High temperature corrosion of calcium hexaaluminate with biomass slag
Von der Fakultät für Georessourcen und Materialtechnik der
Rheinisch -Westfälischen Technischen Hochschule Aachen
zur Erlangung des akademischen Grades einer
Doktorin der Ingenieurwissenschaften
genehmigte Dissertation
vorgelegt von
Dipl.-Ing. Lise Rebecca Loison
aus Schiltigheim
Berichter: Univ.-Prof. Dr.rer. nat. Rainer Telle
Univ.-Prof. Dr.-Ing Jacques Poirier
Tag der mündlichen Prüfung: 07.12.2018
Diese Dissertation ist auf den Internetseiten der Universitätsbibliothek online verfügbar
Chapter 1: Introduction ...... 1
1. 1. Research question ...... 1
1. 2. Objectives ...... 3
1. 3. Acknowledgements ...... 4
Chapter 2: Bibliography ...... 5
2. 1. Industrial context ...... 5
2. 1. 1. Biomass as alternative combustible for the energy production ...... 5
2. 1. 2. Ash related issues at high temperature ...... 6
2. 1. 3. Refractory lining in biomass incinerators ...... 7
2. 2. Degradation mechanisms of refractory lining in contact with biomass slag ...... 8
2. 2. 1. Molten slag infiltration ...... 8
2. 2. 2. Refractory dissolution ...... 9
2. 2. 3. Structural spalling ...... 10
2. 3. Calcium hexaaluminate as refractory component ...... 11
2. 3. 1. Properties of calcium hexaaluminate ...... 11
2. 3. 2. Refractory applications ...... 12
2. 3. 3. Stability with alkaline ...... 13
Chapter 3: Materials ...... 14
3. 1. Presentation of the system ...... 14
3. 1. 1. Refractory castables ...... 14
3. 1. 2. Al2O3-SiO2-CaO-K2O phase diagrams ...... 16
3. 1. 3. Slag structure ...... 21
3. 2. Refractory samples ...... 23
3. 2. 1. Powder ...... 24
3. 2. 2. Matrix ...... 25
3. 2. 3. Castables ...... 30
3. 3. Slag samples ...... 32
3. 3. 1. Wood ash ...... 33
3. 3. 2. CaO-SiO2-K2O ternary slag ...... 34
3. 3. 3. CaO-SiO2 binary slag ...... 35
3. 4. Conclusion ...... 37
Chapter 4: Identification of reactive mechanisms ...... 39
4. 1. Theory ...... 39
4. 1. 1. Gibbs’energy ...... 39
4. 1. 2. Local thermodynamic equilibrium ...... 40
4. 1. 3. Dissolution in the system Al2O3-SiO2-CaO ...... 40
4. 2. Method ...... 41
4. 2. 1. Generation of reactive interface ...... 41
4. 2. 2. Characterization techniques ...... 42
4. 2. 3. Thermodynamic calculations ...... 42
4. 3. Results ...... 43
4. 3. 1. Solubility limit and thermodynamic equilibriums ...... 43
4. 3. 2. Phase identifications ...... 47
4. 3. 3. Post-mortem microstructures ...... 49
4. 4. Conclusion ...... 58
Chapter 5: Determination of kinetics ...... 61
5. 1. Theory ...... 61
5. 1. 1. Regime of kinetics ...... 61
5. 1. 2. Transport of species ...... 62
5. 1. 3. Influence of viscous boundary layer ...... 63
5. 2. Method ...... 63
5. 2. 1. Reaction couple specimens ...... 63
5. 2. 2. Quantitative X-Ray Diffraction ...... 64
5. 2. 3. Corrective factors ...... 67
5. 3. Results and discussion ...... 68
5. 3. 1. Liquid phase formation ...... 68
5. 3. 2. Matrix reaction ...... 72
5. 3. 3. Phases identification in the microstructure ...... 75
5. 4. Conclusion ...... 78
Chapter 6: Microstructural considerations ...... 80
6. 1. Theory ...... 80
6. 1. 1. Microstructural properties influencing the corrosion resistance ...... 80
6. 1. 2. Microstructure development of calcium hexaaluminate ...... 81
6. 1. 3. Expansive reactions ...... 81
6. 2. Method ...... 82
6. 2. 1. Characterisation of the porous network...... 82
6. 2. 2. Corrosion tests ...... 83
6. 2. 3. XRD measurements ...... 83
6. 3. Results and discussion ...... 83
6. 3. 1. Resistance to structural spalling ...... 83
6. 3. 2. Influence of the porosity ...... 86
6. 3. 3. Morphology of calcium hexaaluminate ...... 88
6. 4. Conclusion ...... 90
Chapter 7: Conclusion ...... 92
7. 1. Advantages of calcium hexaaluminate ...... 92
7. 2. Limitations of the experimental work ...... 93
7. 3. Outlook ...... 94
Chapter 1: Introduction
1. 1. Research question
The energy transition from fossil to biomass fuels challenges the refractory producer on two aspects. On the one hand, products used for the incineration of biomass have to withstand a wide variation of chemical compositions due to the different origins of biomass fuels. On the other hand, they also have to be resistant to alkali, since biomass ashes can contain up to
60 wt.- % of K2O [VAS10]. Chrome alumina, andalusite or mullite are the oxide products currently used to line biomass incinerators [BEN07]. Unlike alumina, calcium hexaaluminate has already proven to have a high stability to alkali, avoiding catastrophic volume expansion known as “alkali bursting”. This property is due to its crystal structure, which has a similar density to the products formed during the reactions with K2O or Na2O, hence no volume variation occurs during the reaction. However, the chemical resistance against liquid slag at higher temperatures was hardly tackled [VAZ09;GEN63], despite the relevance for incineration processes, where the remaining mineral content of fuels tend to melt locally and dissolve the refractory components at the lining interface. Beyond the high resistance to alkali bursting, the possible advantages of calcium hexaaluminate have to be defined regarding its stability to molten oxides, in order to consider this refractory phase as an alternative to alumina in refractory products designed for new energy applications.
Moreover, the understanding of the dissolution of calcium hexaaluminate in slag would be relevant in many different cases:
• During the corrosion of cement bonded refractory castable, where dissolution first takes place within the calcium hexaaluminate matrix formed during the first heating through reactions between cement and alumina, providing mechanical strength to the refractory [YIL06]. • During the corrosion of an alumina aggregate with a calcium rich slag, where a protective calcium aluminate layer is built at the reaction interface [ZHA00;BIL15]. • During the implementation of calcium hexaaluminate commercial products in high temperature processes (e.g. steel, petrochemical and chemical industry, ceramic kilns and glass industry)
Therefore, even if commercial calcium hexaaluminate based raw materials are only starting to emerge in the refractory industry, the knowledge of the dissolution mechanisms and kinetics of
1 this phase could be extended to other corrosion issues encountered in various applications. In order to further illustrate this relevance, the Fig. 1-1 shows the dissolution of a calcium hexaaluminate protective layer around an alumina grain surrounded by slag, formed in an andalusite brick after reacting at 1600°C with molten wood ash. The further corrosion of the refractory aggregate is indirect, governed by the dissolution kinetics of the calcium hexaaluminate layer.
Figure 1-1. Dissolution in slag of the protective layer surrounding alumina grain during the corrosion of andalusite brick with wood ash at 1600°C
The presence of calcium hexaaluminate in refractory microstructure is frequent and its chemical stability controls the corrosion properties of the high temperature lining. The microstructural and chemical aspects of this problematic are tackled in this work considering the corrosion behaviour against biomass slag of calcium hexaaluminate and alumina refractory samples at different levels: powder, matrix and castable.
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1. 2. Objectives
The aim of this work is to tackle the corrosion mechanisms of calcium hexaaluminate with alkali containing slags in order to express the advantages of the implementation of this oxide instead of alumina in refractory products for biomass incineration. The axes of examination are drawn based on the definition of corrosion as an interfacial phenomenon (Fig. 1-2), whose kinetics are controlled both by the chemical reactions and by the transport of the chemical species to the reactive interface [BIL15].
Figure 1-2. Articulation of the different aspects of corrosion
According to this definition, three different aspects of the corrosion mechanisms with slag are addressed:
• The prediction of reactions based on thermodynamic equilibrium calculations combined with post-mortem chemical and mineralogical analysis • The kinetics of the dissolution with molten oxide to express the reaction rates • The microstructural characteristics to consider the physical phenomenon of liquid slag infiltration
As this work was achieved within the framework of the Federation for International Refractory Research and Education (FIRE), the resulting statements aim to support potential development of refractory raw material rather than to report about the performance of calcium hexaaluminate under industrial conditions.
The introductory literature review is reported in the next chapter, and the material system used for this study is described in chapter three. The results are presented and explained in the chapter four to six. The fourth chapter focuses on the identification of the chemical reactions of the selected refractory oxides with wood ash and comparative model slags, based on
3 thermodynamic calculations and microstructural analysis. The fifth chapter tackles the corrosion kinetics experimentally using quantitative analysis of X-Ray Diffraction (XRD) patterns to determine the reaction rates. Finally, the sixth chapter includes the microstructural features of the refractory samples to further interpret the kinetics results obtained in the previous part. The discussion is fed by the experimental conclusions to consider the advantages of the use of calcium hexaaluminate raw materials in new energy applications.
1. 3. Acknowledgements
First and foremost, I would like to express my special appreciation to my advisors Prof. Telle and Prof. Poirier, who have made the achievement of this thesis possible. I want to thank Dr. Tonnesen, who gave me great support and guided me with precious advice along the last years. I am also grateful to Dr. De Bilbao for all its contributions and discussions.
I gratefully acknowledge FIRE C2 that made my Ph. D. work possible, for their financial support, but also for the valuable input coming from the industrial partners: Almatis, Alteo and Imerys.
My time in Aachen was enjoyable due to the amazing refractory team: Petra, Volker, Benjamin, Nicolas, Jonas, Simon, Wanja and Tim, they have become my “Ofenhalle” family and they have made the last years extraordinary. A special thanks to Karolina and Barbara for their culinary and female support. I also want to thank the Orléans team, where I appreciated the three wonderful months.
I am grateful to my friends Sophie, Margaux and Alexandre who have been by my side in the last decade, and Robert, who will hopefully be in my life for the next decades.
A special thanks to my family and especially my parents, Anne and Jean-Luc, who have guided me through life with originality and creativity. I am also grateful to my brotherhood : Adeline, Laurent and Loic, each one has been a role model in their own way, they have known me and motivated me from the outset. A loving thanks to their children too, who make life so colourful. I also want to express my gratitude to my family in Le Havre, who have always welcome me with open arms. Finally, I want to mention my grandmother Micheline and my father Jean-Luc who have encouraged me into this Ph. D. journey, but have not waited for the denouement. This work is dedicated to them and the caring love they have given me. A special recognition for my father, who has accomplished his career in material research at the CNRS an have lead me on the R&D engineering path. 4
Chapter 2: Bibliography
2. 1. Industrial context
2. 1. 1. Biomass as alternative combustible for the energy production
Before the 19th century, wood was widely used as combustible to produce heat and energy. Then the industrial era took advantage of the availability of coal, as fossil carbon feedstock obtained from mine exploration. However, this feedstock can be exploited rapidly, while it needs geological timescale to develop and reabsorb the CO2 released during its combustion. Therefore environmental considerations about global warming are willing to reduce the carbon footprint. The search for sustainable alternatives to fossil energy lead to the replacement of coal through biomass in existing combustion plants. Biomass can be defined as plant and animal waste and refers to fuels coming from living species like wood, grass, food, human. Municipal waste or sewage sludge can also be integrated in this category, under the term of semi-biomass.
The CO2 produced during the combustion of biomass is reabsorbed rapidly to develop further feedstock. For instance during the growth of crops, the CO2 from the atmosphere is consumed through photosynthesis, enabling a neutral carbon balance over a human timescale. In addition to the enhanced renewability of biomass, the combustion of biomass can also reduce harmful emissions of air pollutants like NOx and SO2, depending on the selected biomass source.
In Europe, biomass power generation capacity has taken up to two thirds of all generated renewable fuel power, but only 11% of the fuels used for combustion [LI12]. The incineration of biomass is a thermo-chemical process to transform the chemical energy stored in the organic fuel into heat. The conversion processes are categorized into direct and indirect combustion, where liquefaction, gasification, and pyrolysis are involved. To further operate the existing incineration plants, the replacement of coal with biomass has to be progressive and co- combustion of both is often used. Co-firing is an economic and ecological compromise to reduce greenhouse gas emission on a short term basis since it does not require a large capital investment, but still reduces fossil based CO2 emissions in comparison of the combustion of pure coal [SAM01;SON01;ARV10;ZHA15]. Fluidized bed boilers are widely used to convert biomass into energy, due to their high flexibility and high efficiency. To be environmentally and economically interesting, efficient furnaces with highly resistant lining must be developed to encounter the corrosion issues induced by organic fuels.
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2. 1. 2. Ash related issues at high temperature
The inorganic content of solid fuels is not converted during the combustion and gather as solid waste on the lining wall of incinerators leading to slagging and fouling issues, as the deposit respectively melt or vaporized and subsequently condensed [SEG99;FRA05;WER10;XIA11;VAN12; VAS13;VAS14]. While the temperature in fluidized bed boiler is ranged between 800°C and 1000°C, the temperature in gasifier can reach 1400°C. At high temperatures the mineral ashes liquefy to slag and causes the wear of the refractory materials, reducing the service life of the furnace. The interactions with slag remain the main mechanisms of corrosion and wear of the refractory linings, through dissolution and penetration [BAG14]. Organic fuels have a higher mineral content than fossil fuels, hence the corrosion issues induced by slag formation are enhanced through the use of biomass. The wide variety of biomass results in strong changes in its chemical composition, making it difficult to adapt the refractory composition toward specific solicitations (Fig. 2-1).
Figure 2-1. Average chemical composition and its standard deviation of solid fossil fuels and biomass residues according to [VAS10]
The ash composition has an influence on the melting behavior, limiting the operating conditions and influences also its corrosive properties, like its reactivity with the lining or its viscosity
[BEN07]. Finally, the organic origin of the biomass leads to high quantities of K2O and P2O5 in the residual ash, lowering the melting point of the ash and enhancing the slagging and fouling
6 issues [JON07]. The alkali metals contained in the ash have a tendency to volatize and deposit deep inside a refractory lining and represents a major issue, especially in black liquor gasification [KEI06;BEN07;HEA08;FRO13;BOR17]. The reaction with alkali can lead to strong volume expansions at the surface of the lining, leading to structural spalling and rupture.
Moreover, the alkali can react with SiO2 into low melting silicates causing low-temperature slagging [NIU16]. In order to use biomass or waste as carbon feedstock, it is necessary to adapt the refractory lining for increasing service life. The next generation of refractory, which will be optimized for energy application, will have to show a compromise between a good stability to variations of slag compositions in addition to alkali- and phosphate-resistance.
2. 1. 3. Refractory lining in biomass incinerators
The incineration of biomass in combustion chambers originally designed for fossil fuels causes severe damage to plant equipment and challenges the refractory lining optimization [KHA09;TEI12;SAM01;SON01]. Refractory materials are ceramics with heterogeneous microstructures used at high temperature. They can be divided into two types: the prefired shapes (bricks, tiles…) and castable refractories used by casting, gunning or ramming. While fired bricks are often preferred over monolithics for their higher stability, monolithics are also used for reparation of the lining of incinerators or for smaller installations depending on the construction matters. Refractory products are present in several parts of the incineration process (cyclone walls, external heat exchanger, wind box,…). The area prone to corrosion with molten oxide is the combustion zone with temperatures between 750°C and 1000°C [SCH04;RAU10;
VAN12;SCH04]. In gasifiers, the refractory lining has to show a high resistance to H2 and CO vapor and carbide refractories are often used, like SiC [BEN07;POW10]. Silicon carbide exhibits a high resistance to infiltration of gaseous compounds, like alkali vapor, due to the formation of a silica layer through oxidation [POI09;BRO13;COL14]. Nevertheless, the transition from coal to biomass is harmful for SiC products, which decomposes because of the humidity contained in biomass, for instance in wood. In the combustion chamber, oxide refractories can be found mostly based on Cr-Al2O3 compositions, due to its high resistance to molten slag infiltration, explained by the Cr-spinel layer formation at the slag refractory interface with high density and high refractoriness [YAN90;HIR03;BEN07;GEH13]. However, the formation of toxic hexavalent chrome oxide during application force refractory suppliers to find an environmental and recyclable alternative to Cr-containing products. Aluminosilicate or fused alumina based materials have been used for black liquor gasification, with mostly mullite for fired bricks, while bauxite or chamotte are preferred for unshaped refractory compositions. However the
7 expansive β-alumina or kaliophilite formation, as well as the formation of low melting phases, makes their application unsuitable for alkali rich atmosphere, though the implementation of phosphate bonded bricks seems to minimize the formation of expansive phases [HEA08;SCH04; BEN0;POW10;ANT13;OLE14;LI17].
Designing a refractory lining for biomass incineration is challenging due to the variable chemistry of each biomass and a limited amount of studies has been performed to identify acceptable materials for biomass combustion lining.
2. 2. Degradation mechanisms of refractory lining in contact with biomass slag
In addition to the high temperature resistance, refractories must also withstand a severe corrosive atmosphere leading to chemical wear and reducing the lifetime of the furnace lining. The elevated temperatures increase the reactivity of the refractory lining with its environment leading to corrosion issues at the interface with the slag, such as infiltration, dissolution and structural spalling.
2. 2. 1. Molten slag infiltration
Molten slag infiltration refers to the permeation of the fluid penetrating through the material capillaries without chemical reaction [GOT97]. The impregnation of porous refractory microstructures by molten slag constitutes the physical aspect of liquid corrosion. The extent of penetration depends on the material porous network, the viscosity of the slag and the wettability of the solid by the liquid. The relation between those parameters is described by Poiseuille’s law [POI40]:
푟 cos 휃 훾 푙2 = 푡 ( 1 ) 2 휂
Where l is the slag penetration depth, r is the open porosity radius, θ is the contact angle between the liquid and the solid, η is the slag viscosity, γ is the slag surface tension, and t is time. The surface tension and the contact angle are not sensitive to the chemistry of the slag and they can be considered constant at high temperature [ONI80]. The viscosity of the slag is an important parameter to control the corrosion process and is strongly influenced by the chemical composition and the temperature [LEE99;LEE04].
The resistance of a refractory product against slag penetration can be improved by diminishing the interconnected porous network, i. e. the permeability. Consequently, it is important to
8 characterize the microstructure of the original refractory thoroughly prior to any corrosion test. The infiltration of the slag mostly occurs into the matrix, where the open porosity is located. In the case of unshaped products, the final microstructure is strongly influenced by the microstructure of the hydrated monolith. However, while a dense matrix will hinder slag infiltration, it could lead to explosive spalling by increasing the steam pressure inside the pores during the release of hydrate phases on the first heating.
