virtually no effect on density (see Table 2). Effects of this increased substitution on molar volume can be interpolated from Fig. 1 which shows that molar volume changes with increased X value are almost within experimental error (±0.03 cm3 mol1 ). That said, there appears to be an increase in molar volume associated with the initial increase in fluorine content (X ¼ 2:0–2.5) but values for the X ¼ 2:5 and 3.0 substitution levels are almost identical. Microhardness values decrease linearly with increasing substitution of potassium for barium substitution levels (see Fig. 3). In addition, it can be seen from both Table 2 and Fig. 3 that as fluorine substitution levels increase microhardness values decrease.
CLD
7.90 7.85 7.80 7.75 7.70 7.65 7.60 0
0.25
0.5
0.75
1
Z value Fig. 1. Effect of potassium substitution for barium (Z) on molar volume of boron free (Q ¼ 0) glasses with different fluorine substitution levels ( (X ¼ 2:0), M (X ¼ 2:5), } (X ¼ 3:0)).
Glass transition range midpoint temperatures (Tg mid ) data for frit samples are given in Table 2. Fig. 4 shows that whilst they are lowered by potassium substitution, the extents of these decreases are dependant on fluorine substitution level. Thus,
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690
glass transition temperature (°C)
fractional glass compactness
0.585
0.580
0.575
0.570
680 670 660 650 640 630 620 610
0.565 0
0.25
0.5
0.75
1
0
0.25
Fig. 2. Effect of potassium substitution for barium (Z) on fractional glass compactness of boron free (Q ¼ 0) glasses with different fluorine substitution levels ( (X ¼ 2:0), M (X ¼ 2:5), } (X ¼ 3:0)).
6.9
microhardness (GPa)
6.7 6.5 6.3 6.1 5.9 5.7 0.25
0.5
0.75
1
Z value
Z value
0
0.5
0.75
1
Z value Fig. 3. Effect of potassium substitution for barium (Z) on microhardness of boron free (Q ¼ 0) glasses with different fluorine substitution levels ( (X ¼ 2:0), M (X ¼ 2:5), } (X ¼ 3:0)).
for X ¼ 2:0 and X ¼ 2:5 glasses, Tg mid values decrease linearly with increasing potassium for barium substitution (Z), but for X ¼ 3:0 glasses glass
Fig. 4. Effect of potassium substitution for barium (Z) on glass transition temperature of boron free (Q ¼ 0) glasses with different fluorine substitution levels ( (X ¼ 2:0), M (X ¼ 2:5), } (X ¼ 3:0)).
transition temperatures are almost constant up to 75% substitution of potassium for barium. The effect of increasing fluorine substitution levels is to cause significant decreases in Tg mid as X is increased from 2.0 to 2.5. However, for further fluorine substitution (X ¼ 2:5–3:0) decreases are less significant, indicating that the degradation of this property, induced by fluorine, is diminished, except when all barium is substituted by potassium. Fig. 5 shows the effect of Z value on coefficient of thermal expansion (CTE) and indicates that as potassium replaces barium then CTE values also increase. Furthermore, it can also be seen from Fig. 5 that CTE values for the same Z value are only affected by initial increases in fluorine substitution levels (X ¼ 2:5). Further increases (X ¼ 3:0) have little effect. 3.3. Glasses (X ¼ 2:5)
with
constant
fluorine
content
Table 3 shows the compositional variations used to investigate the effects of variations in potassium for barium and boron for aluminium substitution on glass properties, together with
K. Greene et al. / Journal of Non-Crystalline Solids 325 (2003) 193–205
7.1
-6
-1
CTE (10 K )
7.6
6.6
6.1
5.6 0
0.25
0.5
0.75
1
Z value Fig. 5. Effect of potassium substitution for barium (Z) on coefficient of thermal expansion of boron free (Q ¼ 0) glasses with different fluorine substitution levels ( (X ¼ 2:0), M (X ¼ 2:5), } (X ¼ 3:0)).
