CA2315398A1 - Dense refractories with improved thermal shock resistance - Google Patents

Abstract

The present invention discloses a dense refractory material, which includes a spinel matrix and a micro-crack initiating phase dispersed in the matrix. The micro-crack initiating material introduces micro-cracks into the refractory material, which inhibits catastrophic failure as a result of the effects of thermal shock. A method of manufacturing a dense refractory material which includes the steps of mixing precursor oxides for a spinel material, calcining the material, forming the spinel material into a green form of the product and firing the green form to produce the final form is also disclosed.

Description

- WO 99/32417 . PCT/AU98/01049 DFNSE REFRACTORIES WITH IMPROVED THERMAL SHOCK RESISTANCE
The present invention relates to a refractory material and to a method of manufacturing the refractory. material.
A simple definition of a refractory material is one which resists the effects of high temperatures. Commonly, the term refractory material is applied to relatively low cost products that are used in many industrial processes, typically operating at high temperatures, to contain corrosive materials, such as molten metal and slaps. As such refractories are an important class of materials.
The following factors are relevant to the design of refractory materials:
chemical compatibility;
thermal shocks constraints on start-up;
operating conditions;
slag penetration;
hot strength creep resistance; and cost.
Many ceramics materials have properties in common with refractory materials. gor example, ceramic materials are characterised by excellent chemical stability, high hardness and a brittle nature. In comparison with refractory materials, typically ceramic materials have poor thermal shock resistance. The combination of poor thermal shock resistance and high cost limits the use of ceramic 'materials in refractory applications.
There are two options to minimise the effects of thermal shock. The first is to avoid the initiation of cracks and WO 99/32417 PCT/AU9$/01049 - a -the second is to avoid catastrophic crack propagation.
Thermal Shock Damage Resistance Parameters (R), a measure of a material s resistance to the above types of failure, were proposed by Hasaelman (see Introduction to Ceramics, Ringery a~° edition 1976 pp 825-30). The physical properties required to compute Thermal Shock Damage Resistance Parameters are thermal conductivity k, thermal expansion coefficient a. YouaQ~s Modulus E, effective fracture energy y,==, and strength (MOR) a. Specifically, the Thermal Shock Damage Resistance Parameters R' and R
can be expressed as:
kQ
R' _ '-' ( 1 ) Ea and Ey ""= Q~' (a) where R' is the parameter for the resistance to crack initiation and R " " is the parameter for the resistance to crack propagation.
The material characteristics for inhibiting crack formation are high strength with respect to elastic modulus. The requirements for minimising the extent of crack propagation are a high product of work of fracture and elastic modulus with respect to strength. Thus, the design requirements for a material for inhibiting crack formation and crack propagation are different.
=t is known that resistance to catastrophic failure, which is required in refractory applications, can be improved by the introduction of enough cracks of sufficiently large size so that crack propagation takes place semi-statically.

- ~ WO 99/32417 . PCT/AU98/01049 =t is also known that, alternatively, resistance to catastrophic failure can be achieved by the introduction of microstructural inhomoQenieties in any form Which serve as stress concentrators in the material. In this way, cracks will form locally, but catastrophic failure.is avoided as a result of the small average stress in the material.
Conventional refractory materials are designed for chemical stability, thermal shock resistance, and cost. This is achieved through a con~romise between reducing the effective surface area for attack and increasing resistance to crack propagation. Typically, a conventional refractory material has an open structure With between 15 and 20%
porosity. The open structure allows rapid penetration of slaQs and gases but inhibits crack propagation. A schematic representation is shown in Fig. 1.
The shortcomings of this compromise approach to design were recognised by the late Ronald C. Garvie. He proposed that a dense thermal shock resistant material would offer superior performance to a conventional refractory material. To achieve this goal he introduced micro-cracks into the microstructure. This increased the work of fracture for the material by promoting crack branching. The end result was a dense material with the chemical stability of an advanced ceramic and the thermal shock resistance of a porous refractory. This micro-crack toughened coawposite material is disclosed in US patents 5,296,420 and 5,334,563 of Garvie. A schematic representation of the composite refractory material is shown in Fig. 2.
The essential features of the composite material disclosed in US Patent 5,334,563 are that the material have less than 12% porosity and comprise:
a matrix of alumiaa, with 5 to 90% by volume of the alumina grains having a diameter in the range of 15 to 8O microns;
particles of monoclinic zirconia dispersed is the matrix, each dispersed particle comprising an agglomerate of microcrystals which;
(a) are strongly bonded together;
(b) exhibit a strong thermal expansion anisotropy; and (c) a size such that cracks do not form spontaneously within the agglomerates after cooling from high temperatures in the range of 1600°C; and the alumiaa and the monoclinic zirconia being chemically inert with respect to each other within the temperatures used in practice.
The Garvie US Patent also discloses a number of other combinations, such as: mullite as the matrix and zirconia as the dispersed material; silicon nitride as the matrix and boron nitride as the dispersed material; barium titanate as the matrix and zirconia as the dispersed material; silicon carbide as the matrix and boron nitride as the dispersed material; alumiaa as the matrix and aluminium titanate as the dispersed material; spinal as the matrix and zirconia as the dispersed material; and fosterite as the matrix and zirconia as the dispersed material.
The basis of the Garvie US Patent is the addition of a dispersed second phase in a continuous dense matrix with very particular inter-dependence of the respective thermal expansion coefficients of the phases. Specifically, the use of specific grades of monoclinic zirconia as the 8isperaed phase produced as enhanced dilatatioaal/coatractional mismatch is a number of matrices such as alumina or zircon. Aa optimised composition, with respect to thermal shock damage resistance (measured by retained strength) was determined empirically by Garvie to be 8% by weight of zirconia in alumina and 10% by weight is zircon.
Extensive chemical attack of ceramic matrix materials, such as zircon and alumina, limits the use of such ceramic composites in many corrosive industrial applications. These include applications where the ceramic composite is in contact with slaps used in iron and steel making operations.
A further significant problem is the prohibitive cost of production of the composite ceramic materials oa an industrially realistic scale.
It is known that reaction sintering of zircon mixtures can result is the formation of oxide zirconia dispersions (see for example OS Patent 2842447 by Schlotzhauer and wood and Cambier; Baudin de La Lastra, Dilate and Leriche Brit.
Ceram. Soc. Trans. and J. 83 pp 196-200, 1984). As discussed by Cambier et. al. the use of this technique is useful in the manufacture of zirconia is a mullite or alumina mullite matrix. These materials are characterised by high strength with MOR values that can reach 400 to 5001~Pa. =n addition, these materials are typically characterised by pores around the zirconia particles as a consequence of the process. That is, the original zircon particles lose silica to the surrounding matrix. There is a volume decrease reported to be about 20% for the zircon particles converted into zirconia particles. This results in the formation of pores associated with the zirconia grains. Furthermore, Schlotzhauer and Mood (col. 3, lines 11-20 of the US Patent) indicate that the high corrosion resistance of the final products is a consequence of the lack of cracking associated with the inversion of the zircoaia as it is heated or cooled through 1000~C. The presence of such pores would accommodate the volume expansion of zirconia associated with the inversion of the zirconia on cooling without the generation of stresses or strains.
US Patent 4298385 of Claussen and Steeb discloses a method for producing bodies having high fracture toughness and "substantially equal" mechanical strength. This is achieved by the addition of from 4 to 25 volume % zirconia grains ("embedment material") with a diameter from 0.3 to 1.25 lun in as anisotropic ceramic matrix, such as alumiaa. The improvement is the properties of the fabricated products resulted, by way of exaa~le in the case of alumiaa with unstabilised zirconia, from the production of extremely fine micro-fissures and a high fissure density in the products. This was reported to significantly increase the toughness, thermal shock resistance and impact strength as compared to products prepared without the zirconia addition. In addition, it was found that it was preferable to disperse the zirconia within agglomerates of zirconia and matrix phase with a size of 2 to 15 um containing from 4 to 25 volume % (preferably 8 to 25 volume %) of the phase (Col 2, lines 9 to 54). For alumina this is ectuivalent to 5.8 to 32.8 wt% and preferably 11.3 to 32.8wt%.
Furthermore, it is also taught that the use of large embedment material is to be avoided, as the strength is considerably reduced (Col 4, lines 43 to 46). From the results presented in Figure 1 of the US patent it is clearly seen that increasing the particle size embedment materials from 0.3~m to 1.25pm required an increase in the amount of the embedmeat material from 10 vol% to 15 vol%
(14 to 20.6 wt%) and this indicates the benefits of the smaller zirconia grain size. The examyles disclose the use of high vol% of the embedment phase. For example in Fig. 6 1,, - - WO 99/32417 . PGTIAU98/01049 _ 7 _ the vol% ranges from 15 to 25 vol%. It is further reported that such materials are especially suited to high tea~erature Qas turbine elements.
US Patent 4,804,644 of Aaseau, Lawson and Slasor also discloses a material which includes dispersion of zirconia is a matrix, in this case as 0'-sialon matrix. 0'-sialon is a solid solution based on silicon oxynitride (Si,N,o) where there is substitution of Al sad 0 for Si and N
respectively. The OS Patent discloses a number of methods for the preparation of such materials. However, for materials produced according to the methods the zircoaia is is the tetragonal form. It is stated that improvements in properties would result from the transformation of meta-stable tetragonal zirconia to the monoclinic form in response to a tensile stress typically caused by an advancing crack tip. The transformation results is the formation of compressive stresses that tend to close the cracks. =adeed from example 18 of the US Patent. the zircoaia is reported to be is the tetragonal form at room ten~erature. For the zircoaia to be effective the size of the particles must remain small to prevent spontaneous transformation on cooling. There is ao report of the physical properties such as strength aa8 thermal shock resistance of such bodies formed.
An object of the present invention is to provide a refractory material with enhanced corrosion, erosion sad thermal shock resistance which alleviates the disadvantages of the kaov~ra refractory materials discussed above.
Accordir~ to one aspect of the present invention there is provided a dense refractory material which includes a matrix and a micro-crack initiating single czystal phase formed from fused zircoaia dispersed in the matrix.
The tezm "dense" is understood herein to mean that the refractory material has limited open porosity, typically less than 5% by volume.
According to another aspect of the present invention there is provided a dense refractory material which includes a spinal matrix.
The spinal group of materials is understood herein to mean materials that are described by the general formula:
ABsO, where A'' is typically is either singly or in combination Mg, Fe, Zn and Mn and B'' is typically either singly or in combination Al, Fe, Cr and Ma.
Examples of spinals are magnesium aluminium oxide MQA1~0~, magnetite Fe,O~, and chromite FeCr,O~ . An exaa4ple of a "mixed" spinal is MQ(Al,Fe)s0~.
The spinal group of materials have a cubic crystal structure and, therefore, are isotropic. As a consequence, the spinal microstructure is relatively stress-free.
Furthermore. the spinal group of materials is relatively stable at high temperatures and while maintained at temperature.
The spinal may include one or more additional elements. The additional elements may include Li, Mg, Ca, Ti, 1~, Fe, Co, Ni, Cu, Zn, sr and sa, for divalent cations and A1. Cr, Fe, and Mn as the trivalent cations. In addition, spinal phases can exist over a range of coa~ositions with respect to the ratio of the divalent to trivalent cations.
The selection of the additional elements may depend on a Wide range of factors. By way of example, one factor is the environment in which the refractory material will be used. Specifically, in situations where the refractory material will be in contact with molten slag in metal smelting operations, the additional elements may be S selected to optimise the chemical stability~of the refractory materials with respect to the slag. By way of further example, another factor is to include additional elements to assist in the manufacture of the refractory material as a dense refractory material.
A further advantage of spinals is that they can exist over a range of coagposition without a change in phase. For example, magnesium aluminate spinal can be magnesium rich, stoichiometric (Mg to Al ratio of 1:2) or aluminium rich.
This allows the loss of an element from the crystal lattice without decomposition to form a new phase or compound.
Typically, the formation of new phases can result in physical disruption of the refractory body or the formation of less refractory phases. The ability of the spinal to adapt to the environment without a change is phase enhances the stability of the products.
=t is known that spinals, such as magnesium aluminium oxide MgAl~O~ and chromite FeCr=O~ spinals, have excellent corrosion resistance to slaps in metal smelting operations.
However, typically, the spinals are coarse and are use8 as grits or aggregate in refractory bodies for many metal making and cement making operations and not as the matrix of a dense refractory material. Moreover, the refractories that incorporate these spinals are in the form of traditional refractoriea that are characterised by open porosity and are not dense refractory materials.
Furthermore, whilst the C3arvie US patents propose the use of spinals in a matrix of a micro-crack toughened refractory material, the disclosure is speculative sad not supported by examples.