The reduction of melt penetration through the reduction of the wettability is often used in the metallurgical industry. In aluminum production, non-wetting agent are widely added to the refractory composition, in order to avoid the spreading of the metal on the refractory surface, decreasing the infiltration. The wettability of a refractory with a melt can be assessed using hot stage microscopy, however coherent values can only be determined if the dissolution of the substrate in the melt is hindered. Reducing the wettability of the refractory with the slag prevents the infiltration inside the pores through capillarity, and it represents also an advantage to reduce the interactions between the melt and the refractory surface.
2. 2. 2. Refractory dissolution
In contact with melt, the dissolution behavior of the refractory material in the corrosive liquid is crucial to understand the wear mechanisms and predict the lifetime of the lining. Dissolution of ceramics is a complex phenomenon, nevertheless it is widely tackled in the scientific literature, since it occurs in many different fields, for instance the dissolution of bioceramics in body fluids. Once the liquid has infiltrated and wetted the refractory surface, the interface is submitted to a concentration gradient and approaches thermodynamic equilibrium [POI11;GOT97;LEE08]. To reach the thermodynamic equilibrium and reduce the concentration gradient at the interface, the refractory is being dissolved by the liquid until the saturation limit is achieved. The solubility of the crystal in the slag is therefore required to predict the extent of the dissolution and can be deducted from phase diagrams or by thermodynamic calculations, if the data are available [GOT97;LEE08;FRO13]. Beyond the surface reactions, the transport of the reactants towards this interface is also controlling the dissolution behavior, as well as the transport of the reaction products away from the interface [GOT97;LEE01;BIL15;LEE98].
Depending on the state of the corrosion product, two types of dissolution can be distinguished. If the reaction product is volatile, soluble or removed by the melt flow (erosion), then the further dissolution is direct (congruent, homogeneous or active) and the refractory surface is exposed to slag attack with no intermediate solid phase. The atoms diffuse away from the interface forming concentration gradients in the liquid near the oxide surface. If the reaction product is a 9 viscous melt or a solid, which precipitates at the interface, the surface of the refractory is protected and the dissolution is slowed down by this barrier, the further dissolution is indirect (incongruent, heterogeneous or passive). The sequence of precipitated phases at the interface is determined by the primary crystallization areas encountered across the binary system slag- refractory. The rate-determining steps of the passive corrosion are the chemical reactions forming the layer, diffusion through the layer and diffusion through the slag [OH77;LEE98; LEE04].
In the case of a liquid boundary layer, the compositional gradient depends on the mobility of diffusing species into the slag, which is directly linked to corresponding diffusion coefficient [COO62;COO64;OIS65;GUH97;SAR01;LEE01;POI11;BIL15]. For instance, it was shown that the dissolution of alumina in silicate slag can be controlled by the diffusion of Al ions through a viscous liquid boundary layer in some cases or by the diffusion of Ca ions in a solid calcium aluminate layer in other cases [BIL15;SAR01].
The study of dissolution of solid oxides into molten slags is required to understand those interactions. This knowledge of the mechanisms can be completed by material testing to estimate the resistance of refractory to corrosive liquid. Different procedures are described in the European standard 15418, with static corrosion tests like the crucible test or dynamic corrosion tests like the rotary slag test [DIN60;YAN90]. In order to reduce the solubility of the refractory in the melt, the only axis of improvement is to control the chemistry of the liquid at the interface, for instance by generating a viscous adjacent layer preventing further corrosion or by getting closer to the solubility limit of the slag [GOT97;TAI93;GUH97].
2. 2. 3. Structural spalling
Spalling describes the microstructural damages through cracks formation and propagation, encountered when the refractory material does not have the possibility to expand without generating highly compressive stresses. This heterogeneity of dilatation in the refractory material can be due to temperature, in the case of thermal spalling, for instance a steep thermal gradient will raise strong variations of thermal expansions. If the temperature is well distributed, the spalling is structural and occur through the cohabitation of different mineral phases with different thermal expansion coefficient, and the expansive phase (e.g. calcium dialuminate or hexaaluminate) is limited by the more stable phase (e.g. corundum). The volume increase can also occur at a given temperature, in the case of a phase modification, like ZrO2 transition from monoclinic to tetragonal lattice, or in the case of a phase formation, for instance the in situ formation of Mg-spinel in Al2O3 refractory. Those expansions are predictable during the design 10 of the refractory and can be encountered by an appropriate heating program or composition. However, the structural spalling occurs also during the reaction of the refractory with its environment. The typical situation is the strong volume expansion during the reaction of
Al2O3/SiO2 refractories with alkali, known as “alkali-bursting”. This spalling behaviour is described as the result of a reaction between K2O or Na2O and the aluminosilicate matrix, with the formation of the high-specific-volume reaction products, like feldspar or β-alumina, with theoretical volume increase of 15-30 Vol.-% [LEU84;SCU90;ROU01;SCH09]. This issue is well known in cement kilns, but also in coal combustors and gasifiers [SHA88;RIL96].
Structural spalling is intimately linked with the two previous deterioration mechanisms of infiltration and dissolution. The resulting cracking will favour the liquid infiltration and enhance the contact region for the slag to further react with the refractory. On the other hand, if the material has a good resistance against slag infiltration, it will also be less prone to structural spalling [LEE01;BEN07].
2. 3. Calcium hexaaluminate as refractory component
2. 3. 1. Properties of calcium hexaaluminate
Calcium hexaluminate, CaAl12019 (CaO·6Al2O3=CA6), occurs in nature as the mineral hibonite, which was discovered in Madagascar in 1953 and named after a French prospector Paul Hibon [CUR56]. It crystallises in the hexagonal system with space group symmetry P63/mmc and unit cell parameters a=5.5Ǻ, c=22.0Ǻ [UTS88;MAL98]. Just like β-alumina, it belongs to the hexagonal polyaluminate phases, but with a structure of the magnetoplumbite type. This consists of a layered structure composed alternatively of spinel blocks separated by conduction layer (Fig. 2-2). Ca has 12-fold coordination and Al ions are distributed over five independent crystallographic sites, comprising three octahedral, one tetrahedron and a trigonal bi pyramid [NAG10].
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Figure 2-2. Schematic representation of the calcium hexaaluminate crystal structure and its comparison with β-alumina from [CER98]
Calcium hexaaluminate has an incongruent fusion and is dissociated in Al2O3 and CaO.2Al2O3 at the peritectic point of 1875°C (see further Fig. 3-2). The morphology of CA6 crystals or grains shows preferential growth along their basal plane. Calcium hexaaluminate has the same -6 average thermal expansion coefficient than alumina (αav=8.5·10 /°C), but with a high -6 -6 anisotropy (αa=7.3·10 /°C and αc=11.8·10 /°C) which can lead to thermal expansion mismatch when both phases are coexisting [SAN00]. Hibonite is of high importance in cosmology, since it is considered to be one of the earliest condensation products from the solar nebula. This phase also gains applications in ceramic engineering, for applications as pigment, as host structure for nuclear waste or as fiber coatings in alumina based composite [MOR81;HEN86;CIN94;CIN95; CIN96;AN96;LI16].
2. 3. 2. Refractory applications
In refractory products, calcium hexaaluminate phase is known to form upon sintering within alumina products bonded with calcium aluminate cement. Since last decade, it also exists as sintered raw material on the refractory market, so far exclusively supplied by Almatis due to its complicated synthesis [GAR98;OVE05;KOC05;SAK09;SCH11;SCH12;SAL16;TOM17]. This alumina rich binary compound shares similar properties with corundum, except for its lower thermal conductivity and lower Young's Modulus, resulting in a high crack propagation resistance [NAG90;MAC00]. Calcium hexaaluminate exhibits high stability against various solvents like
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H2/CO atmosphere or alkali attack, which makes this refractory phase a promising candidate for fluidized bed boiler [SCH11].
Concerning the resistance against melt, the implementation of calcium hexaaluminate has already exhibited advantages in the metallurgy [SCH11;BUC05;BUH04]. Indeed this phase shows a low solubility in molten steel, limiting the corrosion of the refractory lining. Moreover calcium hexaaluminate provides an anti-wetting effect for aluminum melt, analog to barium hexaaluminate, but with a higher thermal stability [ENT06].
2. 3. 3. Stability with alkaline
Alkalis represents a major threat in many high temperature process, due to their expansive reactions with aluminosilicate products described previously. The high alkali resistance of calcium hexaaluminate represents the main asset for its use in biomass incineration process.
Like alumina, calcium hexaaluminate reacts with alkaline to form β-alumina (NaAl11O17 or
KAl11O17) and residual calcium dialuminate. However, its crystal structure is really close to the one of β-alumina, and only differs in the arrangement of the stabilizing cations and oxygen anions in the mirror planes between the spinel blocks (Fig. 2-2). This property is a major advantage to avoid the typical alkali bursting.
Moreover, a lower penetration of potassium is observed for calcium hexaaluminate based castables and it is attributed to the calcium hexaaluminate crystal structure, which can incorporate alkalis in the vacant positions inside the spinel blocks without reaction [BUC05; SCH11]. Since alkalis are the major threat in most of the biomass incineration processes, calcium hexaaluminate can be considered to withstand those severe corrosive environment.
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Chapter 3: Materials
The examination of a corrosive refractory system is most of the time hindered by the heterogeneity of its microstructure and the diversity of its chemical composition, including both the refractory and the slag. In order to further interpret the post mortem analysis carried out in this work, three axes are chosen to simplify the complex corrosive system:
• A focus on the castable matrix based on formulations without aggregates to characterize the corrosion properties of the fine fraction of the castable. This region tends to be more reactive and prone to infiltration, because of the higher open porosity and the higher presence of contact interface. • A comparison of the corrosion mechanisms between a biomass slag and two model slag
compositions made of CaO, SiO2 and K2O. This analogy allows to relate the reactions observed with available thermodynamic data. • A reduction of the mineralogical phases contained in the refractory matrix to lower the influence of the secondary phases. The samples were formulated to target alumina and calcium hexaaluminate as the only remaining phase after the sintering.
Considering the complexity of the corrosive system, it seems pertinent to address its physicochemical properties prior to corrosion in the present chapter, in order to deepen the interpretation of the results presented in chapters 4 to 6.
3. 1. Presentation of the system
3. 1. 1. Refractory castables
Refractory castables are a loose mixture of a binding system with refractory aggregates. Its microstructure includes pores, matrix (1/3 of the material) and aggregates (2/3 of the material). In the last century, different types of binding systems were developed to ensure the green strength of refractory products, and transform into ceramic bonds upon sintering. The binding agents of unshaped refractories are still in the core of refractory research since their properties have a major impact on the restrictions during heating and on the resulting strength of the material [PIV98;ROU01;SCH04].
Hydraulic binders like calcium aluminate cement, hydratable alumina or silicate cement, react with water at room temperature and provide green strength through the formation of hydrate phases. The most commonly used hydraulic binders are calcium aluminate cements, which
14 react with water at room temperature to form CaOAl2O310H2O (CAH10), 2CaOAl2O38H2O
(C2AH8), 3CaOAl2O36H2O (C3AH6) crystals. The water of hydration is released during heating and increase the steam pressure in the porosity of the refractory castable. The resulting explosive spalling can be countered through a controlled heating and a highly permeable microstructure. Due to this major drawback and the long setting time, the content of hydraulic binders tend to be reduced or replaced by water free solution like colloidal binders, which is a promising approach for industrial applications [LEE98;SCH04].
The alternative of colloidal binders operates through the coagulation of an ultrafine oxide powder in suspension, like silica sol which is commonly used. As the colloidal particles approach each other, the coagulation occurs through the attractive Van der Wall’s forces. Repulsive forces are present simultaneously due to the overlap of electrostatic double layers present on each particle. At zero potential, also known as the isoelectric point, the colloidal particles coagulate rapidly and forms a 3-dimensional polymeric network of colloidal particles, called a gel. The dispersed state of the sol-gel provides a low viscosity vehicle to carry the alumina aggregates, while the flocculated state enables the castable to immobilize in the desired shape in the mould. This binding system surrounds the refractory aggregates and provide the green strength, which is replaced by a ceramic bonding during the subsequent heating. The utilization of colloidal binders in refractory castables has increased in recent years to avoid both the release of hydrates and the local formation of low melting phases due to the presence of CaO from cement. However the agglomeration of fine particles during the mixing of the castable represents a crucial rheological drawback. Although alumina sols are yet barely used as refractory binders, it presents a possibility to increase the purity of alumina rich castables, remaining with only corundum at high temperature, increasing the hot properties as well as the corrosion resistance [PIV98;AND05;ISM06].
The trend towards the reduction of cement content in castables lead to the following categorisation (Tab. 3-1):
Table 3-1. ASTM classification of high alumina refractory castables [ROU01;AST00]
CaO content [wt.-%] Medium Cement Castable (MCC) > 2,5 Low Cement Castable (LCC) 1,0 < x ≤ 2,5 Ultra Low Cement Castable (ULCC) 0,2 < x ≤ 1,0
No Cement Castable (NCC) 0 < x ≤ 0,2
15
The physicochemical properties of the refractory microstructure, like the type of bond, the pore sizes and the permeability have a strong influence on the wear induced by liquid slag. It is commonly accepted that a denser microstructure leads to a higher resistance to corrosion. [SAR01;SCH04]
3. 1. 2. Al2O3-SiO2-CaO-K2O phase diagrams
In order to clearly define the phases mentioned in this work (Tab. 3-2), the available phase diagrams of the system Al2O3-SiO2-CaO-K2O are reviewed here.
Table 3-2. Mineral phases encountered in this work
Name Mineral Structural formula Oxide formula Cementitious notation JCPDS Card N°
α-Alumina Corundum Al2O3 Al2O3 A 46-1212
Calcium Hexaaluminate Hibonite CaAl12O19 CaO·6Al2O3 CA6 25-0122
Calcium Dialuminate Grossite CaAl4O7 CaO·2Al2O3 CA2 23-1037
Calcium Monoaluminate Krotite/Dmitryivanovite CaAl2O4 CaO·Al2O3 CA 77-4147
β-Alumina Diaoyudaoite (K,Na)Al11O17 (K2O,Na2O)·11Al2O3 ßA 72-0587
- Anorthite CaAl2Si2O8 CaO·Al2O3·2SiO2 CAS2 41-1486
- Gehlenite Ca2Al2SiO7 2CaO·Al2O3·SiO2 C2AS 35-0755
Dicalcium Silicate Larnite Ca2SiO4 2CaO·SiO2 C2S 33-0302
- Leucite KAlSi2O6 K2O·Al2O3·4SiO2 KAS4 38-1423 - Kaliophilite KAlSiO K O·Al O ·2SiO KAS 11-0313 4 2 2 3 2 2
The reactions between refractory samples and molten slags are discussed according to the phase equilibrium data supplied by the following systems: SiO2-Al2O3-CaO, Al2O3-CaO, CaO-SiO2,
K2O-CaO-SiO2 and K2O-Al2O3. Others systems are not discussed in this work due to the lack of thermodynamic data.
16
Figure 3-1. Ternary phase diagram of SiO2-Al2O3-CaO from [HAC15]
The ternary phase diagram Al2O3-CaO-SiO2 (Fig. 3-1) is of major importance in ceramic, metallurgical and petrological field [ERI93;OSB60;RAN15]. For high temperature applications, this diagram is the basis either for refractory composition or for slag chemistry. At atmospheric pressure, two ternary compounds can be encountered in this system: Gehlenite (Ca2Al2SiO7) and Anorthite (CaAl2Si2O8) can exist depending on the temperature and composition [MAO06;OSB60].
17
Figure 3-2. Binary phase diagram of CaO-Al2O3 (wt.-%) from [LEA56]
The binary phase diagram CaO-Al2O3 (Fig. 3-2) is in the focus of many scientific works since the beginning of the 20th century and was completed since then by further experimental studies and the development of chemical thermodynamic models [SHE09;KOH68;NUR65;FIL49; WIS55;WIL62;GEN63;JER01;HAL90;SIN90;ELL81]. The binary phases formed in this system are fundamental for the high alumina refractory cement, and contains five binary compounds:
3CaOAl2O3, 12CaO7Al2O3, CaOAl2O3, CaO2Al2O3, CaO6Al2O3, although the stability of mayenite (C12A7) is related to the moisture content [CHO81;JER01;SHE09;RAN15]. The lower melting temperatures of this phase diagram are located at the eutectics between CA and C12A7 and C12A7 and C3A and represent a major drawback for high temperature applications. Calcium hexaaluminate was recognized in 1949, although it was first accepted as a stable phase in 1963, because of its difficulty to form below 1300°C [FIL49;GEN63;HAL90;WIL62;ALL81].
18
Figure 3-3. Binary phase diagram of CaO-SiO2 from [SLA95]
The binary phase diagram CaO-SiO2 (Fig. 3-3) is well established in the literature, although the liquidus close to the CaO rich side is difficult to assess experimentally, while the miscibility gap complicates the thermodynamic modelling of this system [ERI94;HIL90;TAY90;THO07;BLA87].
This diagram shows two congruently melting compounds: CaO·SiO2, with two polymorphic forms (wollastonite and pseudowollastonite) and 2CaO·SiO2, which exists in four polymorphic forms (Ca-olivine, α'-Ca2SiO4, α-Ca2SiO4 and larnite, which forms by cooling α'-Ca2SiO4).
Two incongruently melting compounds belong as well to this binary system: 3CaO·2SiO2
(rankinite) and 3CaO·SiO2 (hatrutite), which is only stable at high temperature.
19
Figure 3-4. Ternary phase diagram of CaO-SiO2-K2O (wt.-%) from [CHE16]
The ternary phase diagram CaO-SiO2-K2O (Fig. 3-4), which is not completed yet, has been examined by Chen et al., with perspectives for biomass combustion applications or steel production [OHM00;CHE16]. The hygroscopicity and volatility of K2O represents an issue in the experimental assessment of this phase diagram. It is worth noticing that the component
3CaO2SiO2 is not mentioned in this ternary phase diagram.
Figure 3-5. Pseudo-binary phase diagram of K2O·Al2O3-Al2O3 from [SLA95] 20
The binary system Al2O3-K2O is often required since potassium occurs as impurity in alumina products. This binary system is similar to the Na2O-Al2O3 system, which is well-known due to the presence of sodium impurities from the Bayer process during the alumina synthesis. The presence of potassium influences strongly the microstructural development during the sintering of alumina, by leading either to β-alumina formation or abnormal grain growth [SUS85]. Despite its name, β-alumina is a binary compound of alumina and alkali metals. The chemical formula of potassium β-alumina is KAl11O17, however almost all β-aluminas exhibit non-stoichiometric compositions [IYI86;IYI89]. The possible range of chemistry for potassium β-alumina is well depicted in Fig.3-5, which is based on the study of [MOY82] and [ROT80]. Potassium β-alumina is known to melt incongruently at 1877°C The thermodynamic properties of this system are continuously assessed: Beyond the scope of ceramic impurities and refractory corrosion, the high ionic conductivity offers β-alumina applications for energy storage [YAZ06;YAZ11;LEC04; KIM18].