property values and crosslink density values for the glasses assuming boron is in either three or fourfold coordination. It is seen from Table 3 that for each boron substitution level (Q ¼ 0, 0.5 and 1.0), densities of glasses decrease as levels of po-
199
tassium for barium substitution increase. The extents of the increases decrease with increasing boron substitution for aluminium. Furthermore, it is seen that for constant Z values the effect of boron substitution for aluminium has a minor effect on density with density decreases of the order of 0.02–0.03 g cm3 accompanying the substitution of one mole of alumina by one mole of boric oxide. Fig. 6 shows the effect of Z value on molar volume for each boron substitution level. It is seen that, for all boron levels, molar volumes increase with increasing potassium for barium substitution. The effect of increasing boron substitution on molar volume is not significant as differences are within the calculation error of ±0.03 cm3 mol1 . Fig. 7 indicates that fractional glass compaction values decrease with increasing potassium substitution levels. Again the effects of aluminium substitution by boron are within the calculation error band (±0.003). Microhardness values for these glasses decrease linearly with increasing potassium substitution level for the Q ¼ 0 and Q ¼ 0:5 compositions (see Fig. 8). However, for the Q ¼ 1:0 substitution level the microhardness values of the baria rich glasses (Z ¼ 0 and Z ¼ 0:75) are similar to each other and to the values for the Q ¼ 0:5 substitution. The
Table 3 Glass compositions, property values and calculated crosslink density values for glasses with fixed (X ¼ 2:5) fluorine content Glass code GðX ; Q; ZÞ
Density (g cm3 )
Molar volume (cm3 mol1 )
Microhardness (GPa)
Tg mid (C)
CTE (106 K1 )
CLD B in fourfold
B in threefold
G(2.5, G(2.5, G(2.5, G(2.5, G(2.5,
0.0, 0.0, 0.0, 0.0, 0.0,
0.00) 0.25) 0.50) 0.75) 1.00)
2.90 2.84 2.77 2.69 2.63
7.68 7.72 7.78 7.88 7.93
6.54 6.43 6.28 6.19 6.14
653 650 645 642 635
6.31 6.34 6.48 6.56 7.00
1.07 1.07 1.07 1.07 1.07
1.07 1.07 1.07 1.07 1.07
G(2.5, G(2.5, G(2.5, G(2.5, G(2.5,
0.5, 0.5, 0.5, 0.5, 0.5,
0.00) 0.25) 0.50) 0.75) 1.00)
2.87 2.80 2.72 2.67 2.61
7.66 7.72 7.82 7.83 7.88
6.23 6.17 6.10 6.06 5.88
627 622 624 624 610
6.53 6.51 6.63 6.91 7.53
1.07 1.07 1.07 1.07 1.07
0.96 0.96 0.96 0.96 0.96
G(2.5, G(2.5, G(2.5, G(2.5, G(2.5,
1.0, 1.0, 1.0, 1.0, 1.0,
0.00) 0.25) 0.50) 0.75) 1.00)
2.85 2.76 2.69 2.64 2.58
7.61 7.73 7.80 7.81 7.86
6.17 6.16 5.94 5.78 5.70
614 611 606 600 581
6.28 6.33 6.68 6.83 7.46
1.07 1.07 1.07 1.07 1.07
0.85 0.85 0.85 0.85 0.85
200
K. Greene et al. / Journal of Non-Crystalline Solids 325 (2003) 193–205
6.9
8.00
6.7
7.90
microhardness (GPa)
3
-1
molar volume (cm .mol )
7.95
7.85 7.80 7.75 7.70
6.3 6.1 5.9
7.65
5.7
7.60 0
0.25
0.5
0.75
1
Fig. 6. Effect of potassium substitution for barium (Z) on molar volume of glasses containing 2.5 mol MgF2 (X ¼ 2:5) for different boron substitution levels ( (Q ¼ 0:0), N (Q ¼ 0:5), r (Q ¼ 1:0)).