It is preferred that the refractory material further comprises a micro-crack iaitiatin~ phase dispersed in the matrix.
=t is preferred that the micro-crack initiating phase be no more than 15% by volume of the material.
It is preferred particularly that the micro-crack initiatiz~gr phase be no more than 10% by volume of the material.
It is preferred that the spinal matrix be at least 80% by volume of the material.
=t is preferred particularly that the spinal matrix be at least 90% by volume of the material.
=t is preferred that the micro-crack initiating phase comprises a dispersion of single crystals.
It is preferred that the micro-crack iaitiatiaQ phase be formed from zirconia.
It is preferred that the zirconia have a particle size in the rsaQe of 5 to 50E,tm.
It is preferred that the zirconia have a particle size in the range of 10 to 20~Im.
=t is preferred particularly that the zirconia be fused zirconia.
The_micro-crack iaitiatinQ phase may be formed from any other suitable material, such as boron nitride and silicon carbide.
=t is preferred that the spinal be manufactured from law cost precursors.

According to another aspect of the present invention there is provided a dense refractory material which includes a spinal matrix and a micro-crack initiating phase dispersed in the matrix.
According to another aspect of the present invention there is also provided a method of manufacturing a dense refractory material product which includes the steps of:
(i) mixing precursor oxides for a spinal material;
(ii) calcining the mixture to form the spinal material;
(iii)forming the spinal material into a green form of the product; arid (iv) firing the green form of the product to produce the final form of the product.
It is preferred that the method further includes the step of mixing the spinal material produced in step (ii) with an additive, such as zirconia. selected to form a micro-crack initiating phase dispersed in the fired product.
According to the present invention, the spinal material is formed by reaction of the precursor oxides. This is typically carried out in the temperature range of 800°C to 1600°C and preferably in the range of 1000°C to 1400°C for dwell times at temperature ranging up to at least 10 hours. Longer times are generally preferred for lower calcination temperatures and shorter times for temperatures is the upper reaches of the range. D~rell times of 1 hour or less are possible for higher temperatures in the range.