3. 1. 3. Slag structure
The high temperature silicate melt induced slagging is the main ash related issue of biomass combustion. The structure of silicate melts has a great influence on slagging issues during biomass combustion, and has been thoroughly studied in other fields such as metallurgical processes or glass sciences [PAR12;MYS05;NIU16].
Silicate slags are built up of Si4+ cations surrounded by 4 oxygens anions arranged in the form of a tetrahedron. Those tetrahedral (SiO4)4- are joined together in chains or rings by bridging oxygens (BO). Silica is considered acidic and acts as a network former: Its presence in slag increase the polymerization of the melt, hence the high viscosity.
Figure 3-6. Schematic representation of network of tetrahedra formed by Si and oxygen atoms from [SLA95]
21
The degree of polymerization depends on the ratio of non-bridging oxgen (NBO) with the number of tetrahedrally-coordinated atoms (T), denoted as NBO/T. Beside the viscosity, this ratio influences other physical properties of the slag such as its thermal conductivity.
The basicity also indicates the degree of polymerization of the melt. The basic oxides such as calcium or potassium act as network breakers, forming non bridging oxygen, O-, and providing free oxygens, O2-. The resulting depolymerization of the melt decreases its viscosity.
In calcium silicate slag, the crystallisation of calcium rich compounds like larnite (Ca2SiO4) has for effect to increase the viscosity, since the concentration of SiO2 in the melt increases VAZ09.
Some values of viscosities for CaO-Al2O3-SiO2 are given in Tab. 3-3.
Table 3-3. Viscosities of Al2O3-SiO2-CaO slag from [GHO10] and FactSage calculations of liquidus temperature
CaO SiO2 Al2O3 Tliq FactSage Tliq Slag Atlas Viscosity 1600°C Viscosity 1500°C [wt.-%] [wt.-%] [wt.-%] [°C] [°C] [Poise] [Poise] 46.7 31.9 21.4 1525 1515 3.5 15 29.5 52.0 18.5 1380 1300 28 90 39.9 29.7 30.4 1547 1510 8.6 29 37.2 34.0 28.8 1461 1450 17 35 29.3 42.7 28.0 1480 1390 24 71
Alumina is considered to be neutral or slightly acidic as it acts as a network former in high basicity slag, forming polymeric units through the substitution of [SiO4] tetrahedron by [AlO4] tetrahedron and [AlO6] octahedron into the melt structure [GHO10;SLA95]. In this case CaO participates in the charge compensation, according to the reaction Ca2+ + 2Al3+ ↔2Si4+ and the basic cation does not act as a network modifier [CHA99;PAU29]. The resulting higher degree of polymerization of the silicate network increases the viscosity of the slag. As all the Al3+ ions have been compensated, the CaO works as network modifier and decreases the viscosity of the melt.
The addition of K2O to CaO-SiO2 melts increases its basicity, hence decreasing its viscosity.
However, if Al2O3 is present in the melt, the effect of K2O is variant [ZHA12;SUK06]. First the 3+ K2O substitutes CaO in the charge compensation of the Al ions. The chemical bonds around
OAl,K are stronger than for OAl,Ca (Fig. 3-7) and stronger interionic forces can also lead to higher viscosities [ZHA12].
22
Figure 3-7. Change of bridging oxygen from OAl,Ca to OAl,K from [ZHA12]
Once all the Ca participating in the charge compensation have been substituted, the further addition of K2O acts as a network modifier and decreases the degree of polymerization of the slag.
The viscosity influences the mobility of the ionic species in the silicate melt. It plays a major role in diffusion-controlled kinetics of dissolution since a high viscosity slag slower the transport of the species toward the interface [SAL07;POI08].
3. 2. Refractory samples
During the reaction of a refractory castable with slag, the diversity of mechanisms occurring simultaneously can prevent a pertinent interpretation of corrosion study. In order to better relate the post-mortem results with thermodynamic calculations, following orientations were performed:
- Homogenization of the mineralogical composition: Refractory materials are composed of different high temperature phases, already present in the raw materials or formed in situ during the sintering. In alumina rich cement bonded refractory castables, a mixture of calcium hexaaluminate and calcium dialuminate can be found beside of corundum, with singular corrosion behaviours. The particular effects of minor phases are difficult to distinguish in the post-mortem analysis. In order to produce alumina and calcium hexaaluminate and compare their performance with a minimal influence of secondary phases, different kinds of binding are used. The alumina matrix is bonded with a boehmite sol converting into alpha alumina during the sintering. The calcia containing matrix is based on a traditional calcium aluminate cement strengthening through the in-situ growth of calcium hexaaluminate. - Separation of the microstructural elements: Beside the diversity of mineralogical phases, refractories are composed as well of different microstructural features with the matrix phase, aggregates and porosity. In this work, the reactions with slags are
23
examined not only on a full-sized range castable, but also on matrix samples mentioned previously, as well as directly on refractory fine powder.
3. 2. 1. Powder
For the determination of the reaction kinetics with the binary slag CaO-SiO2, a commercial powder of sintered calcium hexaaluminate aggregate (Powder CA6) is provided by Almatis. The chemical composition is shown in Tab. 3-4.
Table 3-4. Chemical composition of refractory powder measured with X-Ray Fluorescence
[wt.%] Al2O3 CaO SiO2 MgO K2O Powder CA6 89.7 9.0 0.8 0.3 0.1
The Powder CA6 has as two major components Al2O3 and CaO. The chemistry of Powder CA6 is enriched in CaO compared to the stoichiometry of a pure calcium hexaaluminate (91.6 wt.%
Al2O3 and 8.4 wt.% CaO), hence there is a lack of Al2O3 to fully form calcium hexaaluminate. This is confirmed with X-Ray Diffraction analysis Fig. 3-8.
Figure 3-8. X-Ray Diffraction patterns of Powder CA6
The Powder CA6 is characterized by the strong presence of secondary phases like calcium dialuminate, traces of corundum, and β-alumina. The remaining presence of both alumina and calcium dialuminate confirm the difficulty to fully form calcium hexaaluminate.
24
Table 3-5. Grain sizes of Powder CA6
[µm] d10 d50 d90 Powder CA6 1.0 7.7 35.5
Tab. 3-5 shows the grain sizes measured with laser granulometry and reveals a wide distribution of the calcium hexaaluminate powder with a d50 of 7.7 µm.
3. 2. 2. Matrix
The aim is to produce two types of refractory matrix with comparable microstructures and different mineral composition. This allows to isolate the effect of calcium hexaaluminate as refractory component regarding its potential to replace alpha alumina for specific biomass applications. The alumina and calcium hexaaluminate matrix are formulated with the same grain size, same water amount, but different chemistry and bonding.
In order to obtain two different chemistry for the matrix, two different binding systems are used.
The binding of the Matrix CA6 is provided by calcium aluminate cement and reactive alumina to react into CA6 during the sintering and the weight ratios are adjusted to obtain the stoichiometry of the targeted phase. The Matrix A is bonded with boehmite sol-gel which irreversibly transforms into α-alumina during sintering. The alumina sol-gel is prepared from a commercial Boehmite powder (AlO(OH)) containing acetate. After dissolving the powder in 80°C-85°C hot water, the dispersion is stabilized through the addition of nitric acid and vaporized to a specific amount.
The formulations of the matrix are based on a reference composition of low cement castable developed through the optimization of the Andreasen packing coefficient and of the rheological behaviour. From this standard composition, only the fine raw materials are selected. The proportion of the water is adjusted to obtain the required rheology for the casting of the samples, but also to obtain comparable open porosity after sintering (Tab. 3-6). The raw materials for the castable samples are provided by FIRE C2 partners: the reactive alumina is supplied by Alteo, the tabular alumina and calcium aluminate cement by Imerys, and the calcium hexaaluminate raw material by Almatis. The dispersant used for the defloculation is supplied by BASF.
25
Table 3-6. Formulations of the matrix samples a) sol-gel bonded Matrix A; b) cement bonded Matrix CA6
a) Matrix A [wt.-%] b) Matrix CA6 [wt.-%] Reactive alumina PFR 27.5 Reactive alumina PFR 20.0 Tabular Alumina 0 - 0.045 mm 30.0 Tabular Alumina 0 - 0.045 mm 25.0 Tabular Alumina 0 - 0.5 mm 41.0 Sintered Calcium Hexaaluminate 0 - 0.5 mm 40.0 Sol-Gel 14.0 CA cement Secar71 15.0 Dispersant FS65 0.15 Water 11.0
After the casting, the sol-gel bonded matrix is kept at ambient conditions, while the cement bonded matrix is put to a humid chamber for 48 h. Afterwards, both types of samples are kept in a drying chamber at 110 °C for additional 24 h. The sintering is performed at 1700 °C for 6 h, with an additional dwell time at 500 °C for the cement bonded matrix to release hydrate phases without damaging the samples. The sintering temperature is chosen as compromise for the calcium hexaaluminate formation (see Fig. 3-10) and the obtention of comparable microstructure for both types of matrix (see Fig. 3-11). The obtention of a calcium hexaaluminate matrix is more challenging, because of the difficulty to fully react Al2O3 and
CaO to form CA6. The precise adjustment of the chemical composition as well as the high sintering temperature permits to obtain a matrix sample containing 99 % CA6 according to Rietveld quantification.
The resulting mineralogy of the samples is controlled with X-Ray Diffraction. The sol-gel bonded matrix leads to a sample mainly composed of -alumina with traces of -alumina, which remain after the sintering.
26
Figure 3-9. X-Ray Diffraction patterns of Matrix A a) after 24h at 110°C; b) after 6h at 1700°C
Matrix A is formulated with raw materials containing mainly corundum, but the tabular alumina is also composed of sodium -alumina due to impurities from the Bayer process. Since the setting at room temperature is only provided by the coagulation, no chemical reaction is required and both alpha and β-alumina remain in the sample after drying (Fig. 3-9). After sintering, the reflexes of β-alumina have lower intensities, while the corundum reflexes have increased. The sodium might have diffused in the liquid phase present at the grain boundaries forming -alumina.
27
Figure 3-10. X-Ray Diffraction patterns of Matrix CA6 a) after 24h at 110°C; b) after 6h at 1500°C; c) after 6h at 1700°C
The hydraulic binding provided by calcium aluminate cement is obtained through different chemical reactions confirmed with X-Ray Diffraction (XRD). The XRD patterns are shown in Fig. 3-10, where the main peaks of the different phases are labelled. After drying at 110°C to release the residual water, tricalcium hydroaluminate 3CaO·Al2O3·6H2O is present from the hydration of the calcium aluminates CaO·2Al2O3, CaO·Al2O3, 12CaO·7Al2O3. Calcium dialuminate remains after the hydration, which is more likely to come from the sintered calcium hexaaluminate raw materials. During the heating the hydrate phase decomposes to release H2O and form calcium monoaluminate, supposedly during the dwell time programmed at 500°C. The calcium monoaluminate reacts with the corundum upon further heating to form calcium
28 dialuminate. At higher temperature, the calcium dialuminate and the remaining alumina further react to form calcium hexaaluminate. The temperature of 1500°C is not sufficient to fully react both phases into the hexaaluminate phase. The enhanced sintering at 1700°C succeeds to react totally the calcium dialuminate, although traces of corundum remain.
For the resistance against slag infiltration, some microstructural properties of the matrix are relevant, like the open porosity, the pore sizes and the permeability to gas. The open porosity and bulk density of both matrix are measured according to DIN 993-1 based on Archimedes’ principle, using a vacuum-pressure technique through the determination of the mass of the sample under different water impregnation conditions. The mercury intrusion technique is used to measure the pore-size distribution of the samples.
Figure 3-11. Pore size distribution after sintering at 1700°C/6h
The rheology of the Matrix CA6 requires lower water amounts during the mixing of the raw materials. However, the resulting open porosity is slightly higher than for the Matrix A. This might be due to the poorer tendency for densification of calcium hexaaluminate reported in the literature, leading to higher open porosity compared to alumina [CRI88]. Although a comparable open porosity was achieved for both matrix, the porous network seems to exhibit different features. The use of a colloidal binder lead to greater average pore diameter of 3.0 µm, while the deflocculated cement binding reaches a lower average pore diameter of 1.0 µm.
The permeability to liquid was determined by a multi-point measurement of gas permeability to determine the intrinsic permeability. The exploitation of the results is based on Dranchuk
29 interpretation reported in [DRA68] and its application to refractory is described in [DEB18] and [LOI17].
Figure 3-12. Permeability of alumina and calcium hexaaluminate matrix
The alumina matrix exhibits a permeable microstructure with an intrinsic permeability of 22.7 mD against 4,9 mD for the calcium hexaaluminate matrix (Fig. 3-12). The permeability value represents the interconnected open porosity and is a relevant parameter for the resistance to infiltration.
3. 2. 3. Castables
Qualitative corrosion tests are performed on full sized range castables, to complete the observations based on the matrix microstructure. Compositions of low cement castables are designed both for alumina and calcium hexaaluminate aggregates, i.e. the castables are mixed with the same binding system of hydrated calcium aluminate cement. The cement is mixed with tabular alumina and reactive alumina for the Castable A, and with sintered calcium hexaaluminate, tabular alumina and reactive alumina for the Castable CA6.
All the raw materials are dry mixed for two minutes. The dispersant is dissolved in the water prior to mixing. The whole content is mixed thoroughly for 7 minutes. The rheology of the calcium hexaaluminate composition exhibits a poorer castability. This can be attributed to the higher open porosity of the aggregates, which imprisons the casting water. Therefore, defflocculant is added to obtain a similar flow consistency.
30
Table 3-7. Formulations of Low Cement Castable (LCC)
a) Castable A [wt.-%] b) Castable CA6 [wt.-%] CA cement Secar71 5.0 CA cement Secar71 5.0 Reactive alumina PFR 12.5 Reactive alumina PFR 12.5 Tabular Alumina 0 - 0.045 mm 10.0 Sintered Calcium Hexaaluminate 0 - 0.045 mm 10.0 Tabular Alumina 0 - 0.3 mm 17.5 Sintered Calcium Hexaaluminate 0 - 0.5 mm 26.5 Tabular Alumina 0.2 - 0.6 mm 10.0 Sintered Calcium Hexaaluminate 0.5 - 1 mm 11.0 Tabular Alumina 0.5 - 1 mm 10.0 Sintered Calcium Hexaaluminate 1 - 3 mm 35.0 Tabular Alumina 1 - 3 mm 35.0 Dispersant FS65 0.15 Dispersant FS65 0.1 Water 5.0 Water 5.0
After mixing the slurry is vibrated in lubricated molds of different shapes (prismatic molds of 25x25x150 mm3, cylindric molds of D50xH50 mm3). The cast specimens are cured in humid chamber within the mold for 24 h and demolded for additional 24 h. After 48 h in humid chamber the specimen are placed in drying chamber for 24 h at 110 °C. The dried specimens are fired at 1500 °C for 6 h with a heating rate of 2 °C/min and a dwell time of 6 h at 500 °C to enable the release of the hydrate phases.
Figure 3-13. Micrograph of Castable A after sintering at 1500°C/6h
The resulting microstructure (Fig. 3-13) is composed of the tabular grains with angular morphology embedded in a cement matrix. Aggregates exhibit intragranular porosity, but most
31 of the micropores are located in the bond phase. The spherical pores with a diameter of approximatively 100 µm are generated by air bubbles during casting.
Figure 3-14. Micrograph of Castable CA6 after sintering at 1500°C/6h
Table 3-8. EDX measurements from Fig. 3-14
[wt.%] Al2O3 CaO Na2O MgO TF 86.2 13.8 1 91.3 8.7 2 78.6 21.4 3 90.9 1.4 4.1 3.7
This high magnification of Castable CA6 microstructure (Fig. 3-14) shows calcium hexaaluminate aggregates in a matrix composed of both of calcium hexaaluminate (Tab. 3- 8 EDX #1) and calcium dialuminate (Tab. 3-8 EDX#2). The third EDX measurement (Tab. 3-8 EDX #3) could correspond to the local formation of a liquid phase or a β-alumina phase containing impurities.
3. 3. Slag samples
Slag composition controls not only its viscosity and wettability, but also impacts the reactivity with the refractory material. This influence results from the interaction and contribution of all
32 the slag elements, hence the difficulty to interpret and isolate the role of specific species from the slag. The present study focuses on three different slag compositions:
• A wood ash, whose composition is representative of the average composition of biomass slag • A ternary model slag composed of the three main oxides of biomass slag, i. e. CaO,
SiO2 and K2O in ratio 2:2:1
• A binary model slag composed only of CaO and SiO2 in ratio 1:1 in order to delete the effect of the potassium oxide. While the experiments with the wood ash enable to characterize mechanisms close to application conditions, the use of model slags allows the comparison with predictions from available thermodynamic data.
3. 3. 1. Wood ash
Wood is the most common source of biomass and its implementation in this work allows obtaining an industry-oriented exposure of the refractory samples. The typical composition of this ash is depicted in Fig.3-15 and compared to the composition of the specific wood ash used in this work.
Figure 3-15. Average composition of wood ash from [VAS15] compared with the chemical composition of the wood ash used in the present study
The slag sample is obtained from the household incineration of wood, whose remaining ashes are gathered and grounded <500 microns.
The melting behaviour of the slag is examined with hot stage microscopy. The device consists of three principal units mounted on an optical bench: a light source to illuminate the specimen, an electric furnace with an alumina tube to heat the specimen and a video camera unit to record
33 the shape variations of the specimen. The different stages are recorded according to the European standard 51730.
Figure 3-16. Recording of the melting behaviour of wood ash on alumina substrate according to European standard 51730
The compacted sample of ash undergoes a strong shrinkage upon heating mostly between 1330 °C and 1410 °C (Fig. 3-16), which originates either from the sintering or from the release of volatile compounds. The apparition of melt occurs around 1430 °C right before the complete melting of the sample at 1440 °C.
3. 3. 2. CaO-SiO2-K2O ternary slag
SiO2, CaCO3 and K2CO3 powders are mixed as reagent grades to obtain a composition with
SiO2-CaO-K2O in ratio 2:2:1. The ternary slag CSK is decarbonated at 1000 °C for 12 h and melted at 1400°C/1h for a final homogenization prior to quenching in water (Fig. 3-17).