0.5850
0.5800
0.5750
0.5700
0.5650 0
0.25
0.5
0
0.25
0.5
0.75
1
Z value
Z value
fractional glass compactness
6.5
0.75
1
Z value Fig. 7. Effect of potassium substitution for barium (Z) on fractional glass compactness of glasses containing 2.5 mol MgF2 (X ¼ 2:5) for different boron substitution levels ( (Q ¼ 0:0), N (Q ¼ 0:5), r (Q ¼ 1:0)).
effect of increasing boron substitution level up to an initial level of Q ¼ 0:5 is to decrease microh-
Fig. 8. Effect of potassium substitution for barium (Z) on microhardness of glasses containing 2.5 mol MgF2 (X ¼ 2:5) for different boron substitution levels ( (Q ¼ 0:0), N (Q ¼ 0:5), r (Q ¼ 1:0)).
ardness by 0.2–0.3 GPa. For the Q ¼ 1 substitution level, similar effects arise for the glasses with Z ¼ 0:5, 0.75 and 1.0 but not for the Z ¼ 0 and 0.25 compositions where hardness values are unaffected. The effect of potassium for barium substitution on glass transition temperature is shown in Fig. 9. It is seen that for each boron substitution level there is a decrease in Tg mid , although the effect of Z value on Tg mid for glasses with boron substitution levels corresponding to Q ¼ 0:5 is negligible until complete substitution of barium by potassium (Z ¼ 1:0) is reached. In addition, it can be seen from Fig. 9 that, as boron substitution for aluminium increases, Tg mid values decrease quite markedly. For glasses with Z ¼ 0:5, 0.75 and 1.0, the decreases in Tg mid with Q are linear. However, for the other Z value glasses the decreases in Tg mid are less as boron is substituted for aluminium. Coefficients of thermal expansion (see Table 3 and Fig. 10) increase with increasing levels of potassium substitution, for each boron content. As the level of boron substitution is increased from Q ¼ 0 to 0.5, CTE values also increase. For further increases in boron substitution, the effect on CTE
K. Greene et al. / Journal of Non-Crystalline Solids 325 (2003) 193–205
pansion are similar to those containing no boron and for lower Z (high Ba) values the values of CTE are similar to those for glasses with a boron substitution level corresponding to Q ¼ 0:5.
glass transition temperature (°C)
660 650 640 630
4. Discussion of results
620 610 600 590 580 570 0
0.25
0.5
0.75
1
Z value Fig. 9. Effect of potassium substitution for barium (Z) on glass transition temperature of glasses containing 2.5 mol MgF2 (X ¼ 2:5) for different boron substitution levels ( (Q ¼ 0:0), N (Q ¼ 0:5), r (Q ¼ 1:0)).
7.6
7.1
-6
-1
C TE (10 K )
201
6.6
6.1
5.6 0
0.25
0.5
0.75
1
Z value Fig. 10. Effect of potassium substitution for barium (Z) on coefficient of thermal expansion of glasses containing 2.5 mol MgF2 (X ¼ 2:5) for different boron substitution levels ( (Q ¼ 0:0), N (Q ¼ 0:5), r (Q ¼ 1:0)).
depends upon Z value. Thus, for the high Z value (high K) glasses with one third of the aluminium substituted by boron coefficients of thermal ex-
In general, for each of the boron free glass series in which potassium for barium substitution was investigated, density, fractional glass compactness, microhardness and glass transition temperature values decreased, whilst molar volume and coefficient of thermal expansion values increased. Each of these property trends indicates two structural effects. One is that barium and potassium ions act as network modifiers rather than charge balancing aluminium in the network. The other is that the substitution of potassium for barium results in a decrease in the strength of association between modifier ions and their surrounding elements of the glass network. It is of course non-bridging oxygens (ONB ) which are most strongly coordinated to modifier cations. As indicated in the introduction, the effects of alkali/alkaline earth modifier cation substitutions on glass structure, and therefore properties, can either be viewed from the point of view of a sequential replacement of ONB –MII –ONB bridges with two ONB –MI terminations [4,5] or changes in cation field strength [6–8]. If changes in cation field strength for the substitution of Ba2þ by 2Kþ , are examined (using ionic radii values given by Chiang et al. [19]) then very little change is observed, since the cation field strength of Ba2þ is 110 nm2 and that of 2Kþ is 105 nm2 . Because of this, it appears that the sequential replacement of ONB –Ba–ONB bridges with two ONB –K terminations is responsible for the degradation in properties observed. Such replacements would certainly result in a decrease in the degree of association of the glass network which would facilitate decreases in fractional glass compactness and hardness as the rigidity with which adjacent parts of the glass network could be associated would be substantially decreased. A similar argument would be expected to apply to decreases in glass transition temperatures since segments of the glass would be expected to be able