- WO 99132417 _ PCT/AU98I01049 Typically, the spinal material formed is they milled (if necessary) to form a finely divided powder suitable for deasification in the secondary heat treatment of step (iv).
Typically, the average particle size should be less than ZO
E,Im, preferably less than 5~"tm and more preferably less than 2~,I,m.
Typically, the additive which form8 the dispersed phase is then added to the spinal powder.
Typically, the spinal ponder and the additive are then moulded or formed into the desired shape is a green form in step (iii). This can be done with and without the use of additives to increase the plasticity of the powder facilitating forming into the desired "green" shapes.
The green shape is then heated to effect densification is the firing step (iv). This is typically carried out in the temperature range of 1000°C to 1800°C and preferably in the range of 1400°C to 1600°C for darell times at temperature ranging up to at least 10 hours. Longer times are generally preferred for lower secondary heating tea~eratures and shorter times for temperatures is the upper reaches of the rsage. Dwell times of 1 hour or less are possible for higher tem~peratur~s is the range. Temperatures can also be reduced by use of sintering assists that can be incorporated into the structure of the spinal. However, it is pr~farable that the firing temperature used in the manufacture be at least as high as the expected operating temperature where the refractory is to be used.
Sintering aids may be used to promote densification of the refractory material. These aids can foszn liquids that result in enhanced diffusion rates thereby increasing the densificatioa rate. Where these additives exist as secondary phases in the final microstructure they can exert a deleterious effect on the performance of products. It is well - , WO 99/32417 . PCT/AU98/01049 known that the presence of silica-based glasses and calcium-containiaQ phases can lead a marked decrease in the high temperature properties of alumiaa based refractories.
Appropriate sintering aids may be used to promote deasificatioa at lower temperatures without a loss of performance. The firing cycle of refractory materials can represent a substantial proportion of the cost to manufacture products. Reducing the firing temperature can result is a lower cost to manufacture products. =n addition, it is postulated that improved chemical stability is obtained by using a matrix material that contains the main elemeata of the slag in a solid solution within the crystal structure of the matrix phase or where stable phases are produced as a result of the interaction of elements in the slag with the matrix.
The dense refractory material of the present invention contains micro-cracks in the microstructure after fabrication. These micro-cracks are characterised by emanating from the dispersed phase (typically formed by zircoaia additions) and extending over several grain diameters is the microstructure. Typically the grain size is the order of greater thaw IO yam.
As stated above, the spinal group of materials is defined by the general formula A8s0~ where A'' is typically is either singly or is combination big, Fs, Za and Ma and B3' is typically either singly or is combination Al, ge, Cr and »n. The spinal may include one or more additional elements.
The additional elements may include Li, Mg, Ca, Ti, Ma, P'e, Co, Ni, Cu, Zn, Sr and Ba, for divalent cations and Al, Cr, ge, and Ma as the trivalent catioas. In addition, spinal phases can exist over a range of compositions with respect to the ratio of the divalent to trivalent cations.
As indicated above, zircoaia is the preferred additive to for the dispersed phase. It is well lmowa that certain 8opants can stabilise the high te~oo~erature crystallographic forms of zirconia at room temperature. Typically, dopants include magnesia, ceria, yttria and calcia. For the case of spinals that tyyically include elements that can stabilise the high temperature forms of zircoaia it is surprising that products formed according to the present invention do not contain zirconia particles that are stabilised. Such stabilisation would render the micro-cracking mechanism responsible for the improvement in thermal shock inoperative. From thermodynamic considerations it is believed likely that the dopaats for zirconia would be partitioned between the zircoaia and the matrix. With the exception of certain instances is the finer fractions, this it not observed. Furthermore it is believed that large particle size and hence uareactivity of the zirconia particles is responsible in part or whole for the observed behaviour. This indicates the requirement for large uareactive zirconia particles. Another advantage of such particles is the relative low cost of large uareactive powders and grits as compared to finely divided and reactive powders. Such fine fractions maybe advantageous for other reasons but they do not contribute to the improvement in thermal shock behaviour for refractory bodies as described by the current invention.
The present invention overcomn~es th~a problems of obtaia3ag low cost refractory materials with high erosion and corrosion stability.
Ia relation to the prior art discussed above, the process disclosed by Schlotzhauer and Wood aa8 Cambier is clearly different to that of the present invention. =a the present invention, cracks are deliberately created by the inclusion of the zircoaia (or other dispersed micro-crack initiating phase). That is, the stresses and strains associated with the inclusion of zirconia into the matrix can not be accommodated by the matrix and result in the formation of cracks.
A feature of the present invention is the tolerance to S impurities and the fact that low cost refractory grade precursors can be used. This allows the use of low coat refractory precursors. It is speculated that the finer fractions of the zircoaia materials used are able to react with impurities to produce more refractory phases.
Ia addition, the body disclosed by Claussen sad Steeb is substantively different to that is the present invention.

The Claussen and Steeb body retains high fracture strength and fracture toughaeas. This is achieved by the requirement for the use of a large vol% of micron and preferably sub-microa zircoaia material. The materials produced according to the teachiaga of the present invention are for refractory type applications. A requirement for this type of material is relatively low cost. This typically means below US$5,000 per tonne for the finished product. Such a final price requires the use of inexpensive rsw materials.

Sub-micros zirconia powders are expensive. At the current prices for such zirconia powders at the levels required according to the teachings of Claussen sad Steeb would equal or in some cases exceed the price of the final product produced according to the present invention. In the present invention the amount of zircoaia addition is minimised, thereby allowing bodies to be cost competitive with conventional refractory materials. gor materials of the present invention, the strength of the final bodies is sacrificed for improved thermal shock resistance. Materials of the present invention would not be suitable for applications such as turbine blades. However, bodies as described according to the present invention are eminently i suitable for applications were high chemical resistance and thermal shock resistances are required but where high strength is not require8 such as refractory applications.

- WO 99/32417 _ PCT/AU98101049 It is important to note that is the present invention the presence of <5um fraction of zirconia does not enhance the thermal properties of the refractories. This is demonstrated when ~L S zircoaia particles that contain -2%
particles less than Sum. In the case of the.AFM ZC03 material the parceatage of <5fun zircoaia particles is increased to >20%. However, this does not lead to an increase in performance is terms of thermal shock. =a fact, more ZC03 material is required as compared to the b~L S
zirconia source. This is attributed to the finer fractions being inoperative for the improvement in thermal shock according to the present invention. In addition, it has been observed that the finer fractions of zirconia are preferably attacked in contact with slaps typically encountered in operating conditions. The presence of the finer fractions as taught by Claussea and Steeb are clearly inferior for the applications intended for the current materials for the reasons outlined.
The following examples are used to describe the invention in a non limiting manner.
BXA~~I~ES
Examples 1-9 The object of the Examples 1-9 was to compare the perfoxawace of a micro-crack toughened refractory material in accordance with the present invention which includes a low cost single crystal fused zirconia (AFM Grade 3) dispersed phase with a lmawn micro-crack toughened composite material based on agglomerates of monoclinic zircoaia (~ S) proposed by Garvie.
The raw materials used were as follow8:

WO 99132417 _ PCT/AU98/01049 Raw Materials Component Supplier ..

A1~03 Alcoa (A1000) ZrO~ blEL grade S or AFM C3rade ZC03 The batch size was a nominal 2008. The batches containing MEL S were designated Examples 1-5 and 6-9 for the AFM
containing range. The starting coa4positions are given in the following table.
Starting Composition in parts Example Batch MEL S ZrOs AFM ZrO~ A1s03 5 CMAll 10 - 90 8 CMA15 - 8 9a The alumiaa poarder was combined with zirconia in the proportions given using the milliag conditions as outlined in the following table. The objective for ail batches was to thoroughly distribute the ZrO, rather than reduce the particle size. See the following table for details.

Ball Milling Conditions Time 0.5 hours Powder 0.2 kg Balls 0.5 kg YTZP

Fluid 0.3 1 iso propanol Binder 2.0/2.Og Surcol/alycerol The resultant slurries were dried at 80°C. Segregation of the constituents was avoided by ensuring that slurry viscosity remained high and by uae of a shallow drying pan.
ears were pressed from the dried powder with a geometry suitable for strength testing (MOR), Young's Modulus and work of fracture testing (T~TOF). These were formed by die pressing at a pressure of 551~Pa with a bar die of dimensions 5 x 5lmm for MOR bars and 7.5 x 102mm for wOF bars. The bars were then bagged and cold isostatically pressed to a pressure of 210Mpa. The samples were fired in air on an alumina setter plate. The firing cycle used is given in the following table:
Firing Cycle in Air Heat 80C.h'1 Dwell 1600C for 2 h Cool 100C. h'1 to room temperature All bars were diamond ground to the dimensions required for compliance to the ASTM E399-83 test for work of fracture testing (wOF). This was a nominal 5 x 10 x 85mm. Hare for strength testing (biOR) were also machined in compliance with the ASTM Standard C1161-94 for flexural testing.
Dimensions for these bare were a nominal 3 x 4 x 45mm.
For the determination of Young's Modulus, the ends of the WOF bars were ground square. Young's Modulus was determined WO 99/32417 - PCTlAU98/01049 by a Transient Vibration Method (ASTM Standard C1259-94) of a right prismatic beam in the flexural mode. Densities were determined by the direct measurement method. The results are presented in the following table.
Mechanical properties Example ZrO~ Young~s Modulus FHD
(wt%) (GPs) (g.cai') 1 0 375 3.90 2 4 275 3.92 3 6 165 3.93 4 8 142 3.95 5 10 134 3.96 6 4 363 3.94 7 6 244 3.95 8 8 169 3.96 9 10 158 3.98 The results indicate an increase in fired bulk density with increasing zirconia contest. This is consistent with zirconia having a significantly higher theoretical density as compared to alumina. Young~s modulus was then plotted as a function of zirconia content.
prom the results in Fig. 3, it can be seen that:
(i) there is a decrease in YounQ~s Modulus with increasing zirconia content for both types of zirconia; and (ii) the MEL S material was more effective in forming micro-cracks at lower levels of addition as compared to the AF~t ZC03 material.
The particle size of the two zirconias was determined (see following table).