Figure 3-17. Micrographs of the amorphous slag CSK after melting at 1400°C/1h and quenching in water
34
Table 3-9. EDX measurements from Fig. 3-17
[wt.-%] MgO SiO2 K2O CaO TF 0.3 42.3 20.3 37.1 1 35.4 3.0 61.6 2 43.1 23.2 33.7
The global composition of the slag (Tab. 3-9 EDX#TF) is in accordance with the targeted composition. Despite the quenching in water, crystallization of the mineral Larnite (C2S) containing K2O impurities (Tab. 3-9 EDX#1) could not be prevented, and the composition suffers local heterogeneity (Tab. 3-9 EDX#1 and #2). The liquid demixing was already observed in the bibliography for potassium containing slag [BER09]. The melting temperature of 1224°C is calculated with FactSage for the ternary slag, which is confirmed with hot stage microscopy.
3. 3. 3. CaO-SiO2 binary slag
The binary model slag CS is obtained by mixing CaCO3 and SiO2 in ratio 1:1, prior to the decarbonatation at 1000°C for 12 hours to enable the complete decomposition of CaCO3. To provide a better homogeneity, the mixed powders are molten at 1550°C/1h and quenched in water down to room temperature. The microstructure resulting from quenching is shown in Fig. 3-18.
Figure 3-18. Micrograph of the amorphous slag CS after melting at 1550°C/1h and quenching in water
35
Table 3-10. EDX measurement from Fig. 3-18
[wt.%] SiO2 CaO TF 53.8 46.2 The slag remains amorphous from the rapid cooling and the homogeneous composition shows a slight enrichment in SiO2 compared with the targetted stoichiometry (Tab. 3-10). This slag composition should melt at 1536°C according to FactSage. However, a melting temperature of 1465°C is determined with the hot stage microscopy. The gap is attributed to the dissolution of the alumina substrate in the melting slag. This was verified by SEM imaging (Fig. 3-19) by performing EDX measurements of the molten slag (Tab. 3-11 EDX #TF1) and the alumina substrate (Tab. 3-11 EDX #TF2).
Figure 3-19. Micrograph of the hot stage microscopy sample after melting the slag CS on an alumina substrate
Table 3-11. EDX measurement from Fig.3-18
[wt.-%] CaO SiO2 Al2O3 TF1 32.3 47.9 19.9 TF2 1.1 98.9
On the micrograph the alumina content in the molten slag reaches the value of 19,9 wt.-%. However, this value is measured after the further heating and cooling of the sample, i. e. it does not correspond to the quantity of Al2O3 dissolved during the in situ recording of the melting.
36
FactSage calculations revealed that a quantity of 8 wt.-% of Al2O3 dissolved in the slag should be sufficient to drop the melting temperature of the composition down to 1465°C.
3. 4. Conclusion
This chapter highlights the limitations induced by the experimental possibilities, which might impede the objectives of this research. Those drawbacks are discussed regarding their impacts on the further results.
The matrix samples presented previously represent the most innovative approach of this work, consisting in the realization of matrix with identical microstructure to isolate the contribution of the chemistry. The fulfillment of this objective is based on the premise that similar grain sizes and similar water content should yield similar microstructures. The results confirm that this assumption applies to the open porosity, however the permeability and porosimetry results revealed microstructures with different porous networks.
Regarding the model slag chemistry, the results of EDX (Tab. 3-9 and 3-10) show an acceptable homogeneity. Moreover, the conservation of the K2O in the ternary slag disclaims the possible volatilization of the potassium despite the high synthesis temperature. The high amount of SiO2 as network former in the slag might have sufficiently increase the polymerization of the melt to prevent the release of the potassium in the gas phase. However, the resulting chemistry of the slag shows a gap with the targeted compositions, with enrichments of silica (Tab. 3-9 and 3- 10). Although those differences are unfortunate, the desired compositions were set based on a literature review of biomass compositions [VAS13]. The resulting chemistry should be sufficient to compare the dissolution behavior of alumina and calcium hexaaluminate and to isolate the effect of the potassium. Moreover, due to different mobilities of the constituents of the slags, the bulk composition of the slag does not represent the chemistry at the slag/refractory interface, and the slight deviation should not impact the effects observed at the reaction front.
The experimental parameters of the different corrosion tests are explained in the next chapters with specific conditions according to the different test procedures. A global summary is already provided in Tab.3-12 to give an overview of the experimental plan.
37
Table 3-12. Overview of the corrosion tests carried out in the present work
Wood ash Ternary slag CaO-SiO2-K2O Binary slag CaO-SiO2
1400°C 1550°C 1400°C 1550°C 1400°C 1550°C
Matrix X X Al2O3 Castable X
Powder X X X
CaO·6Al2O3 Matrix X X
Castable X X X
Since the experiments commented in this work are the results of different studies performed in the framework of FIRE C2 collaboration, the testing conditions are not uniform and prevent sometimes a consistent comparison between the results. The testing conditions were not defined prior to the beginning of the project but have rather evolved along the collaboration to address specific questions and to adapt to the partners interests.
38
Chapter 4: Identification of reactive mechanisms
The determination of the corrosion reactions between the refractory and the slag samples is based on post-mortem analyses of the microstructure and phase composition. The results are compared with phase composition at the equilibrium obtained with the thermodynamic calculation software FactSage. The correlation supports the further steps of this work by describing the mechanisms of corrosion.
4. 1. Theory
4. 1. 1. Gibbs’energy
The reactivity between the refractory material and the slag predicts the corrosion extent of a refractory product. In industrial application, the concept of basicity is used to determine if a material is going to react with its environment [BAG14;MUL11]. For refractory applications, and especially in the steel industry, the basicity is known as the CaO/SiO2 ratio, but its expression depends actually on the complexity of the slag [CHU09;SLA95]. A basic refractory oxide will react with greater extent with acidic slag and will be stable in the presence of a basic slag. Behind the notion of basicity gap, the driving force for thermochemical reactions is the concentration gradient at the reaction front, which tends to bring the system towards the thermodynamic equilibrium for a given composition and temperature [LEE02;POI08;SAL07; BIL15]. Thermodynamic equilibriums are calculated by minimizing the free energy function of Gibbs for a given system (2). This function enables to describe the energy of a phase according to enthalpy (H), entropy (S) and absolute temperature (T):
퐺 = 퐻 − 푇푆 ( 2 )
Moreover the Gibbs’ free energy enables also to characterize a phase modification by considering the variation between the initial state and the final state of energy:
∆퐺 = ∆퐻 − 푇∆푆 ( 3 )
When the phases are in equilibrium, ΔG is equal to zero. In the study of corrosion of refractory with liquid slag, the temperature and chemical composition are the key parameters. For systems with less than three compounds, the regions of phase stability can be represented in 2D phase diagrams sections, while for more than three compounds, the use of computational calculations is recommended to obtain the composition at the equilibrium under given conditions [SAL07].
39
4. 1. 2. Local thermodynamic equilibrium
Microstructural effects prevent some areas of the refractory to come in contact with some areas of the slag. The system will not reach the thermodynamic equilibrium globally, but the fraction of the slag which reacts intimately with the refractory at the interface will reach a local thermodynamic equilibrium [LEE01;LEE02;POI08]. The concept of local thermodynamic equilibrium is then required to analyse open system exchanging chemical species and energy with the surrounding [KOR50;KOR57;THO55;THO59;WEI64].
To predict the further dissolution at the interface, the local composition of the liquid has to be considered rather than the bulk composition. As mentioned in the introduction, the kinetics of corrosion depend both on the reaction rates and on the transport rates of the reactants. If the interfacial reactions are faster than the transport of the species, the local thermodynamic equilibrium can be considered. Otherwise the component mobilities have to be taken into account beside the thermodynamic predictions [GUH97;COO64;OIS65;POI08;JOE91;SAL07]. In the case of a mineral dissolved by a melt, the liquid will dissolve solid phases and precipitate new phases until saturation. At the equilibrium, several solid and liquid phases can coexist, whose number is limited by the phase rule determined by Gibbs (3):
푉 = 퐶 + 2 − Φ ( 4 )
Where C is the number of components, is the number of present phases, and the 2 corresponds to the number of factors of the system: most of the time pressure and temperature. If pressure and temperature are fixed, the equation (4) can be simplified (5):
푉 = 퐶 + 2 − Φ ( 5 )
If the phase rule can be applied to the reactive interface then thermodynamic equilibrium is reached locally [POI08].
4. 1. 3. Dissolution in the system Al2O3-SiO2-CaO
The dissolution of solid oxides into molten slags impacts strongly the furnace life and the knowledge of the individual dissolution behaviour of each refractory constituent is required. Many investigations on the dissolution rate of alumina into molten slags have been carried out, in some cases with the aspiration of a better comprehension of the corrosion of refractory [BAT87;TAI93;SRI00]. During the dissolution of alumina in silicate slag Gehlenite, Hibonite and Grossite are reported to be the major reaction products [BIL15;COO64;OIS65;GUH97;BAT87]. In
40
[OIS65], [SAM64] and [SAN90], a liquid boundary layer at the alumina interface was observed. The composition of the liquid boundary layer depends not only on the saturation limits of the refractory in the slag, but also on the relative mobility of the four ions (Ca2+, Al3+, Si4+,O2-).
The dissolution of the mineral in the liquid can also lead to the precipitation at the interface of a solid phase with a lower solubility in the liquid than the initial phase, passivating the further dissolution. The further dissolution is then controlled by the diffusion of the species inside the solid boundary layer. The precipitation of CA2 coherent layer was attested in [GUH97], while the formation of CA6 can disintegrates the interface due to the large differences in the molar volume between alumina and calcium hexaaluminate. Indeed the formation of CA2 and CA6 is known to be linked with the development of considerable stresses due to the volumetric expansion respectively of 13.6% and 3.01% [GUH97]. The CA6 layer can then exhibit different morphologies: Incomplete or dense, penetrating the grains or at the periphery, which depends on the microstructure of the solid which is being dissolved [SAR00;DOM15;CHE16]. With fused alumina for instance, the layer was dense and continuous, while the layer formed at the porous tabular alumina surface was discontinuous. If the layer is disintegrated with the presence of an amorphous phase, then the dissolution might be direct despite the presence of the interfacial layer. However with increasing temperature a thick, continuous and liquid free layer was observed, showing a fully indirect dissolution requiring mass transport through the CA6 interlayer to access the Al2O3 [SAR00;DEB15;DOM15]. The precipitation of calcium aluminate depletes the slag in CaO and the concentration of SiO2 increases, resulting in a more viscous melt [OGU91;PIL03].
While the dissolution of alumina was widely examined, both as single and polycrystalline material, the dissolution of calcium hexaaluminate is rarely tackled in the literature. The dissolution of calcium hexaaluminate in calcium silicate slags was examined in [VAS09], where the formation of calcium dialuminate was observed both as interfacial layer and as precipitation product from the slag. The examination of the reaction between calcium hexaaluminate and silicon carbide showed the formation of viscous liquid phase as well as Anorthite adjacent layer [CIN98].
4. 2. Method
4. 2. 1. Generation of reactive interface
The corrosion of different types of refractory samples is tested together with different slags under static conditions, based on the description of the crucible test mentioned in the DIN EN
41
15148 [DIN06]. Tests are performed in a resistive electric furnace at different temperatures and for different durations, which are indicated in the caption of the presented results. The cooling to room temperature is mentioned due to its relevance to interpret the post-mortem observations. If the sample is quenched, it can be assumed that the amorphous phase detected at room temperature is related to the molten phase at high temperature, while a natural cooling of the corroded sample can cause the crystallisation of the liquid phase. Care must be taken when interpreting the relation between the thermodynamic calculations and the microstructural results. The methodology used is based on the one found in the literature [LEE02;POI11], relating the microstructural evolution (intergranular liquid composition, aspect of reaction products …) over time and temperature to calculations in order to gain a full understanding of the corrosion mechanisms. For a proper translation of the industrial problematic to an exploitable laboratory work, the critical aspects of the observed solicitations have to be isolated and addressed independently before the transfer to material or process improvement. Therefore the conditions of testing diverge from applications conditions, to enable stable statement about the material selection for biomass incinerators.
4. 2. 2. Characterization techniques
In ceramic field, the chemical reactions are examined by means of X-Ray Diffraction (XRD) and Scanning Electron Microscopy (SEM). After performing the corrosion tests, the samples for imaging are cross-sectioned perpendicularly to the slag-refractory interface, using a diamond wheel. The samples are then impregnated with epoxy resin before polishing with standard ceramographic methods. If the samples are quenched, the cracks developed during the rapid cooling hinder a proper cutting, so the samples have to be embedded directly and cut afterwards. The samples are then carbon coated and examined with backscattered electron imaging SEM, coupled with Energy Dispersive X-Ray Spectroscopy (EDX).
The samples prepared for phase identification with XRD are grounded <63 microns with vibrating disc mill. The XRD patterns are recorded on a Bruker D8 automated diffractometer using Ni-filtered Cu-Kα1 radiation (λ=1.5406 Å). The data are collected in the Bragg-Brentano (θ-2θ) geometry between 5 ° and 90 ° (2θ) in 0.01 ° steps at 96 s.step-1.
4. 2. 3. Thermodynamic calculations
FactSage is a thermodynamic modelling package that contains a database of thermodynamic properties. As with all thermodynamic models, FactSage predicts equilibrium and does not take
42 into account kinetic or microstructural factors, which are tackled respectively in chapter 5 and 6. The Gibbs’ energy minimization module EQUILIBRE is used together with the the databases FactPS, FToxid, FTsalt and slag solution. Beside the targeted composition at the thermodynamic equilibria, the solubility limits of the refractory phases in the slag are determined as well. The temperatures for the calculations are always related to the experiments carried out and mentioned in the description.
Since the composition of the wood ash is too complex for calculations and necessitates the contribution of too many databases, impaired by the lack of thermodynamic data, the first calculations are performed only with the global compositions of the model slags CaO-SiO2 and
CaO-SiO2-K2O.
Once the sample has been penetrated, dissolution of the fine phases occurs fairly rapidly, so that the dissolved phases contributes to the composition of the local liquid, which further attacks the remaining solid. To further relate to the thermodynamic calculations, the thermodynamic calculations in a second part are performed with the phase compositions found locally in the microstructure.
4. 3. Results
4. 3. 1. Solubility limit and thermodynamic equilibriums
As the solid refractory oxide is put in contact with the liquid oxides from the slag, dissolution takes place at the interface, as described in chapter 2, with eventually precipitation of solid phases, in order to reach locally the thermodynamic equilibrium. The calculations in Fig. 4-1, show the solubility limits of alumina (A) and calcium hexaaluminate (CA6) in the bulk composition of each slag (CS and CSK) at temperatures of 1400°C and 1550°C. Although the slag CS is not liquid at 1400°C, its saturation with both mineral phases is mentioned, because the diffusion of aluminium ions in the slag can lower its melting point and temporary form a liquid phase which further precipitate to reach the thermodynamic equilibrium composed exclusively of solid phase. Therefore the term “solubility” is used, although no stable liquid phase is predicted by thermodynamic calculations.
43
Figure 4-1. Calculated solubility limits of Al2O3 and CaO·6Al2O3 in the model slags CS and CSK
As mentioned in the theory, the saturation of oxide melts with mineral phases being dissolved is a data provided by phase diagrams. The Fig. 4-1 shows the slag amount necessary to totally transform each oxide. Below this slag amount, the liquid is saturated with the dissolved mineral phase, which will coexist in equilibrium with the reaction products. Above this slag amount, the starting mineral is no more present at the thermodynamic equilibrium and only reaction products remain in the solid or liquid state. The calculations show a lower solubility for calcium hexaaluminate (CA6), which requires more slag to fully react in both slags. A strong difference is observed in the solubility of CA6 and A at 1400°C with the binary slag CS, while the solubility is almost similar for both mineral phases at 1550°C. In this case, it is worth mentioning that the melting temperature of the binary slag is estimated at 1465°C, therefore the bulk slag is in solid state at 1400°C and in liquid state at 1550°C. CA6 is as well less soluble than A in the ternary slag CSK, with 20% more slag required to fully dissolve at 1400°C, and
18 wt.-% at 1550°C. Except in the case of CA6 dissolved by CS at 1550°C, all the other dissolutions presented in Fig.4-1 are indirect, i.e. resulting in the precipitation of solid phases from the melt. The amount of stable liquid phase is relevant in the study of corrosion, and the products of the reactions tackled in Fig. 4-1. are shown in Fig. 4-2 to 4-5.
44
Figure 4-2. Phase composition at the thermodynamic equilibrium during the dissolution of Al2O3 in slag CS a) at 1400°C; b) at 1550°C
The reaction of alumina with CS slag at 1400°C (Fig. 4-2) tends to the formation of two solid phases: calcium hexaaluminate (CA6) in greater amount and Anorthite (CAS2) in lower amount. Although the formation of those phases can involve a temporary liquid phase, from which the mineral will precipitate, no melt is stable at 1400°C (Fig. 4-2-a). At 1550°C (Fig. 4-2-b), the liquid phase is the most stable reaction product, while calcium hexaaluminate is further formed. Anorthite is not present anymore at the thermodynamic equilibrium.
The Fig. 4-3 shows the reaction products of alumina with CSK slag at 1400°C and 1550°C.
Figure 4-3. Phase composition at the thermodynamic equilibrium during the dissolution of Al2O3 in slag CSK a) at 1400°C; b) at 1550°C
45
Although CSK melts at 1224°C, the reaction with alumina at 1400°C (Fig. 4-3-a) only tends to solid products, mainly calcium hexaaluminate, but also traces of potassium aluminium tectosilicates (KAlSi2O6 and KAlSiO4). At 1550°C (Fig. 4-3-b), the molten phase is stable, however in lower amounts than with the CS slag. While the reaction of alumina with CS slag at 1550°C leads to the formation of liquid and calcium hexaaluminate in ratio 6:1 (Fig. 4-2-b), the reaction with CSK inverts the tendency, with the formation of liquid and calcium hexaaluminate (Fig. 4-3-b). The presence of potassium in the slag composition seems to favour the formation of calcium hexaaluminate, most likely due to the composition of the liquid phase varying toward the precipitation of calcium hexaaluminate.
This behaviour is compared with reactions between calcium hexaaluminate and both slags (Fig 4-4 and 4-5).
Figure 4-4. Phase composition at the thermodynamic equilibrium during the dissolution of CaO·6Al2O3 in slag CS a) at 1400°C; b) at 1550°C
The reaction of CA6 with the binary slag CS at 1400°C (Fig. 4-4-a) forms a stable liquid phase and the mineral Gehlenite (Ca2Al2SiO7), although the bulk slag is in solid state at this temperature. At 1550°C (Fig. 4-4-b), the dissolution of CA6 at this temperature is direct and no solid phase is formed.