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K. Greene et al. / Journal of Non-Crystalline Solids 325 (2003) 193–205
to move more readily as the degree of network association decreases due to ONB –K termination effects. With respect to glass transition temperatures, there is a clear anomaly in the effect of Z for the glasses containing the maximum level of fluorine substitution (X ¼ 3). This is discussed later. Increasing thermal expansion values may also be attributable to decreased association levels. However, the situation is somewhat more complex since this property is also related to non-symmetrical vibrations of ions and structural entities, the exact analysis of which is outside the scope of this work. For the glasses containing boron with fluorine contents corresponding to X ¼ 2:5, the general effect of potassium substitution for barium is virtually the same as that for the boron free glasses. Accordingly, the same explanations can be invoked as for boron free glasses. Anomalous property trends (e.g. Tg mid for Q ¼ 0:5 glasses) are addressed below in the paragraphs dealing with the effects of boron content. The effects of increasing fluorine substitution for oxygen on glass properties are similar to those for the substitution of barium by potassium, i.e. all property values disimprove. However, the reasons for the property degradation induced by increased levels of fluorine substitutions are different as fluorine resides in the glass network. As oxygen is replaced by fluorine, potential Al–O–Si crosslinks in the network are replaced by Al–F terminations thus reducing crosslink densities as indicated in Table 2. The property disimprovements observed are therefore consistent with decreased crosslink density trends as has been previously reported for other fluorine containing glass systems [9–11,20–22]. It is clear from the results presented, that, with the exception of hardness, the property values for glasses containing 2.5 and 3.0 mol of MgF2 are very similar for the same potassium content and that glass transition temperatures for the glasses with X ¼ 3:0 are independent of potassium/barium content until all barium is replaced. Given that no weight loss attributable to fluorine was observed for these glasses and the Al:F ¼ 1:1 rule of Hill et al. [11] was adhered to, the level of fluorine substitution can be safely assumed to be that present in the starting glass composition. Thus,
fluorine loss is not the reason for the phenomena observed. Two other possible explanations arise. One involves separation of the glass into a fluorine and modifier (K, Ba and Mg) rich phase and an aluminosilicate rich phase as proposed by Hill et al. [11], Rabinovich [23] and Maeda et al. [24], whilst the other may be related to levels of glass network disruption reaching a maximum at the fluorine content corresponding to X ¼ 2:5. With respect to the possibility of a two phase glass, additional work has shown that the optimum nucleation temperatures of the glasses investigated here are close to their glass transition temperatures indicating that they nucleate via a phase separation process and thus that chemical inhomogeneity may exist in the glasses prior to their nucleation. If this is the case, then the overall glass microstructure would comprise a modifier plus fluorine rich droplet phase in a continuous modifier and fluorine depleted aluminosilicate phase. This might well explain the distinct similarities in molar volume and thermal expansion coefficient values for glasses with X ¼ 2:5 and 3.0, since if much of the excess fluorine (i.e. that corresponding to 0.5 mol MgF2 ) is accommodated in the droplet phase then the matrix phase which controls these particular properties of the glass would be very similar to the bulk glass corresponding to X ¼ 2:5. Such a glass structure would also be consistent with the effects of increased fluorine levels on microhardness. This is because the microhardness values would represent those obtained from a droplet phase–matrix phase composite and as the modifier plus fluorine rich droplet phase would be comparatively soft the overall microhardness values for X ¼ 3:0 glasses would be less than for the X ¼ 2:5 glasses. However, it must be noted that no evidence of modifier plus fluorine crystalline phases were observed during X-ray diffraction and thus the modifier plus fluorine rich droplet phase must be either amorphous or present in very small levels (<2–3%). The alternative explanation is that as the fluorine level is increased from X ¼ 2:0 to 2.5, the extent of network disruption reaches a maximum. This would result in subsequent fluorine for oxygen substitution (X ¼ 3:0) having no further effect on properties such as molar volume and thermal expansion coefficient. Such an argument is how-