- - WO 99/32417 _ PCT/AU98/01049 particle Size of Zirconia Powders Material dio (~) dsa (gym) d9o (~) %<5~

MEL S 10.1 19.1 35.1 ~ 2.1 AFM ZC03 2.4 14.6 36.1 23.7 lot 011709) From the results of the particle size analysis it can be seen that the average size of the MEL S is slightly larger.
However, the greatest difference is is the shape of the particle size distribution. The MEL S material has a sharper distribution. This is clearly be seen by the comparing the d10 values where the AFM material has a much higher concentration of fixes. =t is anticipated that isolated zirconia grains is the microstructure with a particle size less thaw 5~im will contribute little to micro-crack formation. For the AFM material this is almost ~25% of the material.
The data from Fig. 3 was replotted considering only zirconia particles above.SN,m in size. This plot supports the hypothesis that the sub 5N,m grains contribute little to micro-crack toughening.
From the results, it can be sees that the smaller fractions of zircoaia contribute little to micro-crack toughening.
Furthermore, the results indicate that for the particle sizes of zircoaia used, a critical zirconia density is reQuired to initiate micro-cracks.
Examination of etched polished surfaces of the Examples containing AFM ZrOs revealed that this phase was composed of either single crystals or particles with only two or three grains. This indicates that it is possible to use a low cost micro-cracking agent.

- ~ WO 99132417 - PCTIAU98/01049 Examples 10-11 The objective of Facamples 10/11 was to investigate the performance of a micro-crack toughened refractory material S having a spinal matrix in accordance with the present invention produced by the method of the present invention from relatively low cost raw materials.
A cup test was used to evaluate the material is contact with both metal and slag.
The raw materials used to manufacture the cups are given in the following table.
Raw btaterials Component Supplier MgC03 Causmag A1s03 Alcoa (1tA13) ges03 Aldrich 131,005,1 purity 99%+) ZrO~ mEL Grade 8 Causmag is a refractory grade precursor. This material was crushed is a ring mill to produce a powder with a particle size less than 75E.im. 1CA13 alumiaa is also a refractory grade precursor. The price of lCAl3 alumiaa is roughly as order of magnitude less than A1000 alumina as use8 is Examples 1 to 9. The starting coac4positioas are given is the following table.
Starting Composition in Parts 8xample Ng0 ., . A1s03 10a 27.8 66.7 lia 27.2 61.7 Note: Mg source was added as MgC03 The alumina and maQnesite were ball milled (see table).
Ball MilliaQ Conditions Time 16 hours Powder 3 k~

Balls 9 kQ MQ PSZ

Fluid 4.5 ldistilled water Binder No binder was added at this stage The water was removed by pan drying at 80°C. The dried cake of maQnesite and alumiaa were calcined at 1400°C to decompose any carbonates or hydroxides present (see table) and to form a spinal.
Calciaation to Produce Spinal Heat 200C.h'1 Dwell 1400C for 1 hour Cool 200C. h'1 or natural rate Iron oxide sintering aid was added to the pre-reacted spinal. The overall co~ositions are Qivea in the following table:
Overall Compositions is Parts Example Batch Spinal Fe,03 Overall Composition lOb SP805 94.5 (SPOT) 5.5 MQ(A11.9Feo,=)O~

llb SPB06 89.2 (SP02) 10.8 MQ(All_eFeo.s)Os The batches were milled using the following conditions:

Ball Milling Conditions Time 4 hours Powder 3 kg Halls 9 kQ MQ PSZ

Fluid 4.5 l distilled water Binder 1 wt% PVA, 1 wt% glycerol and lwt% Dispex The zircoaia was added to the slurry after milling and immediately prior to spray drying. The zirconia addition to the slurry was at a level of 4 volume % (6.3 weight %). The slurry was continuously stirred prior to spray drying to minimise the effects of settling.
The overall starting compositions are given in the following table:
Overall Compositions in Barts 8xample Batch Spinet wt% ZrOa wt%

lOc SHZOl 93.7 (Mg(All.sFeo.i)O~)6.3 llc SBZ02 93,.7 (Mg(All.eFeo.a)04)6.3 wet bag cold isostatic pressing (CIP) techniques were used to fabricate the cups using tooling consisting of a polyurethane bag and steel mandrel. Lids were fabricated by die pressing followed by bagging and CIP. A pressure of 210Mpa was used for all runs. The cong~oaents were fired using the following firing cycle.

Summary of Firing Cycle in Air Heat 100C .h'1 Dwell 130C far 30 minutes Heat 100C .h-1 Dwell 750C for 60 minutes Heat 100C. h''' Dwell 900C for 60 minutes Heat 100C. h'1 Dwell 1700C 240 minutes for Cool 100C. h'luntil over. the natural cooling rate takes The nominal dimensions of the cups are given is the following table.
Nominal Dimensions of Cups Dimension Nominal Measurement Outside Diameter 55 mm 1.

=nside Diameter 30 mm Overall Height S5 mm Depth of Bore 35 mm Mass 500 g An iron making slag was used for the teat. The lime to silica ratio was in the range of 1 to 1.4 and the F'e0 content in the range of 0.5 to l0%. The pig iron was machined to be a snug fit into the cups. The slag was loaded into the crucibles snd compacted to provide a dense powder bed to protect the metal from oxidation during heat up before melting of the slag. The details of the combined slag-metal-cup test used are given in the following table.

- as -Cup Test Details Slag Addition a0 g Metal Pig Iron nominal 4% carbon Metal Addition 40 g (close fit in bore slag on top) Tea~erature 1700~C

Dwell at Temperature 4 hours Atmosphere Air (static) Cover Loose fitting lid The performance of these materials was excellent with very low dimeasioaal change observed after the test. The cups were essentially single phase materials with the iron incorporated into the crystal structure of the spinal.
There was little slag penetration into the crucibles. Both metal and slag were detected after the test.
8xamples Ia and 13 The objective of $xamples 12/13 was to investigate the effect on performance of variations in composition of the spinal matrix of micro-crack toughened refractory materials in accordance with the present invention.
The raw materials used were as follows:
Raw Materials Coa~onent Supplier MgC03 Causmag Al~Oj AlCOa (KA13 ) ZrOs Magnesium Electron M8L S

The Causmag was crushed in a ring mill to produce an agglomerated powder less than 75E,i,m in size. The starting compositions are given in the following table. The final WO 99132417 _ PCT/AU98101049 composition after the total process is given under the NCOmment8" COlumn.
Starting Composition in Parts Example Composition MgCO, A1,0, Comments 12a SP95 46.7 53.3 Magnesia Rich 70 wt%

Alumina 13a SP96 40.6 59.4 Alumiaa Rich 75 wt%

Alumina The alumiaa and magnesite were mixed in a ball mill (sae table).
Ball Milling Conditions Time 16 hours Powder 0.7 kg Balls 3 kg Mg PSZ

Fluid 1 6distilled water The water was removed by pan drying. The dried cake of magnesite and alumina were calciaed at 1400°C to decompose any carbonates or hydroxides pr~sent and to form a spinal using the conditions as disclosed in Examples 10 and 11.
The spinal was crushed to produce a powder with a dso less than 5E,tm. The milling conditions used were as follows:

WO 99/32417 _ PCT/AU98/01049 _ 27 _ Hall l~illinQ Conditions Batch 12b (SP95) 13b (SP96) Details Time 4 4 hours Powder 0.6 0.6 k9 Balls 3 3 kQ Mg PSZ

Fluid 2 1 ldistilled watsr Hinder Yes Yes l.5wt% PYA, l.2wt%
glycerol and 1wt% Dispsx The zirconia was added to the slurry after milling sad immediately prior to spray drying. The zirconia addition to the slurry was at a level of 4 volume %. The overall starting compositions are given in the following table:
Starting Compositions in parts Example Composition Spiael ZrO~

12c SPZ95 93.7 6.3 13c SPZ96 93.6 6.4 The slurry was spray dried. During spray drying, the slurry was continuously stirred to minimise the effects of settling. Wet bag cold isostatic pressing (C=P) techaiQues were used to fabricate the cups an8 lids as described in Examples 10 and 11 from the dried powder. The firing cycle used to densify the test cups sad lids was the same as described for Examples 10 and 11.
After the cup test of Example 12c, examination of a cross section taken from the crucible revealed extensive damage.
There was swelling of the cup as the result of the formation of internal porosity in the walls of the cup. The ~~tearing~~ of the microstructure is indicative of the formation of the porosity occurring at high temperatures.
The performance was assessed as poor.