The Fig. 4-5 shows the products of the reaction between CA6 and CSK slag.
46
Figure 4-5. Phase composition at the thermodynamic equilibrium during the dissolution of CaO·6Al2O3 in slag CSK a) at 1400°C; b) at 1550°C The calculated phase composition at 1400°C (Fig. 4-5-a) shows the indirect dissolution of calcium hexaaluminate with an increase of the liquid content while calcium dialuminate is formed. The behaviour at 1550°C (Fig. 4-5-b) is similar, although the ratio of liquid phase and solid calcium dialuminate becomes almost 1:1 at the equilibrium. It is interesting to notice that no potassium containing mineral is stable at those temperatures.
To summarize the knowledge acquired with thermodynamic calculations, alumina is less stable in the contact with CaO-containing slag and tends to form CA6 already with small amounts of slag. However, the precipitation of CA6 in the alumina system leads to less liquid phase at the equilibrium. Possible statements for the wood slag will be discussed at the end of this chapter.
4. 3. 2. Phase identifications
The calculated thermodynamic equilibria between CA6 and both model slags (Fig. 4-4 and 4- 5) are compared with the mineralogical composition observed in the Castable CA6 after reacting with the slags at 1550°C for 1h. The corrosion couple consisted in sintered bars of refractory castables covered with pressed bars of slag. After the corrosion test, the samples are quenched in water at room temperature and ground for XRD analysis (Fig. 4-6).
47
Figure 4-6. Castable CA6 a) after sintering; after reacting at 1550°C/1h b) with wood ash; c) with CS slag; d) with CSK slag
After the sintering, the Castable CA6 is composed of :
- Calcium hexaaluminate (CA6) from the raw materials and from the reactions between calcium aluminate cement and reactive alumina
- Calcium dialuminate (CA2) left unreacted from the raw materials - β-Alumina (βA) formed with sodium coming from the Bayer process used for the production of alumina.
After the reaction with wood ash, the intensity decrease of CA6 reflexes comes together with the increase of βA and CA2 (Fig. 4-6-b). The additional β-alumina formed should be composed of potassium rather than sodium, since the wood ash contents 14,7 wt.% of K2O (Fig. 3-15).
The reaction with CS slag only remains with CA6 as the only crystalline phase left (Fig 4-6-c).
This result is in agreement with the Fig 4-4-b, showing the direct dissolution of CA6 with CS slag without formation of solid reaction product. The dissolution of CA2 and Na-βA contained in the castable seems to have only contributed to the formation of either liquid phase or CA6. Thermodynamic calculations verified this results, reporting the formation of calcium hexaaluminate through the reaction of β-alumina with CS slag and the formation of liquid phase through the dissolution of calcium dialuminate.
The XRD pattern resulting from the reaction between Castable CA6 and CSK slag is similar to the one with wood ash, however the formation of K-βA seems to occur in greater extent than
48 for the wood ash. The formation of K-βA is not predicted by the global thermodynamic equilibriums calculated previously (Fig. 4-5-b), hence the formation of potassium β-alumina seems to be controlled by the mobility of potassium toward the interface.
4. 3. 3. Post-mortem microstructures
This part consists of a petrographic study examining the phases and microstructures at the slag/refractory interface. The phases calculated and detected in the previous parts are observed in the microstructure by means of SEM coupled with EDX to determine the morphology and the location of the reactions products.
The Fig. 4-7 shows the interface between Castable CA6 and the molten wood ash after reacting 1h at 1550°C. Just as for XRD analysis, the sample is cooled down to room temperature via quenching in water. Consequently the amorphous phase can be related to the high temperature liquid phase.
Figure 4-7. Micrograph of Castable CA6 after corrosion with wood slag at 1550°C/1h and quenching in water
Table 4-1. EDX measurements from Fig. 4-7
[wt.-%] Na2O MgO Al2O3 K2O SiO2 P2O5 CaO MnO 1 0.4 2.5 52.0 0.3 3.1 2.9 36.3 2.5 2 78.7 21.3 3 1.2 1.9 89.8 5.5 1.1 0.5 4 89.9 0.6 9.6
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The micrograph (Fig. 4-7) reveals three regions:
- The molten slag (Tab. 4-1; EDX #1) with a composition enriched with alumina in
comparison with the bulk composition. This indicates that the CA6 from the castable is already partially dissolved in the molten slag. Small amount of elongated grey crystals
of CA2 are observed in the amorphous phase. The small dark dots are supposedly
crystals of magnesium spinel MgAl2O4 which could have crystallized during cooling
despite the quenching. Hence it cannot be excluded that the presence of CA2 crystals in the molten slag region might come from the cooling. It can happen if the local composition of the melt is really close to the saturation limit of the crystal which precipitates.
- The interfacial layer (Tab. 4-1; EDX#2) composed of CA2. The layer appears disintegrated and discontinuous with porous structure. The cracks are likely to come from the quenching. - The reactive infiltration zone (Tab. 4-1; EDX #3 and #4) where some pore canals are infiltrated with the molten slag (upper bright phase). A complex chemical composition close to K-βA is measured (Tab. 4-1; EDX#3), with also traces of magnesium and manganese which have diffused rapidly from the slag in the refractory microstructure. This mineral is probably the beta alumina phase detected with XRD in Fig. 4-6-b.
This microstructure is compared with the one of the Castable A after corrosion with the same wood ash for 24h at 1400°C (Fig. 4-8). The sample is not quenched down to room temperature, hence the cristallisation observed in the region of the molten slag could be inherent to the slow cooling.
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Figure 4-8. Micrograph of Castable A after corrosion with wood slag at 1400°C/24h and natural cooling
Table 4-2. EDX measurements from Fig.4-8
[wt.-%] Na2O Al2O3 K2O P2O5 CaO 1 0.3 62.3 0.2 1.2 36.0 2 1.2 91.7 4.7 2.4 3 83.3 0.3 16.4 4 83.4 1.7 14.9
Approximately the same zonation can be noticed as in the previous micrograph, with a region of molten slag enriched with alumina (EDX #1), an interfacial layer, and an infiltrated area. The interfacial layer (EDX #3 and #4) has a composition which is neither the stoichiometry of
CA6 nor CA2. It can be assumed that this layer is composed of two monomineral layers parallel to the reaction front, one of CA6 in contact with the alumina and one of CA2 in contact with the slag, as it was already observed in other works [DEB15;DOM15;PAW13;OIS65]. An aggregate of tabular alumina is located at the interface and is surrounded by a reacted layer of K-βA (Tab. 4- 2; EDX #2). The composition of this K-βA is complex with traces of sodium and calcium, which are not necessarily coming from the slag but maybe coming respectively from the tabular alumina and the cement. Despite the natural cooling, thermal expansion mismatch leads to a crack formation at the periphery of the interfacial layer.
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The microstructures of the castables corroded with wood ash are similar, with a region of molten slag enriched in oxide from the dissolved refractory, an interfacial layer composed of calcium aluminate and a diffusion of potassium beyond this layer to react and form β-alumina. The presence of an interfacial layer reveals that the formation of calcium aluminate is faster than the molten slag infiltration in the porous network.
This assumption is confirmed by an infiltration test performed on pressed pellets of Powder
CA6 with wood ash (Fig. 4-9). The pellet is heated with wood ash until 1500°C with a rate of 10°C/min and cooled down to room temperature rapidly. The cooling can be considered as a quenching at air.
Figure 4-9. Micrograph of Powder CA6 pellet after corrosion with wood slag at 1500°C and quenching at air
Table 4-3. EDX measurements from Fig. 4-9
[wt.-%] Al2O3 CaO 1 78.2 21.8 2 91.2 8.8
The interfacial layer is really similar to the previous one in Fig. 4-7, with a thickness of approximately 50 microns and the stoichiometry of CA2 (Tab. 4-3; EDX #1). Below this layer only CA6 is observed (Tab. 4-3; EDX #2) and the molten slag has barely infiltrated the pellet
52 despite the enhanced porous network. The dissolution of calcium hexaaluminate enriched the slag with CaO, leading to the nucleation and growing of the calcium dialuminate crystals. The dissolution/precipitation mechanism is confirmed to be faster than the melt infiltration. Despite the several open capillaries for slag infiltration, the formation of the calcium dialuminate protective layer prevent any further slag penetration.
The Castable CA6 is then tested 1h at 1550°C with CSK slag followed by quenching in water down to room temperature. Although the reactions with CSK slag showed the same resulting mineralogy (Fig. 4-6-d) than for wood ash, the microstructure revealed another distribution of the phases (Fig. 4-10).
Figure 4-10. Micrograph of Castable CA6 after corrosion with CSK slag at 1550°C/1h and quenching in water
Table 4-4. EDX measurements from Fig.4-10
[wt.-%] Al2O3 SiO2 K2O CaO 1 90.6 0.5 0.2 8.7 2 77.5 0.7 21.8 3 48.6 21.4 6.5 23.5
No formation of interfacial layer is observed, and the precipitation of CA2 (EDX#2) occurs locally in the microstructure in the vicinity of the CA6 grains being dissolved (EDX#1). The
53 microstructure of the CA2 is dense and typical of a precipitate, with a size of approximately 20 microns. The CA6 grains have intragranular pores, characteristics of sintered refractory raw materials. The formation of β-alumina is also witnessed in this microstructure (Fig. 4-11).
Figure 4-11. Micrograph of Castable CA6 after corrosion with CSK slag at 1550°C/1h and quenching in water
Table 4-5. EDX measurements from Fig. 4-11
[wt.-%] Al2O3 SiO2 K2O CaO 1 91.0 8.3 0.8 2 53.0 19.7 6.8 20.6 3 90.2 0.7 8.3 0.8
The K-βAl2O3 observed with EDX#1 shows also intragranular porosity, indicating that the reaction has more likely occurred through diffusion of potassium in a solid CA6 grain. On the contrary the crystal observed in EDX#3 shows the same chemistry but its dense microstructure indicates rather a precipitation from the melt.
It seems that the potassium has a higher mobility in the slag than the other two ions (Si4+ and 2+ Ca ) and reacts preferentially with CA6 before any dissolution takes place. The reaction products between calcium hexaaluminate and potassium is calculated with FactSage:
54
Figure 4-12. Phase composition at the thermodynamic equilibrium during the reaction of CaO·6Al2O3 with K2O a) at 1400°C; b) at 1550°C
The β-alumina is a stable reaction product only when CA6 reacts directly with K2O (Fig. 4-12).
The CA6 loses alumina through the β-alumina formation and transforms then into CA2. It is worth noticing that the FactSage database indicates a temperature of vaporization of 1573°C for K2O, implying that the reaction observed in Fig. 4-11 is between calcium hexaaluminate and liquid potassium oxide.
If β-alumina is formed, then the slag CSK becomes depleted in K2O with a composition tending toward the chemistry of slag CS. The thermodynamic equilibriums resulting from the further dissolution of K-βA in CSK and CS slags are as well simulated:
Figure 4-13. Phase composition at the thermodynamic equilibrium during the dissolution of K2O·11Al2O3 in slag CS a) at 1400°C; b) at 1550°C
55
Figure 4-14. Phase composition at the thermodynamic equilibrium during the dissolution of K2O·11Al2O3 in slag CSK a) at 1400°C; b) at 1550°C
As the β-alumina is further dissolved by the slag, the CA6 and CA2 are present at the equilibrium as reaction products (Fig. 4-13 and 4-14).
The following figure shows the microstructure of the Castable CA6 near the penetration front after the corrosion with CS slag at 1500°C for 1h and quenching in water down to room temperature (Fig. 4-15).
Figure 4-15. Micrograph of Castable CA6 after corrosion with CS slag at 1550°C/1h and quenching in water
56
Table 4-6. EDX measurements from Fig. 4-15
[Wt.%] Na2O Al2O3 SiO2 CaO 1 0,4 42,8 30,2 26,7 2 91,4 8,6
The bright slag (Tab. 4-6; EDX#1) has penetrated the fine matrix porosity and the slag composition enriched in alumina indicates the ongoing dissolution of the CA6 (Tab. 4-6; EDX#2) by the melt, without any solid reaction product, in agreement with the XRD pattern (Fig. 4-6-c). However, in the vicinity of the alumina grains, the dissolution lead to the local precipitation of CA6 (Fig. 4-16) which is also consistent with the thermodynamic calculations (Fig. 4-4-b) and the XRD results (Fig. 4-6-c).
Figure 4-16. Micrograph of Castable CA6 after corrosion with CS slag at 1550°C/1h and quenching in water
57
Table 4-7. EDX measurements from Fig. 4-16
[wt.-%] Na2O Al2O3 SiO2 CaO 1 91.5 1.0 8.3 2 0.3 40.5 31.2 28.0
This precipitated CA6 (Tab. 4-7; EDX#2) is dense and shows an oriented growth in a tabular microstructure. Locally this system is invariant with three chemical elements (Al2O3, SiO2 and CaO) distributed in three phases (corundum, calcium hexaaluminate and the melt), hence the liquid composition should be constant.
4. 4. Conclusion
The correlation between post-mortem examinations and thermodynamic calculations enables to determine the reaction mechanisms occurring between the castables and the molten wood ash (Fig. 4-17).
Figure 4-17. Schematic comparison of the monomineral layers formation in contact with molten wood ash; a) Castable A b) Castable CA6
Both Castables A and Castable CA6 show a slag infiltration hindered by the formation of an interfacial layer of calcium aluminates: CA2 in the case of Castable CA6 (Fig. 4-7) and a succession of monomineral layers of CA6 and CA2 in the case of Castable A (Fig. 4-8). Beyond this protective layer, only a few species from the slag have diffused, like manganese, whose high mobility is already attested [SAR01], but also potassium ions which penetrate further in the refractory to form β-alumina.
The comparison with the reactions products detected with XRD (Fig. 4-6) shows similar patterns between the reaction of Castable CA6 with wood ash and the CSK model slag. However the micrographs show a different phase distribution resulting from a different sequence of mechanisms (Fig. 4-18). The composition of the model slag lead to a quick infiltration of the melt in the refractory microstructure (Fig. 4-10 and 4-11). The precipitation of calcium
58 dialuminate is found locally in the vicinity of CA6, which is being dissolved by the model slag, without any formation of protective layer at the interface. The morphology of the precipitated products seems to depend on the slag chemical composition and should be further addressed to deepen this interpretation.
Figure 4-18. Schematic comparison of the corrosion mechanisms observed for the model slag CSK and the molten wood ash
The qualitative characterization of the reaction mechanisms have demonstrated the relevance of the model slag to predict the possible solid corrosion products. However, due to the complexity of the wood ash, the composition of the molten phase does not match the model system. Especially the influence of P2O5, which is a major compound in living organisms and biomass composition, should not be underestimated. This oxide acts as a network former and increase the polymerization of the slag, substituting silicium oxide. Moreover the presence of phosphorus oxide increases the melting temperature of the system, and can also favour the formation of refractory solid phases. A recent study has shown that the increase of P2O5 in biomass slag can encounter the impregnation in the refractory microstructure through the increase of the liquid viscosity [POI18]. Furthermore, the pertinence of the binary model slag is confirmed: The potassium of the ternary slag shows a higher mobility than the other elements and is the first species to react with the castable. Afterwards, the slag is depleted in K2O and the further dissolution mechanisms with the ternary slag are comparable with the mechanisms observed with the binary slag.
Some of the thermodynamic predictions could be proven experimentally. For the binary slag, the direct dissolution of CA6 is indeed observed (Fig. 4-6-c), while the ternary slag dissolved
CA6 indirectly with the precipitation of CA2. Nevertheless, some calculations did not match the post mortem results. In the case of the ternary slag, the thermodynamic model predicts no stable potassium containing solid phase, while the β-alumina formation is confirmed both by XRD
59
(Fig. 4-6-d) and SEM (Fig. 4-11). If the thermodynamic predictions do not corroborate the experimental observations, the kinetics of corrosion are controlled by the transport of the species to the reaction front. Hence the mobility of potassium is higher than the one of the other two constituents of the slag. The formation of β-alumina occurs before the further dissolution with a slag chemistry depleted in potassium.
60
Chapter 5: Determination of kinetics
The knowledge of reaction kinetics is required to predict the service life of a refractory product. This chapter tackles the resistance of calcium hexaaluminate, testing the direct dissolution in binary slag at 1550°C and the indirect dissolution in ternary slag at 1400°C. The comparison with the dissolution behaviour of alumina matrix is also addressed.
5. 1. Theory
5. 1. 1. Regime of kinetics
Dissolution is an activated interfacial process, whose rate is a function of many variables including temperature, interface composition and liquid viscosity. The dissolution at the interface solid/liquid occurs via three steps:
▪ Transport in solution of the reactants to the interface ▪ Surface reaction with adsorption of the reactants ▪ Transport in solution of the reaction products away from the interface
The transport rates of the species (physical process) and the reaction rates (chemical process) control the dissolution kinetics. The reaction rate represents the rate applied by the system to reach the equilibrium state, or a state of lowest free energy called metastable state (Fig. 5-1).
Figure 5-1. Energy barrier diagram with Q the activation energy and H the reaction heat from MCC04
The temperature increases the reaction rate exponentially usually according to Arrhenius equation (6):
푅푒푎푐푡𝑖표푛 푟푎푡푒 = 퐴푒−푄/푅푇 ( 6 )
61
Where A is a preexponential constant, Q is the activation energy in J/mol, T the absolute temperature in K and R the universal gas constant.
If the surface reaction is slower than the transport of the species, the reaction restricts the dissolution kinetics. Otherwise the corrosion is limited by the transport of the elements, i. e. diffusion controlled. The reaction controlled dissolution requires higher activation energy than the transport controlled dissolution. In refractory corrosion, the wear rates are often submitted to a mixed control: Both reaction and transport govern the intermediate regime, which must be addressed experimentally. The slag composition influences which mechanism is predominant on the rate of dissolution. In dynamic corrosion test, for instance the rotary slag test, if the dissolution rate depends on the revolution speed, then the mass transport of solute in the molten slag controls the dissolution rate [TAI93;YAN90;LEE99].
5. 1. 2. Transport of species
The transport of species can occur through the fluid flow or through solid diffusion. Fluid flow can occur as forced convection if the slag is moved by an external source or free convection resulting from differences in the density or surface tension between the saturated slag and the bulk slag [COO64;OIS65]. In the case of a moving slag, the erosion impacts the corrosion process and the viscosity of the slag is an important parameter to consider. If no fluid motion is observed, the species migrate from the interface by diffusion toward the region of lower chemical potential [BAT87]. Diffusion is described by Fick's first law (7) [FIC55;BIR60]:
휕푐 퐽 = −퐷 ( 푖) ( 7 ) 휕푥
Where J is the flux of component in the x direction in mol.m-2.s-1, D is the diffusion coefficient 2 -1 -3 in m .s and ci is the concentration of component i in mol.m . The molar flux due to diffusion is directional and proportional to the concentration gradient.