K. Greene et al. / Journal of Non-Crystalline Solids 325 (2003) 193–205
ever at variance with the crosslink density data given in Table 2 and therefore it strongly appears that the similarity in molar volume and coefficient of thermal expansion values between glasses with fluorine contents corresponding to X ¼ 2:5 and 3.0 is due to the location of the extra fluorine in a modifier rich second phase. For the barium rich boron containing glasses with low Z values (0 and 0.25), increased levels of boron substitution (Q ¼ 0:5–1:0) have no effect on microhardness and decreases in Tg mid are less than for the potassium rich higher Z value glasses. In addition, the coefficients of thermal expansion for the Q ¼ 0 and Q ¼ 1 substitution levels are similar. For all other Z values, the effect of increasing Q value on glass properties is degradative. The most plausible explanation for the anomalous effects of boron for aluminium substitution is that during the melting and subsequent bar casting of these low Z-value glasses, fluorine loss occurred with the result that the property changes due to boron substitution were negated by fluorine loss. This reasoning would certainly be consistent with the weight losses observed for these glasses and the fact that the Al:F ratio for the Q ¼ 1:0 substitution level is 0.8:1 which on the basis of the work of Hill et al. [11], would be expected to demonstrate fluorine loss. The reason why these anomalies only arise for barium rich glasses is thought to be related to the fact that the electronegativity difference between barium and fluorine is less than that between fluorine and potassium. This would result in greater K–F bond strength and thus when insufficient aluminium is present in the glass to fulfil Hill et al.Õs [11] criteria then the presence of excess fluorine in the glass could be strongly associated to potassium with the result that fluorine losses do not occur. Al29 and F19 MASNMR analyses of the potassium rich glasses would be of great benefit in deriving more categoric explanations. If the anomalies referred to above are taken into account, then in general, the substitution of aluminium by boron results in decreases in microhardness and glass transition temperature and increases in coefficients of thermal expansion. The decreases observed for microhardness and glass transition temperature and increases in thermal expansion coefficient are expected to be due to
203
changes in glass structure. If fourfold coordinated aluminium is substituted by fourfold coordinated boron in the glass network the crosslink density values remain the same and therefore, assuming the charge balancing and modification effects are the same, little change would be expected in glass property values. If on the other hand the boron is in threefold coordination as argued in [14] and [15], then two effects arise. These are (a) a reduction in the crosslink density (see Table 3) and (b) release of cations (e.g. 0.5 Mg2þ per AlO4 tetrahedron substituted) from their charge balancing role allowing them to act as modifiers. Thus if boron is in threefold coordination then the glass network becomes more disrupted and the glasses contain greater modifier levels. Both of these structural changes would lead to property disimprovements as the results obtained show. Accordingly, it is possible to conclude that boron is in threefold coordination in the glasses with Q values of 0.5 and 1.0. An interesting effect arises for glasses with Al:F ratios of 1:1. For these glasses (X ¼ 3:0, Q ¼ 0 and X ¼ 2:5, Q ¼ 0:5) the effect of potassium for barium substitution on glass transition temperature is negligible until all barium is replaced by potassium (see Figs. 4 and 9). These effects may be due to the occurrence of a fluorine plus modifier rich phase within an aluminosilicate rich matrix as discussed above. If this is the case, then clearly it only arises when barium is present in the glasses. When barium is fully replaced by potassium then glass transition temperatures fall in a manner which might be expected on the basis of results collected for glasses of lower fluorine content. It is therefore concluded that when barium is present in these glasses then it causes a two phase glass structure to arise and stabilises this glass microstructure.
5. Conclusions For the (BaO, K2 O)–SiO2 –(Al2 O3 , B2 O3 )– (MgO, MgF2 ) glasses studied here the following conclusions can be drawn: 1. Substitution of barium by potassium results in increases in molar volume and coefficients of
204
2.
3.
4.
5.
6.