- , WO 99/32417 - PCT/AU98/01049 - a8 -Hy stark comparison, after the cup test of 13c, examiaatioa of a cross sectioa taken from the crucible showed little evidence of slag attack. There was evidence of slag penetration into the body without any sign of major disruption. The performance was assessed as good.
The results reveal a dramatic effect of stoichiometry on the chemical performance. The aluminium rich materials were superior to the magnesium rich spinals for the slag tested.
Example 14 The objective of the example was to investigate the thermal shock resistance of a micro-crack toughened refractory material having a dispersed single crystal phase in a spinal matrix in accordance with the present invention.
The raw materials used were as follows:
Raw Materials Component Supplier MQC03 Causmag (Milled) A1s03 Alcoa (1CA13 ) ZrO~ AP'M ZC03 The CausmaQ was used as supplied. The particle size of the as-received powder was leas than 75E.tm in size. The starting composition is given in the following table. The composition after calcination is also given in brackets.
Composition of Example 14 (Parts) Composition MgC03 (MQO) A1s03 SP101 (Starting) 41.5 58.5 SP101 (Est.Fiaal) (25) (75) WO 99/32417 . PCT/AU98/01049 The alumiaa and maQaesite were mixed is a ball mill (see table).
Ball Milling Conditions Time 16 hours Powder 5 kQ I

Balls 15 kQ MQ PSZ

Fluid 7.5 1 distilled water The water was removed by pan drying. The dried cake of maQnesite and alumina was calciaed at 1400~C to decompose any carbonates or hydroxides present and to form a spinal.
The firing cycle used Was the same as disclosed in Examples 10 and 11. The powder was milled using the following conditions:
Hall Milling Conditions Time 4 hours Powder 3 kQ

Balls 9 kQ MQ PsZ

Fluid 3 tdistilled water Binder 1 wt% PVA, 1 wt% glycerol and lwt% Dispex The zircoaia was added to the slurry after milling and immediately prior to spray drying. The zirconia addition to the slurry was at a level of 8 weight%. The slurry was continuously stirred prior to spray drying to minimise the effects of settling.
~Tet bag cold isostatic pressing (CIP) techniques were used to fabricate the bars for thermal shock testing from the dried powder. A pressure of 210 MPa was used for all runs.
The components were fired using the firing cycle as outlined for $xamples 10 and 11.

- - WO 99132417 _ PCTIAU98/01049 After firing sad machining, the bars were subjected to a rapid heating using a gas torch passing across the surface of the samples. The flame Was hot enough to cause localised melting oa the surface. Although major cracks were formed, the saa~le was in one piece at the conclusion of the test.
By comparison, under this test, an alumina bar fabricated according to Example 1, resulted is the formation of shards roughly a 1 cm is size. This shows the excellent thermal shock resistance of the bodies fabricated from low cost precursors according to the invention.
Examples 15-17 The objective of Examples 15-17 was to investigate the chemical stability of refractory materials having a spinal matrix in accordance with the present invention.
The precursors used are listed is the following table.
Raw Materials Component Supplier MgC03 A j ax ( Lab f3rade ) A1a03 Alcoa (A16S(i) CaCO, Ajax (Uailab) SiOa 5 micron Mia-u-Sil The calcium sad magnesium carbonates were calcined at 900°C
to decompose the carbonate and hydroxides present before use. The starting compositions are given in the following table:

- - WO 99/32417 _ PCT/AU98/01049 Starting Compositions is parts 8xample Composition MQO A1~0, SiOs Ca0 15 SP04 28.3 71.7 - -16 SPG03 28.I 71.0 1.0 ~ -17 SPD05 28.1 70.9 - 1.0 The calcined magaesite, alumina and silica precursors were mixed in a ball mill (see table).
Ball Milling Conditions Time 16 hours Powder 0.1 kQ

Milling Media 0.5 kQ YTZP

Milling Fluid 0.2 l iso propanol cinder None The powder was pan dried at 80~C to remove the fluid. Bars of nominal fired dimensions 20 mm long and with a square cross section of 5 mm ware fabricated by uniaxial pressing the powder is a steal die followed by cold isoatatic pressing using wet bag techniques at a pressure of 210 MPa.
Samples were densified using the firing cycle as outlined in Exaa~les 10 and 11. The fired bulk densities obtained after firing are given is the following table.

WO 99I324I7 _ PCT/AU98/01049 Fired Bulk Density After girinQ
Example Batch FiriaQ Time FBD AP (%) Temp(~c) (h) (sr.cm') 15 SP04(05)C89 1700 4 3.44 0.5 16 SPG03(02) 1700 4 3.40 0.4 17 SPD05(Ol)C107 1700 4 3.42 0.1 Notes FBD Fired bulk density AP Apparent porosity To determine the chemical stability, the materials were heated in contact with a slag. The test consisted of placing a sample in a crucible and surrounding with pre-mixed slag. The crucible was removed and the sas~ple extracted from the slag at temperature. Details of the test are summarised in the following table.
8t Crucible Test Crucible 15 ml Pt 5%Au Mass of Slag 6 Q

Mass of Sample 4 Q

Test Temperature i550~C

Duration 2 hours The slag was the same as used is Examples 10 and 11. After the slag test, the degree of slag penetration increased in the order of 8xaa4ple 17 > Example 16 > Example 15.
8xample 18 The objective of the example Was to investigate the firing temperature required to produce a refractory material having a spinal matrix in accordance with the present iaveatioa.
The precursors used are listed in the folloaring table.
Raw Materials Compoaeat Supplier MQC03 Causmag A1s03 AlCOa (A13) The maQaesium carbonate was calciaed at 900°C to decompose the carbonate aad hydroxides preseat before use. The atartiag compositions are givea is the following table:
StartiaQ Composition in 8arts Example Compositioa Mg0 A1s03 18 spgao a8.a ~1.~

A vibro milliag techaique was used for mixing aad particle size reductioa of the alumina. (sae table).
Vibro Milliag Coaditioas Time 2 hours Pawder 0.1 kg Milliag Media 0.8 kQ YTZP

Milling Fluid 0.1 t iso propaaol Biader None The powder was pan dried at 80°C to remove the fluid. Discs with a aomiaal greea diameter of 25 mm and mass of 10 g were fabricated. Saa4ples were produced usiag uaiaxial pressiag is steel dies followed by cold isostatic pressing - - WO 99132417 PCT/AU9.8I01049 using wet bag techniques at a pressure of 210 i~Pa. The firing cycle as described in Examples 10 and 11 was used for the densification of the samples with the exception of the maximum temperature and dwell times.
The fired bulk densities obtained after firing for selected temperatures sad times are given in the following table.
Effect of FiriaQ Temperature oa Densification for NA Spinets Example Batch FiriaQ Time FBD %

Temp (h) (Q.~,) Theoretical (C) Density i8a spF20(02)D 1600 1 3. i8 89.1 il 18b SPF20(03)C 1650 1 3.35 93.8 18c SpF20(04)C 1700 1 3.39 95.0 18d SPF20(05)C 1750 1 3.41 95.5 18e SPF20(01)D 1700 4 3.42 95.8 Notes FBD Fired bulk density The results indicate that the temperatures in excess of 1700°C are required to produce high density spinets products.
Example 19 The objective of the example was to investigate the density of micro-crack toughened refractory materials having a dispersed single crystal phase in a spinal matrix is accordance with the present invention.