Fick's second law describes the nonstationary state of flow where the concentration of a fixed region varies with time t (8):
휕푐 휕2푐 푖 = 퐷 푖 ( 8 ) 휕푡 휕푥2
In refractory application, the diffusion coefficients are difficult to take into account, since they depend on the composition and the structure of the material through which diffusion occurs.
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Nagata et al. reported that the diffusion coefficient of calcium is higher than that of aluminium, which is higher than that of silicium [NAG82]. Diffusion of species can also be induced by pressure or temperature gradient [SOR79;LUD56].
5. 1. 3. Influence of viscous boundary layer
In a transport controlled dissolution, a boundary layer can build up at the interface saturated with the reaction products. The dissolution process is then governed by the diffusive transport of species through the boundary layer. The thin layer of solvent adjacent to the solute becomes rapidly saturated and remains saturated during the dissolution process, while beyond this layer the concentration of the bulk solution remains [GUL06]. The flux of component J can be expressed by the Nernst equation [NER04]:
(퐶 − 퐶 ) 퐽 = 퐷 푆 푚 ( 9 ) 훿
-3 Where Cs the saturation concentration of the refractory in the melt in g.m , Cm the concentration of reactant in the melt and δ the effective boundary layer thickness in m.
If the boundary layer leads to formation of a solid interface the further dissolution is indirect and the process is passivated. In this indirect case, the rate limiting step is the solid state diffusion through the solid interfacial layer [ZHA00;GUH97;LEE04;MON04]. The composition of the boundary layer is depending upon the dissolution, whether diffusion or reaction is predominant [MCC04].
5. 2. Method
5. 2. 1. Reaction couple specimens
Two types of samples are used for the kinetics determination:
▪ Powder sample (Powder CA6) for the direct dissolution of calcium hexaaluminate with the binary slag CS. The ex situ measurements are performed at the CEMHTI laboratory in Orléans, in the framework of FIRE C2 collaboration and supported financially by JECS Trust.
▪ Matrix samples (Matrix A and Matrix CA6) for the indirect dissolution of calcium hexaaluminate with the ternary slag CSK.
The corrosion kinetics are tackled mainly through ex-situ trials, where the corrosive interface is quenched prior to XRD measurements for quantification. The ex-situ trials are completed
63 with in-situ determination of the reaction kinetics, where the unreacted powder mix is placed on a Pt-stripe and heated for the high temperature XRD measurements.
For the reaction on powders, two different weight ratios of Powder CA6 : Slag CS are prepared: 60:40 for the ex-situ trials and 85:15 for the in-situ trials. The slag and refractory powders are homogenized in mortar with isopropanol. After the drying of the samples, the powder used for the in situ determination is placed directly on the heating stripe of the high temperature chamber in the diffractometer. The powder for the ex situ quantification is placed in a Pt-crucible and inserted in the resistive laboratory furnace directly at 1550°C for different durations before quenching in water. The ex situ powder trials are repeated to obtain three values per experimental condition.
To test the kinetics of the matrix reactions, the refractory samples (Matrix A and Matrix CA6) are cut in the dimensions of 5x10x30 mm3. The CSK slag is mixed with a few drops of ethanol and pressed uniaxially into a pellet with a surface of 25x5 mm2 and a varying mass of 0.3g, 0.6g and 0.9g to observe the influence of the slag amount on the reaction rates. As for the ex situ determination on powders, the matrix sample covered with the slag pellet is put directly in the furnace at 1400°C for different durations before quenching in water down to room temperature.
In this corrosion test, the quantity of slag is low and might limit the reaction once the slag is saturated with the refractory. Moreover the testing is static, i.e. without movement of fluid, avoiding erosion effects and lowering the transport rates. The temperature distribution is considered homogeneous across the sample volume and the matrix samples are not pre-heated. It can be considered that they undergo a severe thermal shock as they are introduced in the furnace at 1400°C. However the influence of this microstructural damage is assumed to be negligible on the resulting phase composition.
5. 2. 2. Quantitative X-Ray Diffraction
The XRD measurements are performed with a D8 Advanced diffractometer with conditions displayed in Tab. 5-1.
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Table 5-1. XRD measurement parameters
Step Time/step Angular range Scan type D8 Bruker [°] [s] [°] Slag CS Ex situ 0,021 384 15-80 Coupled TwoTheta/Theta CEMHTI, Orléans In situ 0,015 10 32-34,927 PSD fixed GHI, Aachen Slag CSK Ex situ 0,020 384 5-90 Coupled TwoTheta/Theta GHI, Aachen The in situ measurements are performed with the high temperature chamber HTK2000 supplied by Anton Paar and equipped with a platinum heating stripe. The sample is heated with a high heating rate of 200 K/min before the dwell time at 1550°C to obtain a time resolved XRD measurements of the main reflex of calcium hexaaluminate.
The recorded XRD patterns contains information about the ratio of the different phases present in the sample after reacting with slag, i. e. the extent of the reaction. Based on the experimental data, a quantification is required to determine the amounts of the coexisting phases. The quantification of the crystalline compounds is not trivial and can be approximated either through the evaluation of the peak areas or with a Rietveld analysis for more accuracy [WIL62; RIE69;DIC69;CHU74;BRI85;SIN90;FUL99]. This latter method is a global analysis of the diffraction data taking into account both the position and the intensity of the reflexes to compare the experimental input with a simulated diffractogram calculated from a structural model [RIE69;HIL87;BIS88]. In this work, the Rietveld refinement is performed with the software TOPAS from Bruker using the fundamental parameters. The preferred orientation of calcium hexaaluminate is taken into account in the structure card of CA6.
The amorphous content of the ex situ sample should relate to the amount of high temperature liquid phase due to the rapid cooling. Several methods are available for its quantification, with different experimental effort and different accuracy [MAD11;TOR01].
For direct dissolution of Powder CA6 with the binary slag CS, the only reaction product is the liquid phase, whose amount is determined by the method of the internal standard. This method is an indirect calculation, i.e. the sample has to be mixed with a standard powder and the absolute abundances of the crystalline components are used to estimate the amorphous content by difference. This calculation is based on the normalization condition of Rietveld refinement, which calculates the crystal content equal to 100 wt.% (10) [MAD11]:
100 푋 푠 ( 10 ) 푋푎 = (1 − ) (100 − 푋푠) 푋푠,푐
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A laboratory synthetized powder of calcium dialuminate is used as internal standard, whose value for microabsorption is reported in Tab. 5-2
Table 5-2. Absorption coefficient of the main phases with Cu-Kα1
Phase Linear absorption coefficient [cm-1] Corundum 120.6 Calcium hexaaluminate 142.7
Calcium dialuminate 142.8
The method of the internal standard yields consistent results [TOR01], however the powder sample has to be mixed with a perfectly crystallized standard with comparable microabsorption coefficient and with distinctive reflexes without superimposition of the 2θ values. For this last reason, the use of internal standard for the ternary slag CaO-SiO2-K2O is delicate, since five different crystalline phases are expected in the sample after corrosion (alumina, potassium β- alumina, Gehlenite, calcium dialuminate and hexaaluminate). Hence the selection of an appropriate standard might hinder the Rietveld quantification. Therefore, the direct method of Degree Of Cristallinity (DOC) is used to quantify the liquid phase for dissolution with the ternary slag. This method calculates the amorphous amount by relating the surface from the crystalline reflexes and the surface from the amorphous bump (11 and 12) [MAD11]:
퐶푟푦푠푡푎푙푙𝑖푛푒 푎푟푒푎 퐷푂퐶 = ( 11 ) 퐶푟𝑖푠푡푎푙푙𝑖푛푒 푎푟푒푎 + 퐴푚표푟푝ℎ표푢푠 푎푟푒푎
( 12 ) 푋푎 = 1 − 퐷푂퐶
The software EVA from Bruker proposes a tool to calculate the crystallinity based on the area of the peaks. This method can yield a consistent determination of the amorphous content, however it is really dependant on the adjustment of the background and on the angular range used for the calculation. The direct and indirect methods are compared by establishing a
66 calibration curve, where perfectly crystalline alumina and calcium dialuminate powders are mixed with the amorphous slag CS in different ratios (Fig. 5-2).
Figure 5-2. Comparison of the different methods for amorphous quantification on calibration powder mixes
The results confirm the consistency of the indirect method with the internal standard. To obtain results representative of the real amorphous content with EVA, the background is not modified manually and the same angular range of 20-70° is used for all the calculations. It is worth noticing that the reliability of the quantification is limited for small amorphous content, which is the range expected in the indirect dissolution with CSK slag (between 10 and 20 wt.-%).
5. 2. 3. Corrective factors
To further reduce the uncertainty of the amorphous quantification, the liquid amount yield by the reaction is determined by subtracting the amorphous content calculated for the uncorroded sample.
퐿𝑖푞 = ∆퐴푚표푟푝ℎ표푢푠 = 퐴푚표푟푝ℎ표푢푠 (퐶표푟푟표푑푒푑) − 퐴푚표푟푝ℎ표푢푠 (푈푛푐표푟푟표푑푒푑)
Moreover, since deviations exist in the mass of the refractory sample due to the preparation, a corrective factor is applied to the final quantification. The correction is based on the assumption that the slag is the limiting reactant, so that the saturation of the slag occur before dissolving the whole refractory sample. If the mass of the refractory sample is higher, the supplement
67 contribute only to the increase of the uncorroded fraction. This assumption is confirmed by performing the test on two sample sizes, which reveals that the difference in mass corresponds to an increase of the uncorroded section.
5. 3. Results and discussion
5. 3. 1. Liquid phase formation
In order to determine the liquid phase formation during the dissolution of calcium hexaaluminate, powder samples are heated at 1550°C in order to quantify the direct dissolution of Powder CA6 in the binary slag CS (Fig. 5-3). The initial refractory:slag ratio is 60:40 and the dissolved amount is determined after quenching down to room temperature with the method of the internal standard. The liquid content at the thermodynamic equilibrium is calculated with FactSage considering the chemical compositions from Tab. 3-4 and 3-10. Each trial is repeated three times and the average as well as the standard deviation is depicted in Fig. 5-3.
Figure 5-3. Ex situ quantification of the amorphous content after the direct dissolution of Powder CA6 in CS slag at 1550°C
The dissolution of the Powder CA6 in the slag CS is fast: After five minutes at 1550°C, the amorphous content already reaches the value of 84.2 wt.-%. The amorphous content decreases slightly with the corrosion duration, reaching the value of 81.5 wt.-% after 2h. Since no other crystalline phase is observed in the sample beside calcium hexaaluminate, the decrease of the amorphous content is attributed to the precipitation of calcium hexaaluminate from the liquid phase. Since the dissolution of pure calcium hexaaluminate could not yield this precipitation
68 based on thermodynamic laws, it is attributed to the remaining secondary phases in the powder (corundum and grossite), and will be further discussed based on the following in situ measurements. The amount of precipitated calcium hexaaluminate after two hours is 2.7 wt.- %. The standard deviation increases with reaction duration, supposedly because the arrangement between the liquid phase and the remaining calcium hexaaluminate might have a greater impact on the proceeding of the dissolution.
However, the first ten minutes of the dissolution can hardly be tackled with ex situ trials, and in situ XRD measurements are performed on the Pt-heating stripe of the high temperature chamber. The powder CA6 is mixed with the slag CS in ratio 85:15 and heated at 1550°C with a heating rate of 200K/min. At 1550°C, the LynxEye 1D detector equipped with 192 Si-stripes remains open on the main reflex of calcium hexaaluminate, enabling a time resolved recording of its intensity every 10 sec in the angular range 32.00-34.93° [2θ]. The intensity is plotted on the Fig. 5-4
Figure 5-4. In situ recording of the intensity of the calcium hexaaluminate peak by 34.156° [2θ]
The increase of the intensity over time can indicate the precipitation of crystals from the molten phase (Fig. 5-4), in agreement with the ex situ quantification of the amorphous phase (Fig. 5- 3). A scan with a wider angular range 10°-90° is performed at room temperature before and after the in situ measurement. The zoom in 10°-40° range is shown on Fig. 5-5.
69
Figure 5-5. Diffractograms at room temperature of Powder CA6 and slag CS in ratio 85:15 before and after the insitu measurement at 1550°C
The diffractograms before and after the insitu measurements show indeed the increase of the amorphous phase as well as the increase of the calcium hexaaluminate reflexes. However, traces of corundum (35.15°) and calcium dialuminate (25.43°) are present in the raw material prior to the reaction. The hypothesis can be drawn that the traces of alumina and calcium dialuminate dissolved rapidly in the slag to form liquid and calcium hexaaluminate, which is further dissolved.
Both scans are quantified with EVA (amorphous phase) and TOPAS (crystalline phases) and values are given in Tab. 5-3. The absolute amount of amorphous phase has to be consider with care to prevent misinterpretation.
Table 5-3. Quantification of the XRD diffractograms from Fig. 5-5
[wt.-%] Amorphous phase CA6 CA2 A Before 10 min 1550°C 14 80 4 2 After 10 min 1550°C 33 67 Thermodynamic equilibrium 33.4 65.9 0.6
The increase in amorphous content is really high and the resulting ratio of CA6 and amorphous phase is close to the expected composition at the thermodynamic equilibrium. The sample size is lower than for the ex situ trials and the global thermodynamic equilibrium seems to be
70 achieved. As mentioned qualitatively, the traces of calcium dialuminate and corundum disappear which is in agreement with thermodynamic predictions.
The intensities measured at high temperature are compared with the intensities measured at room temperature with the same scan parameters (Fig. 5-6).
Figure 5-6. Scans with Position Sensitive Dectector (PSD) fixed at 32°-34.927° during the dissolution of Powder CA6 and slag CS in ratio 85:15 at 1550°C
The intensities at high temperature are lower than the one recorded at room temperature. The shift of the peak position after heating is due to the thermal expansion of the crystal lattice and is reversible after cooling down to room temperature. However, the peak at room temperature after cooling is slightly shifted to the right compared to the peak position before heating. This shift might result from the change in the sample thickness through the thermochemical reactions linked with density variations.
The comparison of the intensities measured before and after the in situ experiment helps to draw the different steps of the reactions:
▪ During the heating, about eight minutes are required to reach the dwell temperature of 1550°C, the intensity of the calcium hexaaluminate main reflex decreases: The phase is being dissolved by the liquid. The calcium dialuminate and alumina are supposedly as well fully dissolved after the temperature of 1550°C has been reached. ▪ At 1550°C, the intensity increases over time: Calcium hexaaluminate precipitates from the melt, after the dissolution of corundum.
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5. 3. 2. Matrix reaction
In order to further transfer the kinetics results to the corrosion of refractory castables with biomass slag, ex situ quantification is carried out on dense matrix samples in contact with the ternary slag CSK. The results for Matrix A covered with 0.9 g of slag are shown in Fig. 5-7. The standard deviation is only determined on a few samples and can be estimated at 1 wt.-%
Figure 5-7. Ex situ quantification of the corrosion kinetics of Matrix A with 0.9 g of slag CSK at 1400 °C
The decrease of the corundum content is fast: 78 wt.-% corundum remains after 1 h to reach a stable content of 66 wt.-% after 6 h, which keeps constant after a longer exposition. The β- alumina formation is also rapid as 11 wt.-% are formed after 1h. However this phase is not stable and is further dissolved until a content of 2 wt.-% is reached after 16 h. The increase of the amorphous content reaches a maximal value of 10 wt.-% after 6 h and decreases down to 7 wt.-% after 16 h as calcium hexaaluminate is precipitated. The formation of calcium hexaaluminate is slow: No traces of this phase is present until 6 h of reaction, and its appearance occurs simultaneously as the β-alumina content starts to decrease.
As comparison, the results of the reaction between CSK slag and the matrix CA6 are shown in Fig. 5-8:
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Figure 5-8. Ex situ quantification of the corrosion kinetics of Matrix CA6 with 0.9 g of slag CSK at 1400 °C
The effect of decrease and subsequent increase of the calcium hexaaluminate amount is once again observed, with a content of 92 wt.-% after 30 min and 93 wt.-% after 2 h. This behaviour is in agreement with the liquid content, increasing until 6 wt.-% after 30 min and decreasing again down to 4 wt.-% after 2 h. The further decrease of the liquid amount down to 2 wt.-% after 24 h is linked with the precipitation of 6 wt.-% of calcium dialuminate. Just as for the reaction with Matrix A, the formation of potassium β-alumina is observed in lower extent with 3 wt.-% after 2 h, and diminishes down to 1 wt.-% after 24 h.
The global composition at the thermodynamic equilibrium is calculated with FactSage for a composition of 6 g matrix and 0.9 g slag at 1400 °C. The stable phases can be identified as well on Fig 4-3 a and 4-5 a with a slag amount of 13 wt.-%
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Table 5-4. Phase composition at the thermodynamic equilibrium for 6 g matrix and 0.9 g CSK slag at 1400 °C
[wt.-%] A CA6 CA2 KAS4 KAS2 Liq
Matrix CA6 65.2 21.7 13.0 Matrix A 27.5 62.3 7.2 4.3
The remaining phases after 24 h are consistent with the thermodynamic predictions: The calcium hexaaluminate bond phase is more stable than the alumina bond phase (Tab. 5-4). The predicted values are not reached experimentally (Tab. 5-4 and Fig. 5-7 and 5-8), and this calculation is only performed for a qualitative estimation of the stable phases. The predicted phases of the binary system CaO-Al2O3 are detected experimentally. However the calculated potassium containing solid phases (KAS4 and KAS2) does not corroborate the experimental results.
The rates of transformation of alumina and calcium hexaaluminate after 30 min of contact with the slag CSK are determined for the different slag amounts (Fig. 5-9). The rates are expressed as the rate of change of the main phase dc/dt, with c the concentration of the main phase in wt.- % and t the time in h.
Figure 5-9. Rates of transformation for the main phase of Matrix A (Al2O3) and Matrix CA6 (CaO·6Al2O3) after 30 minutes at 1400 °C in contact with CSK slag
Both transformations are fast, but the transformation of alumina exhibits higher rates than the transformation of calcium hexaaluminate. While both rates are similar with 0.3 g of slag, the
74 differences increase with increasing slag amount. At 0.9 g of slag, the matrix A transforms with a rate of 29 dwt.-%/dt, against 13 dwt.-%/dt for the matrix CA6.
5. 3. 3. Phases identification in the microstructure
The observation of the reaction products focuses on the reaction test at 1400 °C for 24 h with 0.9 g of slag CSK. The quantification of the phases obtained with the degree of crystallinity and Rietveld refinement are reminded in Tab. 5-5.