K. Greene et al. / Journal of Non-Crystalline Solids 325 (2003) 193–205
thermal expansion and decreases in fractional glass compactness, microhardness and glass transition temperature values. These effects are due to the replacement of ONB –Ba–ONB bridges by ONB –K terminations, with attendant decreases in the level of association of the glass network and thus its rigidity. Increasing the fluorine content of the glasses results in increases in molar volume and coefficients of thermal expansion and decreases in microhardness and glass transition temperature values, all of which are consistent with a reduction in glass network connectivity due to the replacement of Al–O–Si by Al–F terminations which reduces the rigidity of the glass structure. For boron free glasses with fluorine substitution levels corresponding to X ¼ 3:0, molar volume and coefficients of thermal expansion have similar values to those for glasses containing lower fluorine levels (X ¼ 2:5). This is attributed to the additional fluorine residing in a modifier rich phase resulting in glasses with similar properties since this phase, which makes up only a small volume of the high fluorine glasses exerts little influence on these properties. For glasses containing the maximum level of boron substitution for aluminium, anomalous property values arise when they are barium rich (X ¼ 0 and 0.25). These anomalies can be attributed to fluorine losses during melting and casting which occur because the aluminium to fluorine ratio is less than 1:1. For glasses where fluorine loss does not occur, the substitution of aluminium by boron results in decreases in molar volume, microhardness and glass transition temperature values and increases in coefficients of thermal expansion. These effects of boron substitution on properties arise because boron is in threefold coordination and modifier cations, not required for charge balancing substituted fourfold coordinated aluminium atoms (or indeed boron), are able to modify the glass network even more. Both of these effects reduce the structural rigidity of the boron containing glasses. For boron free glasses containing aluminium and fluorine in the proportion 1:1, glass transi-
tion temperatures are independent of potassium substitution level (Z) until barium is wholly replaced, where there is a marked reduction. It is thought that this phenomenon arises because of the occurrence of a two phase microstructure where fluorine plus modifier regions within the glass are barium stabilised. Acknowledgements Part of this research was funded by the European Commission under contract no. BE4462. Additional funding for a PhD studentship (K.G.) was provided by Enterprise Ireland through a research scholarship and by the University of Limerick. References [1] D.S. Baik, K.S. No, J.S. Chun, J. Mater. Sci. 30 (1995) 1801. [2] S.N. Hoda, G.H. Beall, in: J.H. Simmons, D.R. Uhlmann, G.H. Beall (Eds.), Nucleation and Crystallization in Glasses, Advances in Ceramics, vol. 4, American Ceramic Society, 1982. [3] T. Uno, T. Kasuga, K. Nakajima, J. Am. Ceram. Soc. 74 (1991) 3139. [4] K.E. Wallace, PhD thesis, University of Limerick, 2000. [5] K.E. Wallace, R.G. Hill, J.T. Pembroke, C.J. Brown, P.V. Hatton, J. Mater. Sci.: Mater. Med. 10 (1999) 697. [6] W. Vogel, Structure and Crystallization of Glasses, Pergamon, 1971. [7] W. Vogel, Chemistry of Glass, American Ceramic Society, 1985. [8] R. Ramesh, E. Nestor, M.J. Pomeroy, S. Hampshire, J. Eur. Ceram. Soc. 17 (1997) 1933. [9] S. Likitvanichkul, W.C. Lacourse, J. Mater. Sci. 30 (1995) 6151. [10] A. Rafferty, A. Clifford, R. Hill, D. Wood, B. Samuneva, M. Dimitrova-Lukacs, Paper submitted to J. Am. Ceram. Soc. [11] R. Hill, D. Wood, M. Thomas, J. Mater. Sci. 34 (1999) 1767. [12] G.H. Beall, Mica glass-ceramics, US patent no. 3,689,923, 1972. [13] D.G. Grossman, J. Am. Ceram. Soc. 55 (1972) 446. [14] J. Zhong, P.J. Bray, J. Non-Cryst. Solids 84 (1986) 17. [15] W.F. Du, K. Kuraoka, T. Akai, T. Yazawa, J. Mater. Sci. 35 (2000) 4865. [16] N.H. Ray, Inorganic Polymers, Academic Press, 1978. [17] N.H. Ray, Br. Polym. J. 7 (1975) 307.
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[21] G.M. Singer, M. Tomozawa, Phys. Chem. Glasses 30 (1989) 86. [22] H.M. Mulvihill, PhD thesis, University of Limerick, 2000. [23] E.M. Rabinovich, Phys. Chem. Glasses 24 (2) (1983) 54. [24] T. Maeda, S. Matsuya, M. Ohta, Dent. Mater. J. 17 (1998) 104.