The raw materials us~d were as follows:
Raw Materials Con~poneat Supplier MgC03 Causmag (Milled) A1s03 AlcOa ( ICA13 ) ZrO~ AFM ZC03 The Causmag was used as supplied. The particle size of the as-received powder Was less than 75~un in size. The starting co~osition is given in the following table. The composition after calcinatioa is also given in brackets.
Composition of Example 19 (parts) Composition MgCO, (Mg0) A1s03 SPF106 (Starting) 41.5 58.5 SPF106(8st. Final) (25) (75) The alumina and magaesite were mixed is a ball mill (see table).
Ball Milling Conditions Time 16 hours powder 5 kg Balls 15 kg Mg PSZ

Fluid 7.5 l distilled water The water was removed by pan drying. The dried cake of magnesite and alumina was calcined at 1400°C to decompose nay carbonates or hydroxides present and to form a spinal.
The calciastioa cycle used was the same as used in Examples 10 and 11. After calcination, the powder was milled using - ~ WO 99/3241? PC'T/AU98101049 the following conditions:
Ball Milling Conditions Time 16 hours Powder 3 kg Balls 9 kg b~Q PSZ

gluid 4.5 distilled water Hinder 1 wt% PVA, 1 wt% glycerol and 1wt% Dispex The zirconia was added to the slurry after milling and immediately prior to spray drying. The zirconis addition to the slurry was at a level of 5.1 volume % (8.0 weight%).
The slurry was continuously stirred prior to spray drying to minimise the effects of settling.
Wet bag cold isostatic pressing (C=8) techniques were used to fabricate samples for densification studies. A pressure of Z10 MPa was used for all C1P runs. The components were fired using the following firing cycle except for 19d which used the firing cycle as disclosed for Examples 10 and 11.
Summary of Firing Cycle in Air Heat 100C.h-~

Dtaell Max temp for 50 minutes Cool 100C. h'1 until the natural cooling rate takes over.

The densities obtained after firing at selected temperatures is given in the following table.

$ffect of FiriaQ Temperature on Deasification 8xample Batch Firing Time FBD AP
Tea4p (~C) (h) (Q.cni (%) 3) 19a SPZ106(04)C185 1500 I 2.96 16.5 19b spz106(02)Fr~R 1600 1 3.40 5.5 19c SPZ106(O1)C183 1700 4 3.50 2.0 Examples 20 and Z1 The objective of Examples 20/21 Was to investigate the effect of variations in composition of the spinet matrix and firing temperature on the density of micro-crack toughened refractory materials in accordance raith the present invention.
The raai materials used are Qivsn in the following table.

Raw Materials Component Supplier MQC03 CausmaQ

A1~03 Alcoa ( iCAl3 ) Fes03 Aldrich ( 31, 005,1 ) ZrO~ M8L Orade S

The CausmaQ Was crushed is a ring mill to produce a porader with a particle size less than 75Etm. The final composition after the total process is given under the Comments column.
Starting Composition is Parts Example ~ MQO ~ A1~03 ~ Fe~O, ~ Comments 20s 27.8 66.7 5.5 MQ(A11.9Feo.1)Oe 21a 27 .2 61. 10 . MQ(All.egeo.z 7 8 ) Oe Note: MQ source was added as MQCO, The maQaesite, alumina and ferric oxide were ball mille8 (see table).
Ball Milling Conditions Time 16 hours Powder 3 kQ

Balls 9 kQ MQ PSZ

gluid 4.5 1 distilled water The water was removed, by pan drying at 80°C. The dried cake of maQaesite, alumina and ferric oxide were calcined at 1400°C to decompose any carbonates or hydroxides present and form a spinal as described in 8xamples 10 and 11. The batches were milled using the following conditions:
Ball Milling Conditions Time 16 hours Powder 3 kQ

Balls 9 kQ MQ BSZ

gluid 4.5 distilled water Hinder 1 wt% PVA, 1 wt% glycerol and lwt% Dispex The zirconia was added to the slurry after milling and immediately prior to spray drying. The zirconia addition to the slurry was at a level of 4 volume % (6.3 weight%). The slurry was continuously stirred prior to spray drying to minimise the effects of settling. The overall starting compositions are given in the following table:

overall Compositions in t~leiQht percent%
Example Batch Spinal wt% ZrOs wt%

20b SBZ03 93.7 (MQ(A11,9Feo.1)O~) 6.3 21b SBZ04 93.7 (MQ(All,sFeo.s)O~) 6.3 S Dfscs were fabricated by die pressing followed by bagging and wet bag cold isostatic pressing at pressure of 210 MPa.
The samples Were fired using the following firing cycle except for 20f an8 21f which used the firing cycle as disclosed for Examples 10 and 11.
Summary of giring Cycle in Air Heat 100~C.h'1 Dwell Max temp for 60 minutes Cool 100~C.h'i until the natural cooling rate takes over.

The densities obtained after firing at selected temperatures are given in the following table.

- ~ WO 99/32417 - PCTlAU98I01049 8ffect of Firing Temperature on Densificatioa 8xample Batch Firing Time FBD AP
Ted (C) (h) (Q.cm'') (%) 20c SBZ03(04)C185 1500 1 3.48 2.5 20d SBZ03(03)C184 1550 1 3.54 1.8 20e SBZ03(02)FLR 1600 1 3.56 2.0 20f SBZ03(O1)C183 1700 4 3.55 1.8 21c SBZ04(04)C185 1500 1 3.62 1.9 2Id SBZ04(03)CI84 1550 1 3.60 2.0 21e SBZ04(02)FhR 1600 1 3.64 1.6 21f SBZ04(O1)C185 1700 4 3.55 2.8 A coa~arison of the fired bulk deasities after firing for examples 18 to 21 is given in the following table.
8ffect of Firing Temperature oa Bulk Density SQ.caa') $xample Coa~ositioa Tet~ /Time 1500/1 C h 1700/4 18 Stoichiometric 3.18 3.39 19 A1,0, Rich 2.96 3.39 3.50 20 Fe ContalainQ 3.48 3.54 3.56 3.55 21 Fe Containing 3.62 3.60 3.64 3.55 IO

From these results, it can be seen that there is a siQaificant advantage in the use of alumina rich spinals over stoichiometric compositions. ~n additioa, the addition of iron results in a siQaificaat decrease in the sintering i temperatures.
E~cam~les 22-40 The objective of 8xamples 22 to 40 was to investigate the beneficial effect on densification of selected sintering aids.for the spinal matrix of micro-crack toughened refractory materials in accordance with the present invention.
The raw materials used were as follows:
Raw Materials Sintering Aid Component Code Supplier - MgC03 Causmag - A1~03 AlCOa (~CA13 ) - ZrOs AFM ZC03 Ca0 CaC03 BDH AaalaR

TiOs TiOs HDH

MnOs MnO~ 8DH

Ni0 N30 BDH

Cu0 Cu0 HDH

Cr~O, Cr~03 ALDRICH

Fe~03 Fe~03 BDH

COO COO BDH

Za0 Zn0 AnalaR

Sr0 SrC03 BDH

The Causmag was crushed in a ring mill to produce an agglomerated powder less than 75E,~m is size. The starting compositions are given in the following table.

- WO 99/32417 _ PCT/AU98/01049 - 4a -Starting Composition of the Spinal in parts Composition MgCO, Also, Comments SPFla6 44.7 55.3 Stoichiometric spinal The alumiaa and magaesite were mixed in a ball mill (see table).
Ball Milling Conditions Time 36 hours Powder 5 kg Balls 15 kg Mg PSZ

Fluid 7.5 ldistilled water The water was remove8 by pan drying. The dried cake of magnesite and alumiaa were calcined at 1400~C to decompose any carbonates or hydroxides present and to form a spinal using the conditions as disclosed in 8xamples 10 aa8 11.
The spinal was crushed to produce a powder with a dso less than 5~tan. The milling conditions used were as follows:
Ball Milling Conditions SPFla6 catch Details Time 3ah hours Powder 5 kg Balls 15 kg Mg PSZ

Fluid 5.5ldistilled water Binder 1wt% PVA, lwt% glycerol and lwt% Dispex To examine the effect of the dopant concentration on the sintering behaviour for each dopaat, batches were produced at a level of 5 mol% and 10 mol% with respect to the pre reacted spinal (SPFla6). The following batches were WO 99/32417 _ PCT/AU98/01049 S
produced. The exception was Example 23 that was made reaction sintering of the oxides.
Starting Compositions Exaao~le Batch 5mo1% lOmol%

23 SQEZ02 Cu0 26 SQgZ03 N10 27 SQEZOd TiOs 28 SQgZ04 T10=

29 SQaZ05 ge~03 30 SQHZ05 ge~p3 31 SQaZ06 Crs03 32 SQHZ06 Crs03 33 SQ8Z07 Ca0 34 SQF'Z07 Ca0 35 SQEZOS Co0 36 SQFZ08 Co0 37 SQEZ09 Za0 38 SQFZ09 Zn0 39 SQEZ10 Sr0 40 SQFZ10 Srp The zirconia powder was added dust I5 minutes before the end of milling to prevent any decrease of zirconia particle size. The coacentratioa of zirconia was 8 wt% of the total batch.