Table 5-5. Quantitative XRD of the matrix corroded with 0.9g CSK after 24h at 1400°C
[wt.-%] A ßA CA6 CA2 C2AS Liq
Matrix CA6 89 6 3 2 Matrix A 67 2 24 7
After 24 h of corrosion, both samples are containing four phases composed of four chemical compounds (Al2O3; CaO; SiO2 and K2O). According to Gibbs' rule the present system is invariant, meaning the composition of the liquid phase is constant.
A second corroded sample is kept to be embedded in resin and cut perpendicular to the reaction front. The sample is polished and coated with carbon for SEM observations (Fig. 5-10 and 5- 11)
Figure 5-10. Micrograph of Matrix CA6 after 24h at 1400°C in contact with 0.9g CSK slag and quenching in water
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Table 5-6. EDX measurements from Fig. 5-10
[wt.-%] Na2O Al2O3 SiO2 K2O CaO 1 78.3 21.7 2 37.0 22.1 40.9 3 91.2 8.8 4 0.5 72.1 11.1 1.2 15.1
Four different phases can be observed in the microstructure (Fig. 5-10):
▪ Polycrystalline calcium hexaaluminate (Tab. 5-6 EDX#3). The grain boundaries are visible and well distributed among this phase. ▪ The intergranular liquid phase can be assumed from Tab. 5-6 EDX#4 where the heterogeneous composition with the strong impurity content indicates the presence of an amorphous phase, supposedly at the grain boundaries of the calcium hexaaluminate being dissolved. The grain boundary liquid should be saturated with calcium hexaaluminate. ▪ The calcium dialuminate (Tab. 5-6 EDX#1). Although the kinetics results attest the precipitation of the calcium dialuminate (Fig. 5-8), the one observed on Fig. 5-10 is more likely to come from the calcium aluminate cement with the presence of intragranular pores, which is unusual for precipitates. ▪ Gehlenite (Tab. 5-6 EDX#2) with uniform aspect, resembling an amorphous phase.
Those four phases correspond to the XRD measurements, showing the coexistence of calcium hexaaluminate, calcium dialuminate, Gehlenite and slag (Tab. 5-4).
A high magnification of the corroded alumina matrix is shown in Fig. 5-11
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Figure 5-11. Micrograph of Matrix A after 24h at 1400°C in contact with 0.9g CSK slag and quenching in water
Table 5-7. EDX measurements from Fig. 5-11
[wt.-%] Na2O Al2O3 SiO2 K2O CaO 1 1.0 42.1 30.0 8.1 18.9 2 90.9 0.2 8.9
The tabular alumina grains are observed on the Fig. 5-11 with intragranular porosity and a isometric shape. The morphology of the calcium hexaaluminate (Tab. 5-7 EDX#2) is rather uniaxial after growing from the melt. The chemistry attests the presence of potassium from the slag inside the calcium hexaaluminate reaction product. The Tab. 5-7 EDX#1 shows a various composition with alkaline and glass forming silica, which could correspond to a liquid phase. The EDX measurement has a sensitivity of 10 µm and it can be noticed from the micrograph that a solid phase is also contributing to the first EDX measurement (Tab. 5-7 EDX#1). Taking into account the high amount of potassium (Tab. 5-6 EDX#1) and the correlation with the β- alumina detected with XRD (Tab. 5-4), the small crystal (1x10 µm2) could be potassium β- alumina being dissolved by the slag.
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In both matrix, it is worth mentioning that the ratio K2O:SiO2 is much lower than the one from the bulk slag (1:2), and it could be assumed that some potassium is released in the gas phase after the decomposition of the β-alumina.
5. 4. Conclusion
The kinetics of the matrix reactions with slag involve a complex aspects of corrosion and have been addressed experimentally. For both alumina and calcium hexaaluminate matrix, the dissolution is fast, followed by a slow precipitation of an intermediate compound, calcium hexaaluminate in the case of alumina and calcium dialuminate in the case of calcium hexaaluminate (Fig. 5-7 and 5-8). The formation of those products is in agreement with thermodynamic predictions (Tab. 5-4). The trials allow to confirm the dissolution/precipitation as reaction mechanism, since the formation of a transitional liquid phase is detected, both with XRD (Fig. 5-7 and 5-8) and with SEM (Fig. 5-10 and 5-11). The focus set on the direct dissolution of Powder CA6 in the binary slag (Fig. 5-3) highlights possible effect of the alumina impurities in calcium hexaaluminate product: Their dissolution is fast, and leads to a secondary formation of calcium hexaaluminate in the microstructure. This assumption was already established in the previous chapter based on microstructural observations (Fig. 4-16) and is yet attested by ex situ and in situ trials on powder (Fig. 5-3 and 5-4), as well as on the matrix (Fig. 5-8). For the powder experiments, dissolution kinetics are faster than for the matrix due to the greater contact area. In this case the limits of ex situ measurements are encountered, since the dissolution takes place in the first minutes of contact (Fig. 5-3). For the matrix experiments, in situ XRD measurement cannot be considered and the potential crystallisation of the liquid phase upon cooling has to be taken into account.
The main kinetic effect determined experimentally is the rapid formation of the β-alumina due to the potassium contained in the slag (Fig. 5-7 and 5-8). On contact with the molten slag, the refractory first reacts with the most mobile ion from the melt. Although this formation in contradiction with thermodynamic predictions was already assigned to a higher mobility of the potassium in the previous chapter, it is yet attested that this β-alumina phase is not stable. The further dissolution of the β-alumina is observed once all the potassium from the slag has been consumed (Fig. 5-11). The low amounts of potassium measured with EDX after 24h (Tab. 5-7 and 5-6) point out the possible volatilization of the potassium as the β-alumina is decomposed.
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This could also be deducted from the XRD results (Fig. 5-7 and 5-8), where the amounts of β- alumina formed are lower than the expected quantities (Fig. 4-12-a).
The kinetics study highlights a better performance of calcium hexaaluminate matrix in comparison with the alumina matrix. The XRD experiments show a better stability of the calcium hexaaluminate, with 89wt.-% remaining phase after 24h, against 67wt.-% for the alumina matrix (Tab. 5-5). Those results are in agreement with thermodynamic calculations predicting lower extent of reaction for the calcium hexaaluminate than for the alumina with same slag amount (Tab. 5-4). Moreover the liquid content is lower in the calcium hexaaluminate matrix (Fig. 5-8), slowing the transport of the reactants toward the interface. Many conclusions can be drawn to explain the lower dissolution of the CA6 matrix:
▪ The lower liquid amount of calcium hexaaluminate matrix could be explained by the
better stability of the precipitated Gehlenite in the matrix CA6 while the Gehlenite formed in the alumina sample is further dissolved after 24h (Fig. 5-7 and 5-10). ▪ The enhanced liquid formation in the alumina matrix could also result from the enhanced β-alumina formation. Indeed the reaction of alumina with potassium forms β- alumina as the only reaction product, which is unstable at those conditions. This phase has to be further dissolved and contributes to the liquid formation. In contrary, the reaction of calcium hexaaluminate with potassium forms not only β-alumina, but also calcium dialuminate (Fig. 4-12), which is stable and can exist as a solid phase without being dissolved.
▪ It can also be assumed that the poorer formation of β-alumina in the matrix CA6 results from an enhanced amount of potassium in the gas phase. Indeed, the potassium could
hardly be observed in the post-mortem microstructure of the matrix CA6 (Fig. 5-10), explaining the lower solubility (Fig. 4-1). A higher volatilization not only depletes the melt in alkali but also lowers the slag amount, participating as well to the decreased dissolution.
Beside the lower extent of reaction for the calcium hexaaluminate matrix, the rates of transformation are also higher for alumina, especially with increasing slag amount (Fig. 5-9). However, the poorer performance of the alumina regarding the resistance to dissolution might result from the more permeable microstructure (Fig. 3-12) and the influence of the porous network on the corrosion extent has to be addressed in the next chapter.
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Chapter 6: Microstructural considerations
After determining the corrosion reactions and their kinetics, this chapter considers the different microstructural features involved in the corrosion process.
6. 1. Theory
6. 1. 1. Microstructural properties influencing the corrosion resistance
The characteristics of the capillary system of a refractory product is a determining factor in the corrosion resistance, as it can hinder the slag infiltration, or on the contrary, increase the surface reactions [KAI71;KAL72]. Therefore, refractory producers tend to reduce the surface area potentially in contact with slag to lower the rate of dissolution. For instance the reduced corrosive wear can be achieved through the selection of fused aggregate over sintered aggregate, raw material with larger grain sizes and reduced open porosity in the matrix.
Figure 6-1. Schematic representation of the porous network according to [SAL07]
The Fig. 6-1 shows the different types of porosity present in the refractory microstructure. The closed pores do not contribute to the slag penetration, while the interconnected open porosity influences strongly the corrosion resistance of the refractory. Therefore the surface area of the porosity and the connection between the pores are important characteristics to consider by determining respectively the pore size distribution and the permeability prior to the corrosion test [SAL07;MCC04].
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6. 1. 2. Microstructure development of calcium hexaaluminate
Calcium hexaaluminate is formed by solid-state reaction after mixing the particles of reactants. During the sintering of calcium aluminate cement bonded refractories, the formation occurs according to the reaction (13):
( 13 ) 퐶푎퐴푙4푂7 + 4퐴푙2푂3 → 퐶푎퐴푙12푂19
The reaction begins at around 1000-1100°C with the interdiffusion of the involved ions and requires high firing temperature to be fully completed [SIN02;TUL03;IAN08;PIE15;DON16;SAL16]. 2+ The CA6 growth takes place through Ca transport anisotropies along the basal plan leading to an enhanced texture. The resulting platelet morphology generated from a two dimensional nuclei is responsible for an enhanced intragranular porosity, by analogy with abnormal grain growth [JON84].
In the case where alumina and calcium carbonate are used as precursors, the formation of low melting phase like mayenite (C12A7) can yield a liquid phase in the microstructure, which is also contributing to the oriented growth of platelets. The space required to grow platelets is relevant as well, confirming the influence of processing conditions, like the green packing density or the homogeneous distribution of the reactants, to allow an equiaxed growth of calcium hexaaluminate [AN96;DOM98;DOM01;VIS05;NAG06;CHA09;DEL12;LIU14;LIU15;SAL16]. It is established as well that an increased sintering temperature favorizes the further growth of platelet into equiaxed grains, supposedly through welding with parallel oriented neighbors.
The CA6 tabular elongated shape is a major asset in calcium aluminate cement bonded refractory castable since it behaves as crack deflectors increasing the product toughness and thermal shock resistance. The plate-like grain morphology is even used to toughen alumina ceramics [ASM98]. Concerning the corrosion properties, the anisotropic crystal structure might exhibit different dissolution rates depending on the crystallographic plane [MCC04].
6. 1. 3. Expansive reactions
A reaction product is expansive if it requires more volume than the reactant. In this case, the volume expansion induces strains in the microstructure leading to the cracking and rupture of the material, known as structural spalling.
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Table 6-1. Phase densities according to [BUC05]
3 [g/cm ] A CA6 ßA Density 3.99 3.79 3.37
Tab. 6-1 shows the density of the different phases encountered in this work. Based on those values, it can be predicted that the reaction from alumina to calcium hexaaluminate or β- alumina leads to a volume expansion though the formation of lower density phases. The theoretical volume increase of the reaction from α-alumina to β-alumina is determined in [BUC05]:
11Al2O3 (s) + K2O (liq) → K2Al22O34 (s)
102 g/mol 1200 g/mol
3.99 g/cm3 3.37 g/cm3
3 3 25.6 cm /mol 356.1 cm /mol
푉퐾 퐴푙 푂 − 11푉퐴푙 푂 ∆푉 = 2 22 34 2 3 = +26.5% 11푉퐴푙2푂3
By reacting into β-alumina, α-alumina expands 26.5 % of its initial volume. This volume increase is calculated at the crystallographic scale. In the microstructure, this expansion can take place in the open porosity or it can, on the contrary, generate compressive strains among the matrix.
6. 2. Method
6. 2. 1. Characterisation of the porous network
The characterisation of the matrix capillaries is based on different standard techniques. The open porosity is measured according to the European standard for refractory samples 993-1. The pore size is determined with mercury porosimetry with the Autopore IV 9500 Series from Micromeritics. Samples are grounded and sieved <63µm to determine the density of the refractory by means of helium pycnometer using the Ultrapyc 1200e from Quantachrome.
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6. 2. 2. Corrosion tests
To investigate the microstructural changes generated by corrosion with model slag, some of the samples used in Chapter 5 are further examined in the present chapter. To address the influence of the porous network on the dissolution rate, additional samples with enhanced open porosity are produced by adding water to the dry composition during the casting of the matrix samples. The corrosion setup consists of a cut slice of the refractory matrix covered with a pressed pellet of 0.9 g slag CSK. The testing at high temperature is performed by introducing the corrosion couple in a resistive furnace at 1400°C and quenching in water down to room temperature after the targeted reaction time.
For the comparison with biomass slag, cylindrical samples of matrix (H10mmxD50mm) are covered with a pressed tabled of 5 g of wood ash. The samples are heated in a resistive furnace with 2 K/min and kept 2h and 24h at 1550°C before cooling naturally down to room temperature.
6. 2. 3. XRD measurements
To consider the influence of the microstructure, space resolved XRD is performed after sectioning the corroded samples perpendicular to the refractory/slag interface. The sample is placed on the XYZ stage of the diffractometer. A snout is mounted on the incident X-Ray beam with a diameter of 1 mm, and three XRD measurements are performed at a given distance from the interface before adding the recorded signals to obtain a diffractogram corresponding to a surface of 1x3mm2. The scan is performed from 5 °-90 ° [2θ] with step of 0.02 ° and 384 s/step.
6. 3. Results and discussion
6. 3. 1. Resistance to structural spalling
After reacting at 1400°C with CSK slag for 24h, the matrix samples revealed strong differences in their geometry (Fig. 6-2 and 6-3). The samples are kept massive for space resolved XRD using a XYZ table as sample holder.
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Figure 6-2. Space resolved XRD of Matrix A after 24h in contact with CSK slag at 1400°C and quenching in water
At the surface of the alumina matrix, which was covered with slag, alumina is present with high amounts of calcium hexaaluminate, potassium β-alumina and amorphous phase. The measurement performed at a distance of 5 mm from the interface shows an attenuated presence of the reaction products. The gradient in the mineralogical composition between the interface and the core of the sample is strongly marked. The difference of volume between the reaction products and the alumina leads to the curved shape of the sample after cooling and the cracking observed on the Fig. 6-2.
Figure 6-3. Space resolved XRD of Matrix CA6 after 24h in contact with CSK slag at 1400°C and quenching in water
While the alumina sample shows a strong deformation after corrosion (Fig. 6-2), the calcium hexaaluminate matrix remains dimensionally stable (Fig. 6-3). The XRD measurements show
84 as main phase calcium hexaaluminate with only traces of Gehlenite and Grossite, confirming the low reactivity of the calcium hexaaluminate matrix observed in the previous chapter. The silica containing phase is only present at the interface. Since the extent of reaction at the interface is low, only a slight difference is observed in the XRD measurements of the surface and of the core of the sample. At a distance of 5 mm from the interface, calcium hexaaluminate is present with traces of calcium dialuminate. The density variations are measured on the grounded samples with helium pycnometer (Tab. 6-2).
Table 6-2. Density measurements of the matrix samples
Before After [g.cm-3] corrosion with CSK 1400°C/24h Matrix A 4.0 3.4 Matrix CA6 3.7 3.4
The starting density for the Matrix A is higher, with a value of 4.0 g.cm-3 than for the Matrix -3 CA6, with a value of 3.7 g.cm , in agreement with the theoretical densities (Tab. 6-1). After corrosion, both matrix samples exhibit the same density of 3.4 g.cm-3. The density measured after corrosion is the sum of all the single contributions of the different phases coexisting the sample. The reaction with CSK slag generates low density products, which require higher volume than the original phase. The density gradient is lower in the case of the Matrix CA6, and the reaction products can expand in the open porosity without macroscopic volume expansion.
The observations after corrosion with the model slag are compared with the resulting sample geometry after corrosion with wood ash (Fig. 6-4).
Figure 6-4. Matrix samples after corrosion with wood ash at 1550°C a) Matrix A after 2h; b) Matrix A after 24h; c) Matrix CA6 after 24h
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In the case of the alumina matrix, a strong structural spalling is visible after corrosion with wood ash, while calcium hexaaluminate seems to react without volume variations, hence without damaging the original microstructure of the refractory product.
The volume variations are calculated based on geometry measurements after the natural cooling of the corroded samples (Fig. 6-5).
Figure 6-5. Post-mortem volume expansion after corrosion with wood ash at 1550°C
After corrosion with wood ash, the calcium hexaaluminate sample shows a low expansion of 2 Vol.-% which remains constant after different testing durations. On the contrary, the alumina matrix exhibits a strong volume expansion of 16 Vol.-% after 2h at 1550°C, which increases up to 36 Vol.-% after 24h.
6. 3. 2. Influence of the porosity
In order to test the influence of the matrix porosity on the corrosion behaviour, different matrix are cast with different amount of water. The open porosity is measured with water according to the European standard 993-1 and with mercury porosimetry (Fig. 6-6)
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Figure 6-6. Open porosity of Matrix CA6 after sintering at 1700°C/6h according to the water amount added for the casting
The open porosity measured with water yields higher porosity values, since they also include the macroporosity, which is in minority in the refractory matrix. The intrusive mercury porosimetry has access only to the micropores, explaining the slightly lower values. The open porosity increases linearly with the amount of water added during the mixing of the matrix components. No significant variations of the pore diameter is measured with the mercury porosimetry, and the values are shown in Tab. 6-3.
Table 6-3. Average pore diameter of matrix CA6 after sintering at 1700°C/6h measured with mercury porosimetry
Water amount [wt.-%] 10 11 12 13 14 Average pore diameter [µm] 0.6 1.0 1.3 1.3 1.8
The different samples are corroded with 0.9 g of slag CSK as described in the previous chapter to determine the rate of dissolution of calcium hexaaluminate after 2 h and after 24 h. The values are compared in Fig. 6-7 .
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Figure 6-7. Rate of dissolution of Matrix CA6 in CSK slag at 1400°C according to the open porosity
After 2 h in contact with CSK slag, the rate of dissolution is strongly influenced by the open porosity of the matrix. The original Matrix CA6 compared previously with the alumina matrix exhibits an open porosity of 21.5 % and the dissolution rate after 2 h is evaluated at - 2.7 dwt.- %/dt. The matrix with 28 % open porosity, shows after 2 h a dissolution rate of - 5.9 dwt.-%/dt, with a linear increase of the dissolution speed according to the open porosity. After 24h, the dissolution rate is close to zero, and the open porosity does not have any influence anymore. The open porosity influences the rate of dissolution at the beginning of the corrosion process, with lower dissolution rates for lower open porosity. As the reaction advances, the dissolution rates decrease strongly declining the impact of the open porosity. It is consistent with the literature, confirming that the rate of dissolution increases linearly with the percentage of apparent porosity [SCH04].
6. 3. 3. Morphology of calcium hexaaluminate
As mentioned in the introduction of this chapter, the calcium hexaaluminate phase can exhibit different types of morphology. The Rietveld refinement performed in the previous chapter to determine the amount of reaction products can also determine qualitatively the shape of the calcium hexaaluminate crystals. A tool is proposed in the TOPAS interface to take into account the preferred orientation of one phase in the quantification of the crystalline compounds. Indeed the quantification can be distorted if the texture is not considered for the refinement [DIC69]. A coefficient is then attributed to the (0 0 1) direction planes, indicating the degree of texture. A
88 coefficient close to one refers to equant crystals, while a coefficient smaller than one indicates the presence of platelet (Fig. 6-8).
Figure 6-8. Coefficient of preferred orientation for the calcium hexaaluminate phase obtained from the Rietveld refinement performed in the previous chapter after corrosion with CSK slag at 1400°C.
The coefficient indicates two types of morphology for the calcium hexaaluminate phase, according to the two types of growth:
▪ The sintered calcium hexaaluminate obtained either as sintered aggregate or from the sintering between calcium aluminate cement and reactive alumina. Those calcium
hexaaluminate crystals present in the Matrix CA6 are hardly oriented and can be described as equiaxial with a coefficient close to one (Fig. 6-8). ▪ The precipitated calcium hexaaluminate obtained from the dissolution of alumina matrix in the slag CSK. Those crystals are textured and the coefficient lower than one indicates a platelet morphology (Fig. 6-8). The growth of those platelet in the intergranular melt is observed with SEM (Fig. 6-9)
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Figure 6-9. Micrographs of calcium hexaaluminate formed in Matrix A after corrosion with CSK slag at 1400°C/150h and quenching in water
Table 6-4. EDX measurements from Fig. 6-8.
[wt.-%] Na2O Al2O3 SiO2 K2O CaO 1 0.5 75.2 11.3 0.9 12.2 2 1.9 38.3 34.7 3.5 21.7 3 1.6 54.5 24.3 2.5 17.2 4 91.4 8.6
The morphology of the calcium hexaaluminate as reaction product is visible on the Fig. 6-8. The platelets are precipitating from an intergranular liquid phase in the vicinity of the tabular alumina being dissolved. The EDX#3 reveals a chemical composition close to the stoichiometry of calcium hexaaluminate, while the EDX #1 and #2 is comprising both calcium hexaaluminate and the liquid phase composed of the slag elements, depleted in potassium and enriched in alumina and sodium (Tab. 6-4).
6. 4. Conclusion
Different amounts of water are added during the mixing to the dry composition of the calcium hexaaluminate matrix in order to determine the influence of capillaries network on the corrosion. The resulting open porosity after sintering of the matrix is correlated to the additional water amount (Fig. 6-6). After testing the corrosion with CSK slag, it is observed that the enhanced open porosity increases linearly the rate of dissolution (Fig. 6-7). The objective of this trial is to determine the contribution of the matrix microstructure to the dissolution rates values, in order to assess if the better performance of Matrix CA6 observed in the previous
90 chapter could be attributed to the lower permeability and smaller pore sizes of the microstructure. However the addition of water to increase the open porosity does not lead to any significant increase of the pore size (Tab. 6-3). The comparison with the microstructure of Matrix A is not achieved and the influence of the microstructural features in the high corrosion resistance of Matrix CA6 remains undefined.
Two types of morphology could be observed for the calcium hexaaluminate phase. The sintered calcium hexaaluminate exhibits an equiaxed crystal shape without preferred orientation, while the precipitated calcium hexaaluminate is textured and exhibits a platelet morphology
(Fig. 6- 8). The equiaxed grains obtained in Matrix CA6 after sintering are not representative of the morphology usually obtained in calcium aluminate cement bonded castable. According to literature [DOM01], this could be attributed to the enhanced sintering temperature of 1700°C, leading to the growth of platelets into equiaxed grains.
The plate-like morphology is observed in the alumina matrix, after undergoing a strong volume expansion. The structural spalling is not attributed to the alkali reaction from the potassium contained in the slag, but rather to the reaction with the calcia (Fig. 6-2). The expansive formation of calcium hexaaluminate is due to the lower density of the phase formed. The volume increase from corundum to calcium hexaaluminate can be calculated:
6Al2O3 (s) + CaO (liq) → CaAl12O19 (s)
102 g/mol 668 g/mol
3.99 g/cm3 3.79 g/cm3
3 3 25.6 cm /mol 176.2 cm /mol
푉퐶푎퐴푙 푂 − 6푉퐴푙 푂 ∆푉 = 12 19 2 3 = +14.7% 6푉퐴푙2푂3
Though the theoretical volume increase is lower than the one caused by alkali reaction, the volume expansion for the transformation of alpha alumina into calcium hexaaluminate is considerable. Since calcium hexaaluminate exhibits a low density and a low reactivity, no structural spalling is observed after corrosion in Matrix CA6, neither with the model slag, nor with the wood ash.
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Chapter 7: Conclusion
7. 1. Advantages of calcium hexaaluminate
The resistance to structural spalling of calcium hexaaluminate is confirmed in this study. Although alkali bursting is not hold accountable for the expansive reaction, corrosion tests on alumina matrix samples result in strong volume expansions. The XRD measurements allow identifying the responsible reaction as the calcium hexaaluminate formation upon the dissolution of alumina. The precipitation of calcium hexaaluminate is not only linked with an increase of the specific volume of the solid phase, but the growth is also strongly anisotropic, generating compressive stresses in the microstructure, leading to enhanced cracking. Those effects are not observed in the calcium hexaaluminate matrix microstructure: The geometry of the sample is stable in contact with alkali containing model slag and wood slag, which is a strong asset for its potential application.
The enhanced resistance to structural spalling of calcium hexaaluminate is attributed to a better chemical stability. After being in contact with the liquid slag, the calcium hexaaluminate matrix shows low extents of reaction. Thermodynamic calculations support the hypothesis of the lower reactivity: Upon reaction with the same amount of CaO-SiO2-K2O slag, the remaining quantity of calcium hexaaluminate at the equilibrium is higher than the remaining quantity of alumina. This confirms that calcium hexaaluminate is less prone to react with such slag composition. Moreover the rate of dissolution determined experimentally shows slower kinetics for calcium hexaaluminate than for alumina. This better chemical resistance would extend the service life of the refractory lining in biomass incinerators.
The reduced reactions for calcium hexaaluminate samples can also be attributed to the lower permeability of the microstructure. The porous network of both materials are characterized prior to the corrosion test. Alumina samples show higher permeability values, as well as higher average pore diameter, which offer more surface to react and show indeed a greater extent of reaction with CaO-SiO2-K2O slags. The sol-gel bonded alumina castable has restrictive rheological properties which might impair the microstructural properties of the monolith On the contrary, the reduced pore size is attributed to the rheological properties of the cement containing castable, which is well deflocculated by using a commercial polymeric dispersant. Lower permeability and pore size hinder slag infiltration and enhance the resistance against corrosive wear.
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7. 2. Limitations of the experimental work
The corrosion of refractory materials involves complex systems, and the influence of the single parameters is difficult to sort out. The main issue of the experimental work is to vary the chemical composition of the refractory product, either powder, matrix or castable, without changing its physical properties to enable a clear statement about the properties of the mineral phase. Variations in the chemistry implies the implementation of different raw materials. In turn, varying raw materials generate discrepancies in the physical properties, which also influence strongly the resulting corrosion behavior of the product. For instance in the case of the matrix samples, the sintering behavior of the calcium hexaaluminate and alumina raw materials is different. It results in a more porous microstructure with lower interconnected capillaries in the first case, while the alumina microstructure have lower open porosity but higher permeability values. The changes in the microstructure are difficult to master in order to isolate the intrinsic properties of the mineral phases.
For a proper interpretation of the post-mortem analyses, the evolution of the chemical composition at the refractory/slag interface has to be well known. The selection of bulk slag and refractory chemistry is not sufficient to control the interfacial reactions, since the different mobility of the ions from the slag can enrich the composition at the interface with the more mobile species. The comparison between the experimental results and the calculated thermodynamic equilibrium is hindered since the local composition at the interface cannot be deducted from the bulk composition. For instance, while the thermodynamic laws have not predicted the formation of potassium β-alumina during the corrosion with CaO-SiO2-K2O slag, the formation of this phase is observed experimentally due to the higher mobility of the potassium. Though the influence of the reactants transport is known, its quantification is not trivial and the determination at a given time of the interface composition can only be achieved empirically in such a complex system.
A well known impediment in the study of high temperature corrosion is the ex situ nature of most of the analyses, especially the phase determination on massive corroded samples. The ex situ determination of the reaction extent has two drawbacks. First, the reactions with fast kinetics cannot be tackled experimentally with ex situ trials, since the time used for heating and cooling is comparable to the time needed for the reaction to take place. Second, the effect of the cooling on the reactions is not controlled. In this case, partial crystallization or cracking of the samples have to be taken into account for the interpretation. To encounter those effects, in situ XRD measurements can be performed on powder reactants using a high temperature
93 chamber. The time required for the acquisition can be shortened by using a curved detector to enable timed resolved determination of the phase composition [DOM15;POI17]. The comparison of the reactions on powders and massive samples can be a solution to enable on one hand in situ measurement, and still take into account the physical properties of the microstructure.
7. 3. Outlook
This study provides a support to understand the corrosion with molten ashes, the main solicitation for refractory linings in combustion plants. The difference of corrosion behaviour between calcium hexaaluminate and alumina is addressed to further understand their performances in contact with alkali containing slag regarding their thermodynamic and microstructural features. Resistance to alkali vapors represents the second main solicitation in biomass incinerators and the resistance to alkali of calcium hexaaluminate is further tackled in a parallel PhD work carried out at the CEMHTI laboratory in Orléans, in the framework of FIRE C2. The knowledge gathered from this complementary research characterizes both the behaviour of alumina rich materials in the presence of alkali vapors and their resistance against the low melting phase generated from their condensation at the refractory interface.
Laboratory tests are preliminary to enable the first statements about the corrosion mechanisms of calcium hexaaluminate with a defined slag composition. In the mean time, a calcium hexaaluminate castable has been developed and samples have been used to line the post combustion chamber of an household waste incinerator in Houston, United States of America, also as a contribution to FIRE C2. This trial will take place until 2019 and the performance of calcium hexaaluminate will be compared with other materials usually used in the incineration of organic feedstocks. The present work should contribute to the interpretation of the future industrial results through the comparison of the post-mortem microstructures and mineralogy. The alternative use of calcium hexaaluminate as refractory material for new energy applications will then be fully defined with its advantages and limitations at both laboratory and industrial scales.
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LIST OF FIGURES
Figure 1-1. Dissolution in slag of the protective layer surrounding alumina grain ...... 2
Figure 1-2. Articulation of the different aspects of corrosion ...... 3
Figure 2-1. Average chemical composition and its standard deviation ...... 6
Figure 2-2. Schematic representation of the calcium hexaaluminate crystal structure and its comparison with β- alumina from [CER98] ...... 12
Figure 3-1. Ternary phase diagram of SiO2-Al2O3-CaO from [HAC15] ...... 17
Figure 3-2. Binary phase diagram of CaO-Al2O3 (wt.-%) from [LEA56] ...... 18
Figure 3-3. Binary phase diagram of CaO-SiO2 from [SLA95] ...... 19
Figure 3-4. Ternary phase diagram of CaO-SiO2-K2O (wt.-%) from [CHE16] ...... 20
Figure 3-5. Pseudo-binary phase diagram of K2O·Al2O3-Al2O3 from [SLA95] ...... 20
Figure 3-6. Schematic representation of network of tetrahedra ...... 21
Figure 3-7. Change of bridging oxygen from OAl,Ca to OAl,K from [ZHA12] ...... 23
Figure 3-8. X-Ray Diffraction patterns of Powder CA6 ...... 24
Figure 3-9. X-Ray Diffraction patterns of Matrix A ...... 27
Figure 3-10. X-Ray Diffraction patterns of Matrix CA6 ...... 28
Figure 3-11. Pore size distribution after sintering at 1700°C/6h ...... 29
Figure 3-12. Permeability of alumina and calcium hexaaluminate matrix ...... 30
Figure 3-13. Micrograph of Castable A after sintering at 1500°C/6h ...... 31
Figure 3-14. Micrograph of Castable CA6 after sintering at 1500°C/6h ...... 32
Figure 3-15. Average composition of wood ash from [VAS15] ...... 33
Figure 3-16. Recording of the melting behaviour of wood ash on alumina substrate according to European standard 51730...... 34
Figure 3-17. Micrographs of the amorphous slag CSK ...... 34
Figure 3-18. Micrograph of the amorphous slag CS ...... 35
Figure 3-19. Micrograph of the hot stage microscopy sample ...... 36
Figure 4-1. Calculated solubility limits of Al2O3 and CaO·6Al2O3 ...... 44
Figure 4-2. Phase composition at the thermodynamic equilibrium ...... 45
Figure 4-3. Phase composition at the thermodynamic equilibrium ...... 45
Figure 4-4. Phase composition at the thermodynamic equilibrium ...... 46
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Figure 4-5. Phase composition at the thermodynamic equilibrium during the dissolution of CaO·6Al2O3 in slag CSK ...... 47
Figure 4-6. Castable CA6 a) after sintering; ...... 48
Figure 4-7. Micrograph of Castable CA6 after corrosion ...... 49
Figure 4-8. Micrograph of Castable A ...... 51
Figure 4-9. Micrograph of Powder CA6 pellet after corrosion ...... 52
Figure 4-10. Micrograph of Castable CA6 ...... 53
Figure 4-11. Micrograph of Castable CA6 ...... 54
Figure 4-12. Phase composition at the thermodynamic equilibrium ...... 55
Figure 4-13. Phase composition at the thermodynamic equilibrium ...... 55
Figure 4-14. Phase composition at the thermodynamic equilibrium ...... 56
Figure 4-15. Micrograph of Castable CA6 ...... 56
Figure 4-16. Micrograph of Castable CA6 ...... 57
Figure 4-17. Schematic comparison of the monomineral layers formation ...... 58
Figure 4-18. Schematic comparison of the corrosion mechanisms ...... 59
Figure 5-1. Energy barrier diagram ...... 61
Figure 5-2. Comparison of the different methods ...... 67
Figure 5-3. Ex situ quantification of the amorphous content ...... 68
Figure 5-4. In situ recording of the intensity ...... 69
Figure 5-5. Diffractograms at room temperature of Powder CA6 and slag CS ...... 70
Figure 5-6. Scans with Position Sensitive Dectector (PSD) fixed at 32°-34.927° ...... 71
Figure 5-7. Ex situ quantification of the corrosion kinetics ...... 72
Figure 5-8. Ex situ quantification of the corrosion kinetics ...... 73
Figure 5-9. Rates of transformation for the main phase ...... 74
Figure 5-10. Micrograph of Matrix CA6 after 24h at 1400°C ...... 75
Figure 5-11. Micrograph of Matrix A after 24h at 1400°C ...... 77
Figure 6-1. Schematic representation of the porous network according to [SAL07] ...... 80
Figure 6-2. Space resolved XRD of Matrix A after 24h ...... 84
Figure 6-3. Space resolved XRD of Matrix CA6 after 24h ...... 84
Figure 6-4. Matrix samples after corrosion with wood ash at 1550°C ...... 85
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Figure 6-5. Post-mortem volume expansion after corrosion with wood ash at 1550°C ...... 86
Figure 6-6. Open porosity of Matrix CA6 after sintering at 1700°C/6h ...... 87
Figure 6-7. Rate of dissolution of Matrix CA6 in CSK slag at 1400°C ...... 88
Figure 6-8. Coefficient of preferred orientation for the calcium hexaaluminate phase ...... 89
Figure 6-9. Micrographs of calcium hexaaluminate formed in Matrix A ...... 90
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LIST OF TABLES
Table 3-1. ASTM classification of high alumina refractory castables [ROU01;AST00] ...... 15
Table 3-2. Mineral phases encountered in this work ...... 16
Table 3-3. Viscosities of Al2O3-SiO2-CaO slag from [GHO10] and FactSage calculations of liquidus temperature ...... 22
Table 3-4. Chemical composition of refractory powder measured with X-Ray Fluorescence ...... 24
Table 3-5. Grain sizes of Powder CA6 ...... 25
Table 3-6. Formulations of the matrix samples ...... 26
Table 3-7. Formulations of Low Cement Castable (LCC) ...... 31
Table 3-8. EDX measurements from Fig. 3-14 ...... 32
Table 3-9. EDX measurements from Fig. 3-17 ...... 35
Table 3-10. EDX measurement from Fig. 3-18 ...... 36
Table 3-11. EDX measurement from Fig.3-18 ...... 36
Table 3-12. Overview of the corrosion tests carried out in the present work ...... 38
Table 4-1. EDX measurements from Fig. 4-7 ...... 49
Table 4-2. EDX measurements from Fig.4-8 ...... 51
Table 4-3. EDX measurements from Fig. 4-9 ...... 52
Table 4-4. EDX measurements from Fig.4-10 ...... 53
Table 4-5. EDX measurements from Fig. 4-11 ...... 54
Table 4-6. EDX measurements from Fig. 4-15 ...... 57
Table 4-7. EDX measurements from Fig. 4-16 ...... 58
Table 5-1. XRD measurement parameters ...... 65
Table 5-2. Absorption coefficient of the main phases with Cu-Kα1 ...... 66
Table 5-3. Quantification of the XRD diffractograms from Fig. 5-5 ...... 70
Table 5-4. Phase composition at the thermodynamic equilibrium for 6 g matrix and 0.9 g CSK slag at 1400 °C 74
Table 5-5. Quantitative XRD of the matrix corroded with 0.9g CSK after 24h at 1400°C ...... 75
Table 5-6. EDX measurements from Fig. 5-10 ...... 76
Table 5-7. EDX measurements from Fig. 5-11 ...... 77
Table 6-1. Phase densities according to [BUC05]...... 82
Table 6-2. Density measurements of the matrix samples ...... 85
98
Table 6-3. Average pore diameter of matrix CA6 after sintering at 1700°C/6h measured with mercury porosimetry ...... 87
Table 6-4. EDX measurements from Fig. 6-8...... 90
99
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