_ , WO 99132417 - PCTIAU98/01049 Milling Conditions for Nixing pears Type Ball milling powder see compositions 0.3 kg Media Y-TZP balls 0.9 kg Fluid =sopropaaol 750 ml 8iader No binder Milling time without ZrOs 16 h after adding ZrOs 15 min The slurry including the media was poured from the milling containers into glass containers. The liquid was removed by pan dsyiag in vacuum at 70 °C for 20 h. To remove the milling media from the powder and for a better distribution of the zirconia particles in the powder sieving was applied. The powder was passed through two sieves with a grid size of 4000 Ei,m sad 600 Elm.
Pellets were fabricated from the granulate8 powder batches.
The fired pellets were a nominal 20 mm in diameter with a nominal mass of 10 g for deasificatioa studies and bars nominally 20mm long by 5 mm by 5mm. The samples were uniaxially pressed followed by wet bag cold isostatic pressing at 2i0 Mpa.
The components were fired using the following firing cycle.

- WO 99/32417 . PCT/AU98101049 Summary of girinQ Cycle in Air Heat looc .h'1 Dwell 130C for 34 minutes Heat 100C .h'1 Dwell 750C for 60 minutes Heat 100C .h'1 Dwell 900C for 60 minutes Heat 100C .h'1 Dwell Maximum for minutes Cool 100C. h'= until the natural cooling rate takes over.

The maximum temperatures evaluated for the density studies were 1400°C, 1500°C, 1600°C and 1700°C. Samples for the dip test were sintered at 1700°C.
The results of the densification studies are given in the following tables.
Effect of Dopant on gired Bulk Density ( 5m~ol%) Example/ 22 23 25 27 29 31 Firing Temp SPZ126 SQEZ02 SQEZ03 SQEZ04 SQGZ05 SQGZ06 C
Cuo Nio T30s Fes03 Cr~o3 1400 2.a4 1.89 2.26 3.48 2.45 2.28 1500 2.59 2.35 2.58 3.52 2.93 2.55 1600 3.43 3.34 3.48 3.50 3.51 3.43 1700 3.54 3.30 3.53 3.40 3.51 3.53 WO 99/32417 _ PCT/AU98/01049 Effect of Dopant oa Fired Bulk Density (5mol%) Continued Example/ 33 35 37 39 Firing Temp C SQEZ07 SQEZOS SQEZ09 SQEZ10 Ca0 Co0 Za0 Sro 1400 2.30 2.25 2.25 2.28 1500 2.93 2.63 2.64 2.59 1600 3.40 3.53 3.33 2.62 1700 3.32 3.56 3.58 3.40 Effect of Dopaat oa Fired Bulk Density (lOmol%) Example/ 22 24 26 28 30 FiriaQ Temp C SPZ126 SQFZ02 SQFZ03 SQFZ04 SQHZ05 Cu0 N30 TiOs Fes03 1400 2.24 2.47 2.24 3.36 2.61 1500 2.59 2.93 2.57 3.52 3.12 1600 3.43 3.58 3.53 3.35 3.53 1700 3.54 3.47 3.60 3.26 3.52 Effect of Dopaat oa Fired Bulk Density (lOmol%) Continued Example/ 32 34 36 38 40 Firing Temp C SQHZ06 SQFZ07 SQFZ08 SQFZ09 SQFZ10 Cr~03 Ca0 Coo Zn0 Sr0 1400 2.31 2.34 2.24 2.29 2.18 1500 2.55 3.13 2.65 2.60 2.35 1600 3.40 3.17 3.55 3.29 2.66 1700 3.55 - 3.57 3.59 3.32 From the results it can be clearly seen the beneficial effects of CuO, TiOs , N30, Fes03, Co0 on the densificatioa of MCT spiaels. It is important to note that the addition - - WO 99/32417 _ PCTIAU98I01049 - 4' - _ of Ti and Ca resulted in these phases being detected as discrete second phases by enemy dispersive analysis (EDS) in conjunction with the scanainQ electron microscope (SEM).
Typically the Ti had reacted with the finer zirconia fractions inevitably present to form a new secondary phase.
The performance of the different doped spinals was investigated using a slag dip test. The conditions were the same as 8escribed in Examples 15-17.
nip Test Conditions Cruc ible Pt 5%Au Mass of Slag 6Q

Mass of Sample 4Q

Soak Temperature 1500~C

Soak Time 120 minutes The results after testing are given in the following table.

Observations after the Dip Test 23 26 a8 Cu0 Ni0 T30, Good resistance. Good resistance. Complete Slag penetration Slag penetration infiltration of slag limited to 50-200um limited to 50-200um with slag composition in No ZrO,in this No ZrO, is this zonecentre of the zone. samples the same as original slag with additional Ti detected.

FesO, CrzO, Co0 ~atermediate with Good resistance. Good resistance.

evidence of slag Slag penetration Sla ~ penetration wetting all grain boundaries. l~ited to 50-200um limited to 50-2001un Thickness of the GB No ZrOs in this zoneNo ZrO,in this zone lass than is the case of Ti additives From the results it can be seen for the slag tested that the performance of the Cu, Ni, Cr and Co were good but the performance of the Ti doped samples was inferior for the slag tested. As noted, the use of Ti as a sintering assist resulted fa a second phase as detected by 8DS technique.
However it is postulated that in different slaps the Ti containing phase could~exhibit much greater resistance to slag penetration. This is consistent with the selection of the sintering assist for the intended application.
Example 41 to 45 The object of Exaaoples 41 to 45 Was to illustrate the i effects of addition of zirconia in the present material and to indicate the mechanism.
The raw materials were treated prior to use. The alumiaa powder Was dried at 120°C for approxianately 16 hours. The maQnesite was calcined to remove any carbonates and hydroxides present.
Raw materials Coac~oaeat Supplier MQC03 Ajax (flnilab Lab Reagent) .A1~03 Alcoa (A168(3) ZrOz mEL Grade S

The calciaatioa details are given in the following table.
Summary of Calcinatioa giriaQ Cycle in Air Heat 200C .h-1 .

Dwell 900C for 3 hours Cool 200C . h'1 until the natural cooling rate takes over.

The dried alumiaa powder was combined with the magnesium oxide produced from the calcinatioa process in proportions given in the following table.

- - WO 99I324i7 - PCT/AU98/01049 Startinst Compositions $xample 41 42 43 44 45 MQO (Q) 54.8 53.97 53.1 52.2 51.3 A1~03 (Q) 138.7 136.5 134.3 132.0 129.9 Binder (Q) 2.0/2.0 2.0/2.1 2.0/2.0 2.1/2.0 2.0/2.0 =so propanol 400 400 400 400 400 (ml) Mill media 799.4 801.7 801.1 800.9 799.3 (Q) Mill time 14.5 14.5 15 15 15 (hr) At the end of the milling period the mills were removed from the rack, opened sad the ZrOs added. The mills were then returned to the rack and given an additional 30 minutes of rotation, the objective being to thoroughly distribute the ZrO~ rather than reduce the particle size.
The Quantities of the MEL c3rade S ZrOs added are given in the table.
Zircoaia Additions 8xample M8L g (Q) 41 SPZ11 6.42 42 SBZ12 9.57 43 SpZl3 12.69 44 S8Z14 15.76 45 SPZ15 18.80 The slurries were separated from their respective milling sad the slurries dried in a vacuum oven at 80°C and 200kPa.
Segregation of the constituents was avoided by ensuring that the slurry viscosity remained high and by use of a shallow 8ryiag pan. Drying time was ~24hours for all batches.
For all batches, the dried powder cake was gently crushed by use of a small cumber of YTZP balls in a coarse granulating sieve then through sieves of decreasing size with the final size being 500yim.
The sieved batch weight losses were between 1.0 and 4.0%
and the milling media weight loss was zero.
Four bars were pressed from all batches. These were formed by 8ie pressing with a bar die of dimensions 7.5 x 101.9 mm and a die face pressure of 35MPa (25 bar hydraulic pressure). The bars were then bagged and cold isostatically pressed to 210MPa. The firing cycle used to deasify the samples is given in the following table.
Summary of Firing Cycle in Air Heat 100C .h'1 Dwell 130C for 30 minutes Heat 100C .h'1 Dwell 750C for 60 minutes Heat 100C .h-1 Dwell 900C for 60 minutes Heat 100C .h'1 Dwell 1700C 240 minutes for Cool 100C. h-1 until the natural cooling rate takes over All bars were diamond machined to nominal dimensions of 5.0 x lO.Omm cross section and approximately 80mm length. A
second firing cycle was carried out on all test bars to remove all traces of the hold dower wax used in the machining process.

- 5a -A series of mechanical tests was carried to determine various physical properties. The Modulus of Elasticity is Tension (Young's Modulus, E) was determined by the analysis of a transient flexural vibration of a beam of the test material. The Impulse Excitation of Vibration (ASTM C1259-95) technique used With the arindo-Soaic~. The fracture toughness To, and the fracture energy determinations, Ys (Initiation Fracture Energy) and ~y"~= (Work of Fracture) were determined on the same bars which had been notched to roughly half the depth. These were than tested is 3 point SENB geometry in accord with ASTM E399-83. Strength a, was determined by a four point flexural strength test (MOR) oa the tested SENB bar halves. The thermal shock damage resistance parameter R' and R " " (from Hasselmaa), were calculated.
Q,2 The retained strength after thermal shock Was also determined. Bars cut for the flexural test referred to above, were quenched from 700~C into boiling water. These were then tested for strength and normalised retained strength, aR/a, computed.

WO 99/32417 - PCT/AU9$/01049 $xample 41 42 Property SPZ11 SPZ12 Fired Density, p (gcm') 3.42*0.01 3.47*0.00 Youay~s Mod. 8 (C3Pa) 238.8*2.0 242.2*0.8 Strength, MOR. a (MPa) 179*9.8 162.2*1.0 Fracture Toughness, R~ 1~1m''~' 2.38*0.32 2.58*0.11 Fracture Energy, ~y",o~ ( Jm ' 31. 7 *4 31. 9 *2 ) .1 .1 Thermal Coad. R (yPm'1R'i)15.0 15.0 Coeff. Thermal Exp. a (x10 sK-=)7.6 7.6 R~ 1.48x10' 1.32x10' R~n 2.36x10' 2.94X10' Retained Strength,*aR (Mpa) 11.9*1.0 19.9*4.2 Normalised, ** Q,~/a - 7% 12%

Example 43 44 d5 Property SPZ~13 SPZ14 SPZI5 Fired Density, p (gam') 3.49*0.00 3.50*0.02 3.50*0.00 Young's Mod. 8 (Opa) 239.4*0.1 74.6*1.1 46.4*0.2 Strength, MOR. a (tea) 126.0*6.3 47.811.5 29.9*1.3 Fracture Toughness, MNm'~s 2.58*0.15 1.82*0.04 1.41*0.03 Rig Fracture Energy, (Jm'') 31.313.1 62.5*3.4 60.6*1.3 ~~t Thermal Coad. R (Wbn'''R'1)15.0 15.0 15.0 Coeff . Thermal Exp.(x10'~1C'1)7. 6 7 . 6 7 . 6 a R~ 1.04x10' 1.26x10' 1.27x10' Rpn 4.72x10' 2.04x10-' 3.15x10-' Retained (MPa) 34.0*1.3 34.711.42 20.811.7 Strength,*a*

Normalised,** a,~/a - 27% 73% 70%
f --a...~,.on ,.,~a "~re~uQzu measures ay MvR after a quench from 700~C into mater at 100'C.
** " the retained strength (*) divided by the pristine strength.

- - WO 99I324I7 . PCTIAU98/01049 The flexural strength before and after thermal shock and normalised retained strength after thermal shock testing as a function of zircoaia content are shown in Figs. 4 and 5.
From these results the addition of zirconia as taught in the present invention increases the thermal shock resistance at the expense of the absolute strength of the body.
From the results it can be seen that in stark contrast to Claussen and Steeb (US Patent 4,298,385), products according to the present invention undergo a substantial decrease in the absolute strength and the fracture toughness decreases with increasing zirconia content. This behaviour is the opposite as observed with the Claussen and Steeb materials that showed no significant loss in strength and as increase in fracture toughness. However, the thermal shock behaviour chows that the mechanism that operates in these materials enhances the thermal shock resistance as demonstrated by increased retained atrenQth after the thermal shock test.

Claims (22)

CLAIMS:

1. A dense refractory material which includes a spinal matrix and a micro-crack initiating phase dispersed in the matrix.

2. The material defined in claim 1 wherein the micro-crack initiating phase is no more than 15% by volume of the material.

3. The material defined in claim 2 wherein the micro-crack initiating phase is no more than 10% by volume of the material.

4. The material defined in any one of the preceding claims wherein the spinal matrix is at least 80% by volume of the material.

5. The material defined in claim 4 wherein the spinal matrix is at least 90% by volume of the material.

6. The material defined is any one of the preceding claims wherein the micro-crack initiating phase includes a dispersion of single crystals.

7. The material defined in any one of the preceding claims wherein the micro-crack initiating phase is formed from zirconia.

8. The material defined is claim 7 wherein the zirconia has a particle size in the range of 5 to 50 µm.

9. The material defined in claim 8 wherein the zirconia has a particle size in the range of 10 to 20µm.

10. The material defined in claim 7 wherein the zirconia is fused zirconia.

11. A dense polycrystalline spinal refractory material with a dispersed monoclinic zirconia second phase that has an apparent lower thermal expansion coefficient as compared to the matrix on cooling from the fabrication or operating temperature of the refractory to room temperature and results in the formation of a stable network of micro-cracks that inhibits catastrophic failure as a result of the effects of thermal shock.

12. The material defined in claim 11 wherein the micro-crack network is characterised in that the micro-cracks extend over several matrix grain diameters in the microstructure.

13. The material defined in claim 11 or claim 12 wherein the monoclinic zirconia particles are in the range of 5 to 50µm.

14. The material defined in any one of claims 11 to 13 wherein the monoclinic zirconia particles constitutes 3 to 15 volume % of the total volume of the material.

15. A method of manufacturing a dense refractory material product which includes the steps of:

(i) mixing precursor oxides for a spinal material;

(ii) calcining the mixture to form the spinal material;

(iii)forming the spinal material into a green form of the product; and (iv) firing the green form of the product to produce the final form of the product.

16. The method defined in claim 15 further includes a step of mixing the spinal material produced in step (ii) with an additive, such as zirconia, selected to form a micro-crack initiating phase dispersed in the fired product.

17. The method defined is claim 15 or claim 16 wherein step (ii) is carried out in a temperature range of 800°C to 1600°C.

18. The method defined in claim 17 wherein the temperature range is 1000°C to 1400°C.

19. The method defined in any one of claims 15 to 18 further includes a step of milling the spinal material produced in step (ii) to an average particle size less than 10µm.

20. The method defined is claim 19, when dependent on claim 16, includes adding the additive to the spinal material after the milling step.

21. The method defined is any one of claims 15 to 20 wherein step (iv) includes firing the green form in a temperature range of 1000°C to 1800°C.

22. A dense refractory material which includes a matrix and a micro-crack initiating single crystal phase formed from fused zirconia dispersed is the matrix.